Novel HfB2-SiC-MoSi2 composites by reactive spark plasma sintering

Novel HfB2-SiC-MoSi2 composites by reactive spark plasma sintering

Journal of Alloys and Compounds 809 (2019) 151705 Contents lists available at ScienceDirect Journal of Alloys and Compounds journal homepage: http:/...

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Journal of Alloys and Compounds 809 (2019) 151705

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Novel HfB2-SiC-MoSi2 composites by reactive spark plasma sintering S. Ghadami a, E. Taheri-Nassaj a, *, H.R. Baharvandi b a b

Department of Materials Science and Engineering Tarbiat Modares University, PO Box 14115-143, Tehran, Iran School of Metallurgy and Materials, College of Engineering, University of Tehran, Tehran, Iran

a r t i c l e i n f o

a b s t r a c t

Article history: Received 21 May 2019 Received in revised form 1 August 2019 Accepted 4 August 2019 Available online 9 August 2019

In situ HfB2-SiC-MoSi2 composites were fabricated using HfB2, Si, Mo, and C powders as starting materials. Fully-dense composites achieved by Reaction Spark Plasma Sintering (RSPS) method at 1850  C under 40 MPa for 5 min. Thermodynamic calculations and microstructural studies proved that SiC and MoSi2 second phases were formed and homogeneously distributed according to the following reactions: SiþC]SiC and Moþ2Si]MoSi2. In addition, the presence of in situ SiC and MoSi2 reinforcement phases have also confirmed by field emission scanning electron microscopy, X-ray diffraction pattern and considering thermodynamics calculations made by HSC software. Besides, the mechanical investigations revealed that in comparison with monolithic HfB2, the hardness and fracture toughness of these composites were increased dramatically, and reached up to 25.2 GPa and 5.1 MPa m1/2, respectively. © 2019 Elsevier B.V. All rights reserved.

Keywords: HfB2-SiC-MoSi2 composites Reactive spark plasma sintering Mechanical properties

1. Introduction Amidst recent decades, lots of research for finding engineering materials which are standing against high temperature and severe environment, have been continued [1e4]. Among them, carbide, nitride, and boride of transition metals (Hf, Zr, Ti, Ta) with a melting point over 3000  C are called Ultra-High Temperature Ceramics (UHTCS) [5]. Although carbide compounds have a higher melting point compared to boride compounds, borides have been more practical on account of high oxidation resistance. Moreover, numerous studies in the case of UHTCS have been focused on ZrB2 [6e10]. Compared to other borides, HfB2 has the highest melting point, hardness, electrical and thermal conductivities. Therefore, HfB2 is the best candidate for high temperature applications such as the leading edge of re-entry vehicles, scramjet components, and refractory parts in aerospace applications, where operating temperatures can be reached to 2000  C [11,12]. It is obvious that owing to its strong covalent bonding, low selfdiffusion coefficient, and surface oxides and impurities, it is very difficult to densify HfB2 without applying a pressure and adding some additives during sintering process. As it is well known, Reactive Spark Plasma Sintering (RSPS) is one of the unprecedented methods being utilized to reduce the time and the temperature of sintering process [13]. In this method, in situ phase or phases newly

* Corresponding author. E-mail address: [email protected] (E. Taheri-Nassaj). https://doi.org/10.1016/j.jallcom.2019.151705 0925-8388/© 2019 Elsevier B.V. All rights reserved.

formed during the sintering process which can contribute to full densification and enhance the mechanical properties of ceramics. Only limited investigations have been carried out on RSPS of HfB2based composites. Gurcan et al. [14] fabricated monolithic HfB2 using HfO2 and B as the starting powders via RSPS method at 2050  C for 30 min under 60 MPa according to the following reaction: 3HfO2 þ 10B ¼ 3HfB2 þ 2B2O3(g)

(1)

where reaction product (B2O3) could be eliminated by its vaporization at the range of the temperature above 1200  C. They reported that the maximum relative density of HfB2 was 90%. Wang et al. [15] achieved 98.7% of relative density for HfB2-SiC composites by utilizing RSPS technique at 1600  C under a pressure of 40 MPa for 10 min with powder mixture composed of HfSi2, B4C, and C. Where both HfB2 and SiC phases emerged from the exothermic reaction as following: HfSi2 þ B4C þ 3C ¼ 2HfB2 þ 4SiC

(2)

However, other HfB2-based composites studies were performed by adding some dopants such as WC [16e19], Si3N4 [20], B4C [21,22], HfC [23,24], TaSi2 [25], MoSi2 [26,27], and SiC [24,28,29] which were focused on the role of dopants on microstructure and evaluated in situ phase forming and interactions between the constituents. None of them has studied on the formations of in situ second phase (or phases) in HfB2-based composites separately.

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In the given work, fabrication and densification of HfB2-SiCMoSi2 composites from starting powders of HfB2, Si, Mo, and activated carbon via reactive spark plasma sintering method have been investigated. We demonstrate that how in situ second phases can be effective on density, microstructure, and mechanical properties. 2. Experimental methods Commercial HfB2, Mo, Si, and activated carbon as the starting powders were selected to synthesis HfB2-15 vol%SiC-15 vol%MoSi2 composites. To obtain a precise composition of final composites, volume fractions calculations were done to achieve final composition according to theoretical densities of 11.2 g/cm3 for HfB2, 3.2 g/ cm3 for SiC, and 6.26 g/cm3 for MoSi2. Tables 1 and 2 show the characterization of as-received powders and the composition of the monoclinic HfB2 (Hf0) and HfB2-15 vol%SiC-15 vol%MoSi2 (Hf1) composites, respectively. Mixed powders were milled by highenergy planetary mill using WC-Co cup and balls. The milling process was accomplished after 3 h at 300 rpm. In order to restrict preferential precipitation of constituents with different densities, the milling process was carried out under an ethanol atmosphere. The weight ratio of balls to powders was 3:1. The slurry was dried in air and then the powders sieved using mesh number of 100. For synthesis composite and formation in situ reinforcement phases, sintering process was done using a commercial spark plasma sintering apparatus, at 1850  C for 5 min in the vacuum of 0.05 mbar under the applied pressure of 40 MPa according to the following exothermic reactions: Si þ C ¼ SiC

(3)

Mo þ 2Si ¼ MoSi2

(4)

The pellet-like specimens (25  8 mm2) were ground using Cubic Boron Nitride (CBN) rotary disk. Before SEM characterization, the specimens' surfaces were ground and polished with SiC abrasive papers and a fine diamond paste to yield a mirror-like surface. Then, the microstructure of samples was directly examined without thermal or chemical etching step. The relative density was calculated for each specimen according to the ratio between the bulk density employing Archimedes' method and theoretical density using role of mixture which was based on the final compositions in the specimens. To detect and distinguish in situ formed phases, phase analysis was carried out using X-ray diffraction pattern (XRD, Philips, Model: X'Pert MPD, Tube: Co, and l: 1.78897 Å), and field emission scanning electron Microscope (FESEM, TESCAN, Model: MIRA) equipped with energydispersive spectroscopy (EDS). Because of the noticeable difference between the effective atomic number of HfB2, MoSi2, and SiC phases, the clearly different areas (light to dark areas) were observed and analyzed using back-scattered electron modes. The grain size of the specimens was estimated based on the line intercept method (ASTM E112-13) utilizing ImageJ software. Furthermore, thermodynamics calculations were done to realize phase formation during RSPS process and to support the extracted results from microstructural studies using HSC software. Micro-

hardness was measured by a Vickers indenter with 0.3 kg applied load for 10 s on polished sections, according to ISO6507 standard. To ensure results' reliability, for each specimen, 20 indentations were made and 40 diagonal lengths were measured. The fracture toughness (KIC) calculations were done using Evans and Charles equation based on the measurements of the radial crack length produced by Vickers [30]:

KIC ¼ 0:16 ðc=aÞ3=2 ðHVÞ1=2

(5)

where KIC is the fracture toughness (MPa m1/2), HV means Vickers hardness (GPa), c is the average half-length of the crack acquired in the tips of the Vickers marks (m), and a is the average half-length of indentation diagonal (m). 3. Result and discussion 3.1. Characterization and preparation conditions of materials Micrographs and XRD Patterns of as-received HfB2 powder and the combination of powders after the milling process are illustrated in Fig. 1. The trace of HfO2 impurity is detectable in both HfB2 starting powder and mixing powders. Based on calculations were performed by ImageJ software, the average particle size of asreceived HfB2 powders and mixing powders were estimated below 20 mm and 1 mm, respectively. After the milling process (Fig. 1c), no reactions were carried out between powders and no newly phases were detected. Even though, WC peak is visible in XRD pattern which came from WC-Co media during the milling process. Kalish and Clougherty [31] demonstrated that the impurity could be incorporated into powders mixture from milling media which is affected to the final density. It has been reported that the undersized particle and extended particle distribution could be a noticeable effect on the sintering behavior of ceramics which are specially fabricated using pressure assistant techniques [15]. The higher densities of specimens attributed to the reduced particle size of starting powders result from the high-energy milling process. 3.2. Synthesis of in situ phases Fig. 2 illustrates the diagram of reaction possibility between starting powders which is simulated according to RSPS condition extracted by HSC software. Si could react with C as well as Mo reacts with Si. Hence, SiC and MoSi2 could be formed during the temperature up to 1850  C. Pampuch et al. [32] studied the formation of SiC using Si and C powders by hot press method. They have shown that when the powders mixture of Si and C are heated above 1000  C in argon atmosphere, the reaction between Si and C is completed at 1300  C in solid state diffusion mode and lead to SiC powder formation. On the other hand, Deevi et al. [33] proved that MoSi2 tetragonal phase could be formed from the reaction between Mo and Si according to the reaction (4). From Fig. 2 there is a possibility of Mo2C phase formation based on the reaction (6) as following:

Table 1 Characterizations of starting powders. Row material

Purity (%)

Mean particle size (mm)

Manufacture (country)

HfB2 Mo Si C

95 99 99.5 99

20 10 10 10

Beijing Cerametek Materials (China) Merck (Germany) Merck (Germany) Shanghai activated carbon (China)

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Table 2 Specification and preparation condition of HfB2-based composites.
SiC (vol%)

MoSi2 (vol%)

Relative density (Closed porosity) (%)

Temperature/Pressure/Dwell time of RSPS Process ( C/MPa/min)

Vickers Hardness Matrix Grain Size Fracture Toughness (GPa) (mm) (MPa m1/2)

Hf0 Hf1

0 15

0 15

95 (5) 98.6 (1.4)

1850/40/5 1850/40/5

18±0.62 25:2±0.61

5.1 3.25

3±0.28 5.1±0.13

Fig. 1. SEM micrographs of (a) as-received HfB2 powder, (b) mixed powders after 3 h milling and (c) XRD patterns for (a) and (b).

Fig. 2. Calculated multiphase equilibrium using HSC using HfB2, Si, C, and Mo as starting materials according to RSPS processing condition (P~5 Pa) and (b) close up of thermodynamic equilibrium products enclose by the dashed line.

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2Mo þ C ¼ Mo2C

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(6)

In fact, if the reaction (6) proceeded, the reaction product could react with Si. Hence, MoSi2 and in situ SiC phases were formed according to the reaction (7) as following: Mo2C þ 5Si ¼ SiC þ 2MoSi2

(7)

Nonetheless, HfB2 does not react with reinforcements or other dopants when used as the based-ceramic matrix composites. However, decomposition reactions of SiC and MoSi2 are possible based on thermodynamics, it seems that the kinetic energy required for activation of these reactions is not adequate. Therefore, the decomposition of SiC and MoSi2 could be very slow up to 1850  C. Sciti et al. [34] demonstrated that MoSi2 could be remained in hot-pressed HfB2-MoSi2 composite. In other similar research, Wang et al. [15] reported that in situ SiC phase was not decomposed and homogeneously distributed through the microstructure after SPS process at 1600  C. When the powders mixture (HfB2, Mo, Si, and C) sintered up to 1850  C during RSPS process, three steps could happen: at the first step, HfB2 was not reacted with Si, Mo, and C. In the second step, reactions (3) and (4) could be progressed in parallel until whole Si content consumed led to formation of in situ SiC and MoSi2 reinforcements phases. And finally, if the reaction (6) was accomplished, in situ Mo2C would form then react with residual Si (reaction (7)), and consequently in situ SiC and MoSi2 phases are formed. Fig. 3 illustrates the diagram of Gibbs energy as a temperature function for reactions (3) to (7) via HSC software according to standard state. Based on thermodynamics, if there are several reactions, initially a reaction with large negative delta G value will proceed. Amongst from the reactions (3) to (6), the reaction (4) has the largest negative delta G value at the range of 0e1850  C and then can take place initially. From Fig. 3, the reaction (4) has a turning point at 1400  C and after this point, delta G value increases dramatically. It can be concluded that the reaction (4) could mostly be accomplished at 1400  C. The delta G value of this reaction was calculated 128.417 kJ which was completely fit with findings of Zakeri et al. [35]. By comparing reactions (3) and (6), the delta G value of reaction (3) is lower than reaction (6) in the temperature

range of 0e1200  C (the delta G value of reactions (3) and (6) were calculated 59.94172 kJ and 59.22711 kJ, respectively). When the temperature rises up more than 1200  C, the delta G value of reaction (6) is more negative. Furthermore, it could be deduced that the reaction (6) has more chance to take place in the temperature range of 0e1200  C. Based on this hypothesis, throughout sintering process up to 1850  C, the reaction (4) could be mostly accomplished and lead to MoSi2 formation. Among from the reactions (3) and (6), the reaction (3) could proceed up to 1200  C and SiC could be formed. Above this temperature the reaction (6) was dominant; hence, Mo2C could be formed. If the reaction (6) accomplished, Mo2C could react with Si according to the reaction (7) and lead to the formation of MoSi2 and SiC at above 1200  C. According to the reaction (7), formations of these phases were perfectly completed at 1400  C. 3.3. Density The relative densities of specimens are presented in Table 2. The calculated relative densities of Hf0 and Hf1 were 95% and 98.6%, respectively. These results showed that the density of Hf0 is 8.2% more than of the maximum density of SPSed monolithic HfB2 reported by Anselmi-Tamburini et al. [36]. On the other hand, the density of Hf1 is close to the final density. By and large, non-oxide ceramic powders have thin films of oxide layer that oxide content increased if the powder size close to nano-scale. In other words, by decreasing powder size of HfB2 particles, the specific surface areas of powders increases and leads to the increase of chemical activity of powders. Hence, there is a higher chance of the reaction between HfB2 and O2 as following reaction: HfB2 þ 2.5 O2(g) ¼ HfO2 þ B2O3

(8)

Besides, HfO2 would be formed during the milling process as previously described. The whole HfO2 content can be covered HfB2 particles as a film which obstacle mass transfer during RSPS process which negatively affected on densification. This scenario is favorably supported by the findings of other researcher [37]. Monteverde and Bellosi [38] reported that carbon could be penetrated to structure during the consolidation process. Penetrated carbon from graphite mold could be reacted with HfO2 impurities and caused to eliminate them which effectively enhanced densification mechanism according to the given reaction: HfO2 þ 3C ¼ HfC þ 2CO(g)

(9)

It has been reported that this reaction (9) could take place above 1700  C [39]. Based on calculation extracted by HSC software in the standard state, this reaction could completely happen at 1700  C and DG of 17.303 kJ (Fig. 4). Furthermore, HfO2 impurities could be also reacted with incorporated WC from the milling process according to the reaction proposed by Ni et al. [16]: HfO2 þ WC ¼ HfC þ 3W þ 2CO(g)

Fig. 3. Changing Gibbs free energy by temperature for reactions (3, 4, 6 and 7) at standard state (P~1 atm).

(10)

From Fig. 4, HfC phase could be formed from the reaction (10) at 2000  C and DG of 15.684 kJ. In competition between reactions (9) and (10), with larger negative delta G in the temperature range of RSPS process, the reaction (9) could be progressed predominantly. Thermodynamics prediction for the formation of HfC according

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Fig. 4. Gibbs free energy of reactions of HfO2/WC and HfO2/C as a function of temperature at standard state (P~1 atm).

to RSPS conditions (temperature up to 2000  C and mild vacuum about 0.00005 bar) by HSC software are illustrated in Fig. 5. It is clearly concluded that the reducing pressure from 1 bar (standard state) to 0.00005 bar (vacuum pressure in this work) caused to reduces the favorable temperature of reaction (9) and (10) to about 1100  C and 1200  C, respectively. It is explicitly worth noticing that with the ability to remove surface oxide, RSPS technique played a significant role in sintering mechanism and ameliorating densification as schematically shown in Fig. 6. Therefore, increasing densities of specimens are mainly attributed to remove HfO2 impurity. It has been reported that SiC and MoSi2 are advantageous roles on densification behavior of HfB2-based composites during sintering process [34,40]. In situ formed MoSi2 was very soft at the final temperature during RSPS process (about 50  C lower than its melting point) and could be located between the porosities in HfB2 skeleton. Interesting to realize that B2O3 from the reaction (8) has a melting point of 450  C which was liquid at the sintering temperature. Moreover, Si powder particles are coated by a silica layer (similar to described for HfB2), which can solve in B2O3 and cause the formation a liquid borosilicate phase according to equilibrium phase diagram [41]. This newly formed liquid phase could stream

Fig. 6. Schematic drawing of eliminating HfO2 impurities during RSPS process.

through the capillaries, filled remaining porosities, removed residual impurities, and enhanced relative densities of specimens. It is also worth to note that the evaporation-condensation mechanisms will vanquisher if the oxide-impurities are remains in HfB2 skeleton. These mechanisms cause grain growth instead of lattice diffusion. 3.4. Hardness The micro-hardness and the fracture toughness of the RSPSed composites calculated by the indentation method are given in Fig. 7. The enhanced micro-hardness of composite owning to four factors as below: (1) Density; Sciti et al. [42] reported that residual porosity has a negative effect on density. With concerning the fact that

Fig. 5. Thermodynamic equilibrium calculated at RSPS processing condition (P~5 Pa) for (a) HfO2/C reaction and (b) HfO2/WC reaction.

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3.5. Phase composition and microstructure

Fig. 7. The dependency of hardness and fracture toughness values of Hf0 and Hf1 RSPSed at 1850  C.

there is a direct link between porosity and density, the hardness value of as-sintered composites enhances by increasing density. It is notable that in this work, as one of the powerful methods for sintering, Spark Plasma Sintering (SPS) has greatly contributed to full densification of the specimens. (2) The intrinsic hardness of constitutive phases; It has been reported that the hardness values of HfB2, a-SiC, and MoSi2 are 21.2, 32, and 10e12 GPa, respectively [43,44]. Based on rule of mixture, by incorporating the second phase with a higher hardness value, the hardness of composite will increase. In this work, the mean value of the Vickers hardness of SPSed Hf1 was 25.2 GPa, more than Hf0 (18 GPa) because of the formation of SiC as a second phase. (3) Reduction in matrix grain size as reported by Zener-Smith equation [45]:



4r 3f

(11)

where Z is the matrix grain size (mm), r is the average size of the second phase particles (mm), and f is the volume fraction of the second phases (SiC and MoSi2). The formation volume fraction of SiC and MoSi2 particles led to a decrease in mobility of matrix grain boundaries and reduce HfB2 grain size, subsequently (see Table 2). But this point would also be mentioned that decreasing in particle size causes an increase in strength, according to Hall-Petch relation. (4) Formation in situ phases with excellent characterization; As discussed earlier, MoSi2 and SiC reinforcement phases were formed during RSPS process by exothermic reactions. Unlike common composites, in situ composites have strong bonding between reinforcement and matrix. Because of such a strong interface between the reinforcement phases (MoSi2, SiC) and matrix (HfB2), Hf1 composite became much stronger and its hardness increased to 25.2 GPa. This point should be mentioned that a weak interface is one of the reasons of failing, especially about composites with two or more phases if the thermal expansion coefficient of phases is significantly different. A higher hardness value of the Hf1 attributed to controlling and tailoring the microstructure by the formation of in situ MoSi2 and SiC particles which intimately contacted to HfB2.

XRD patterns of RSPSed specimens are given in Fig. 8 a. It is evidently clear that the only HfB2 and HfC crystalline phases for Hf0 as well as HfB2, HfC, MoSi2, and SiC crystalline phases for Hf1 are detected. Corresponding pick intensity of HfB2, the main crystalline phase, is sharp. In contrast, the intensities of predicted SiC and MoSi2 picks are less sharp, due to both lower densities and vol% of SiC and MoSi2, compared to HfB2. No corresponding peaks of starting materials (Si, C, and Mo) are detected, which proved that equations (3), (4) and (7) are completely proceeded. It has been reported that residual Si has a negative effect on mechanical properties of UHTCs ceramic [35]. As a result, the excellent mechanical properties of Hf1 attributed to fully completed reactions (3) and (4) and no residual Si phase. The probability of taking place of reaction (6) is increased by the trace of HfC peak as already discussed (see section 3.2), and no detectable oxide phase such as HfO2 is found. In order to detect any solid solution phase, X-ray diffraction pattern at high diffraction angles (100 <2q < 120 ) was carried out for Hf1 (Fig. 8 b). From this Fig, a shoulder distinguishes next to HfB2 peak on the right side at 2q ¼ 114.8 . It should be noted that based on ICDD: 00-017-0917, this pick does not belong to MoSi2 tetragonal phase. However, the K-Alpha radiation is stripped from XRD patterns using Xpert software; it is possible the shoulder belongs to Mo peak which is merged into HfB2 afterward. Therefore, it could be concluded that the shoulder corresponds to (Hf, Mo)eB solid solution which was formed during RSPS process. Since the covalent radiuses of Hf and Mo are 1.44 Å and 1.30 Å, respectively [46,47]. Hence, the formation of the solid solution of (Hf, Mo)eB could be possible. In a similar research on pressureless sintered ZrB2eSiC composites with Mo additive, Yan et al. [48] reported that the (Zr, Mo)eB solid solution was formed. The back-scattered SEM micrographs at the different magnification of RSPSed composites (Hf0, and Hf1) are shown in Fig. 9. The microstructure of monolithic HfB2 includes HfB2, HfC phases, and some porosity, according to EDS analysis. From Fig. 9 c, it is clearly seen that there are four regions: dark region, grey region, light region, and white region. EDS analysis (Fig. 9 d) revealed that SiC, MoSi2, HfC, and HfB2 refer to dark zone, grey zone, white zone, and light region, respectively. Consequently, SiC and MoSi2 phases in situ formed through reactions (3, 4, and 7), and homogeneously dispersed in HfB2 matrix. Similar to Hf0, the trace of HfC was detected and no distinguished HfO2 phase was discovered in microstructure of Hf1. The mean grain size of Hf0 and Hf1 composites were measured 5.1 mm and 3.25 mm using ImageJ software, respectively. Moreover, the mean grain size of the composites are reduced by incorporating the second phase, as described in the previous investigations [49,50]. In this study, in situ formation of SiC and MoSi2 caused decreasing in grain size (see section 3.4 (3)). By comparing Fig.9 a, and Fig.9 b, it is clearly seen that the mean grain size of Hf1 RSPSed composite is less than Hf0 which support this scenario. Monteverde [51] described that the reinforcement particles can hinder the HfB2 grain growth. Furthermore, the interaction between the grain boundaries and solid solution (Hf, Mo)eB may inhibit grain boundary migration and then caused to grain size is reduced. This point should be noted that the fast heating rate and short holding time during RSPS process could effectively reduce the mean grain size of Hf0 and Hf1. The back-scattered electron microscope image of the microstructure of Hf1 in high magnificent and the Line-Scan EDS analysis corresponding to the marked line are shown in Fig. 10 a, and b, respectively. Aside from the limitation of EDS analysis for

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Fig. 8. XRD patterns of HfB2-based composites RSPSed at 1850  C (a) Hf0 and Hf1 (b) close up high angle XRD pattern in the 2q range of 100e120 for Hf1.

Fig. 9. Back-scattered electron images of polished surface for (a) Hf0, (b) Hf1, (c) Closer view of (b) and (d) EDS pattern spot scan for spot 1, spot 2 and spot 3 in (c).

identification light element such as boron (B), other elements such as Hf, Mo, Si, C, and O are detected by EDS analysis. Evidently, SiC is strongly bonded to HfB2 grains and no trace of impurity and transition phase are found in the interface between reinforcement particle and HfB2 grains according to line-scan EDS analysis (Fig. 10 b). It could be concluded from this result that, no reactions took place between HfB2 and starting materials (Mo, Si, and C) before in situ formation of SiC and MoSi2. Moreover; no reactions could happen between HfB2 and reinforcements particles after in situ formation of SiC and MoSi2. Another result which could be deduced from this Fig is formation solid solution (ss) between SiC and HfB2 grain. Based on line-scan EDS analysis, at the starting point (SiC grain zone) only Si, C, and negligible of Hf and Mo peaks are detected. Right next to SiC grain, a small size zone (about 1 mm) is found. Getting distance from the starting point, the intensity of Hf

and Mo peaks are increased as well as the intensity of Si and C peaks are decreased which means that (Hf, Mo)eB solid solution could be formed. It could be hypothesized that HfB2 grain has a substructure, similar to a core-shell structure. A small amount of Mo could be replaced on HfB2 structure caused to (Hf, Mo)eB solid solution shell. The above-mentioned results are in agreement with the findings of other researchers [27,52] and phase composition is detected by X-ray diffraction in Fig. 8. Fractured surface of both Hf0 and Hf1 composites are presented in Fig. 11 a, and b, respectively. As can be seen from Fig. 11 b, after fracture, the HfB2 grains perfectly remained unchanged. in other words, the crack caused to rupture propagated through grain boundaries owing to grain boundaries were weaker than that inside the grains. Smooth fractured surface attributed to the inter-

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Fig. 10. (a) High magnification Back-scattered electron image of Hf1 (b) EDS line-scan corresponding to the marked line starting from SiC grain showing formation of ss ¼ solid solution (Hf, Mo)eB.

Fig. 11. SEM fractographs of (a) Hf1 showing trans-granular fracture mode, (b) Hf0 showing inter-granular fracture mode and (c) Closer view of (b) indicating pull-out of HfB2 grain.

granular fracture mode which played a dominant role of failure. After in situ formation of SiC and MoSi2 reinforcement particles in Hf1 composite, the strength of grain boundaries was enhanced, therefore, the crack caused to rupture preferring to propagate through the inside of grains. Rough fractured surface attributed to the trans-granular fracture mode which played a dominant role of failure.

Fig. 11 c presents the close-up fractured surface of Hf0. As presented by the dashed square; the vacancy of HfB2 grain is visibly exposed. Worth noticing that the grain pulled-out during rupture is consistent with inter-granular fracture mode theory. No trace of porosity is found in microstructure after RSPS process which is adjusted by the calculated results from density measurements.

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mechanical properties. Due to effective roles of WC and C by eliminating oxideimpurities as well as Mo by formation (Hf, Mo)eB solid solution phase, the relative density of HfB2-15 vol%SiC-15 vol%MoSi2 composite was improved remarkably and reached 98.6%. Microstructural studies and XRD analysis revealed that SiC and MoSi2 reinforcement phases are completely formed. The Vickers hardness and fracture toughness values of the composite were calculated about 25.2 GPa and 5.1 MPa m1/2, respectively. Remarkable fracture toughness is mainly attributed to crack pining, crack branching, crack bridging and the needle-like shape of SiC particles. The fracture surface of the composite was finally proved that the fracture mode was trans-granular. References

Fig. 12. SEM image of indentation crack propagation of Hf1 showing toughening mechanisms including crack bridging, crack branching and crack deflection.

3.6. Fracture toughness The fracture toughness values of Hf0 and Hf1 composites are also shown in Table 2 and Fig. 7. As can be seen, by in situ formation of SiC and MoSi2 reinforcements particles, the calculated fracture toughness had a dramatic increase and reached 5.1 MPa m1/2. Generally speaking, the amount of fracture toughness value of monolithic HfB2 ceramic is low (for this study about 3 MPa m1/2) due to the high strength covalent bonding nature of HfB2 ceramic. When some reinforcement like SiC and MoSi2 are added to monolithic HfB2 ceramic, the fracture toughness of composite was increased by reduction energy of the tip of the crack. Through crackpropagation, the crack was deflected as shown in Fig. 12. This mechanism could be affected more impressively when reinforcements particles are formed during RSPS process. When the crack strikes to SiC or MoSi2 reinforcements particles, more energy is wasted by crack deflection, crack pinning, crack bridging, and crack branching due to such strong intimate bonding between reinforcements and matrix. It means the fracture toughness is improved. It has been explicitly noted that SiC has two polycrystalline structures. The high temperature phase, alpha-SiC, has a needlelike morphology, whereas the low-temperature phase, beta-SiC, has a spherical morphology. In this case, SiC particles in situ formed at high temperature during RSPS process. Therefore, it can be anticipated that the morphology of SiC would be needle-like. In previous similar study [50], the effect of reinforcement particle morphology on the fracture toughness was investigated. As a result, with generating longer obstacles, needle-like SiC particles could effectively improve fracture toughness. 4. Conclusions HfB2-15 vol%SiC-15 vol%MoSi2 composite has been densified by Reactive Spark Plasma Sintering (RSPS) using HfB2, Si, Mo and C as starting materials at 1850  C under 40 MPa for 5 min. Reinforcement particles (SiC and MoSi2) were formed by in situ reactions during RSPS process which played a significant role to promote

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