Int. Journal of Refractory Metals and Hard Materials 41 (2013) 507–516

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Int. Journal of Refractory Metals and Hard Materials

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Thermal shock behavior of ZrB2–SiC composite ceramics withadded TaSi2Shubin Wang a,⁎, Cheng Xu a, Yongbin Ding a, Xinghong Zhang b

a School of Materials Science and Engineering, Beihang University, Beijing 100191, PR Chinab Centre for Composite Materials, Harbin Institute of Technology, Harbin 150001, PR China

⁎ Corresponding author. Tel./fax: +86 10 8231 6500.E-mail address: shubinwang@buaa.edu.cn (S. Wang

0263-4368/$ – see front matter © 2013 Elsevier Ltd. Allhttp://dx.doi.org/10.1016/j.ijrmhm.2013.06.010

a b s t r a c t

a r t i c l e i n f o

Article history:Received 6 March 2013Accepted 30 June 2013

Keywords:ZrB2-based ultrahigh-temperature ceramicsThermal shock behaviorMechanical propertiesOxidation

The thermal shock behavior of ZrB2–SiC composite ceramics was remarkably enhanced by the addition ofTaSi2 particles. TaSi2 enhanced the residual strength within the temperature range resulting from thecrack-healing effect, especially at ~1400 °C. Residual fracture toughness increased with increasing TaSi2 con-tent from 1000 °C to 1600 °C. The addition of soft phase TaSi2 positively affected the toughening mechanism.Considering the consumption of TaSi2 at ~1400 °C, fracture toughness of the TaSi2-containing ZrB2–SiC basedceramic decreased with increasing temperature. The formation of a compressive stress zone beneath the sur-face oxide layer increased fracture toughness at 1600 °C. The cohesive structure of the TaSi2-containingZrB2-based material inhibited crack formation and retained higher mechanical properties. Apparent surfacechanges were observed as a result of oxidation after severe thermal shock test. From the viewpoint of ther-modynamics, thermal oxidation resistance improved with the addition of TaSi2.

© 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Ultrahigh-temperature ceramics (UHTCs) such as borides and car-bides were developed in the 1960s [1]. ZrB2–SiC ceramics are UHTCsthat have an ultrahigh melting point, low electrical resistivity, highthermal conductivity, high Young's modulus, and good resistance tochemical erosion or corrosion attack. These unique characteristicsmake ZrB2–SiC ceramics useful in the thermal protection and propul-sion systems of hypersonic aerospace vehicles [2–5]. However, intrinsiccharacteristics of ZrB2–SiC ceramics, such as low fracture toughness andthe related poor toughness-induced thermal shock resistance, continueto hinder their wide application, particularly in harsh environments [6].

Thermal shock resistance depends on several mechanical andthermo-physical properties, such as fracture toughness, flexural strength,elastic modulus, coefficient of thermal expansion of materials, and oxi-dation resistance at high temperatures [7]. Numerous studies have beenconducted on improving the oxidation resistance of ZrB2–SiC-basedceramics, but reports on the fracture toughness and thermal shock re-sistance of this material system are limited. ZrO2 particles, ZrO2 fiber,AlN particles, graphite flake, carbon fiber, and SiC whiskers are com-monly used as toughening materials [8–14].

Recently, Opila [15] reported that the addition of TaSi2 significant-ly improves the oxidation resistance of the ZrB2–SiC matrix. Adding20 vol.% TaSi2 to ZrB2–SiC clearly improves the oxidation resistanceat 1627 °C in air, which is attributed to the effects of Ta but not ofSi. The TaSi2 composition is consumed faster than the ZrB2–20 vol.%

).

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SiC compositions, which were exposed to similar oxidation heat treat-ments at 1627 °C. This phenomenon is attributed to the melting ofTa2O5 (1785 °C) and/or compounds of Ta2O5 and ZrO2. The improvedoxidation resistance is evidence of phase separation on the amorphoussurface layer. Opila also found that adding TaSi2 to ZrB2-based UHTCscan reduce the densification temperature in the hot-pressing processand improve sintering performance. Speyer [16] found that Ta additionsto ZrB2–B4C–SiC in the form of TaB2 and/or TaSi2 increase the oxidationresistance throughout the entire evaluated spectrum of temperatures,confirming that TaSi2 is a more effective additive than TaB2.

A few studies on the effect of adding TaSi2 on the thermal shockbehavior of ZrB2–SiC nanocomposite ceramics have been reported.Our previous research demonstrated that ZrB2 with 5 wt.% SiC(100 nm particle size) exhibits good sintering performance, strength,and fracture toughness at room temperature [17]. Thus, this study in-vestigated the effect of TaSi2 content on the microstructure, mechan-ical properties, and thermal shock behavior of the composite. ZrB2–SiC–TaSi2 composite ceramics (ZSTx) were prepared by spark plasmasintering (SPS).

2. Experimental procedure

2.1. Material preparation

Commercially available ZrB2 powder (~10 μm, N99%; Harbin Insti-tute of Technology, China), TaSi2 powder (N99.9%; Alfar Aesar, USA),and β-SiC powder (~100 nm, N99.9% pure; Kaier NanotechnologyDevelopment Co., Ltd., Hefei, China) were used as raw powders. Pow-der mixtures with varying TaSi2 content (0 wt.%, 5 wt.%, 10 wt.%, and

Table 2Physical properties of ZrB2–SiC–TaSi2 ceramics.

Designation Measured density(g/cm3)

Relativedensity (%)

σf (MPa) KIC (MPa·m1/2)

ZS 5.52 94.81 595.21 ± 13.21 5.44 ± 0.12ZST5 5.73 96.09 511.43 ± 19.65 6.06 ± 0.10ZST10 5.74 97.47 477.36 ± 12.24 6.74 ± 0.17ZST15 5.80 98.01 443.67 ± 21.36 7.91 ± 0.10

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15 wt.%) were ball milled for 24 h. After mixing, the slurry was driedin an oven and milled in an agate mortar and pestle. The powder mix-ture was sintered using an SPS equipment (SPS-T1050, Japan). Thepowder mixture was loaded into a graphite die (40 mm inner diam-eter) lined with graphitized paper. The chamber for sintering was ina vacuum (b10 Pa). During the sintering process, the specimen washeated from room temperature to 600 °C for 3 min and from 600 °Cto 1700 °C at 100 °C/min for a holding time of 10 min. A pressure of35 MPa was applied during the heating process, and 50 MPa wasused during the holding time. The composition of the prepared sam-ples is shown in Table 1.

2.2. Thermal shock test

The samples were heated from 1000 °C to 1600 °C in air and heldfor 10 min to eliminate any temperature gradient. The samples werethen quenched in a water bath. A thermal shock test was carried outby measuring the residual flexural strength and fracture toughness ofthe specimens. For the fracture toughness test, notches were made oneach sample after the thermal shock by wire-cut electrical dischargemachining.

2.3. Characterization

The density (ρ) of each sample was measured by the Archimedes'method, whereas the relative density was estimated by the rule ofmixture, assuming that the true densities were 6.09 g cm−3 forZrB2, 3.21 g cm−3 for SiC, and 8.83 g cm−3 for TaSi2. The crystallinephases were identified by X-ray diffraction (XRD; D/max 2000, Rigaku,Japan). Themicrostructure featureswere analyzed by scanning electronmicroscopy (SEM;mod. CS3400, Cambridge, UK) and energy dispersivespectroscopy (EDS; mod. INKA Energy 300, Oxford Instrument, UK). Allspecimens were ground and polished with diamond slurries down to1 μm. The edges of all specimenswere chamfered tominimize the effectof stress concentration due to machining flaws. The flexural strength(σf) was tested by three-point bending at room temperature with3 mm × 4 mm × 36 mm chamfered bars, using a span of 30 mm andcrosshead speed of 0.2 mm/min (Instron 5565). The fracture toughness(KIC) was evaluated by the single-edge notched beam test (crossheadspeed of 0.02 mm/min) using 2 mm × 4 mm × 22 mm bars on thesame jig that was used for strength.

3. Results and discussion

3.1. Properties and microstructures of the ZrB2–SiC–TaSi2composite ceramics

The effect of TaSi2 on the relative density of ZrB2–SiC–TaSi2 isshown in Table 2. The addition of TaSi2 particles clearly improvesthe relative density of ZrB2–SiC ceramics. The surface microstructuresof ZS and ZST15 are shown in Fig. 1(a) and (b), respectively. Theboundaries of the particles of ZST15 are inconspicuous comparedwith those of ZS.

As shown in Table 2, adding TaSi2 significantly improves fracturetoughness, but simultaneously reduces flexural strength. The mea-sured fracture toughness values increase from 6.06 MPa·m1/2 to7.91 MPa·m1/2, with the content of TaSi2 increasing from 5 wt.% to

Table 1Nomenclature for the ZrB2–SiC based ultra-high temperature ceramics.

Composition Designation

ZrB2–5 wt.% SiC ZSZrB2–5 wt.% SiC–5 wt.% TaSi2 ZST5ZrB2–5 wt.% SiC–10 wt.% TaSi2 ZST10ZrB2–5 wt.% SiC–15 wt.% TaSi2 ZST15

15 wt.%, which is higher than the values of the ZrB2–SiC compositematerial. Although the introduction of TaSi2 reinforcement phasecan increase fracture toughness, it can also reduce flexural strengthcompared with ZrB2–SiC. The measured flexural strength values de-crease from 595.21 ± 13.21 MPa to 443.67 ± 21.36 MPa, with thecontent increasing from 0 wt.% to 15 wt.%. The increase in toughnessmay be due to the weaker interface bonding between the ZrB2 particlesand the TaSi2 reinforcement phase. Considering the different coeffi-cients of thermal expansion between the ZrB2 (CTE = 6.7 × 10−6/K)and TaSi2 particles (CTE = 14 × 10−6/K) [18,19], the addition of TaSi2particles to ZrB2–SiC results in the development of thermal stress uponcooling from the processing temperature. Thermal stress may weakenthe interphase boundaries, leading to a fracture along such boundaries,which can enhance crack deflection and branching aswell as stress relax-ation near the crack tip [20]. Given the addition of a second soft phase, theflexural strengths of the ZrB2–SiC–TaSi2 ceramics decrease accordingly,even though the relative density of the TaSi2-reinforced material ishigher. The mechanical properties of ZrB2–SiC–TaSi2 at ambient temper-ature are similar to our previous research [17].

3.2. Thermal shock behavior

Fig. 2(a) shows the residual strengths of the ceramics after thewater-quenching tests. All the specimens showed a drastic reductionin flexure strength compared with their initial strengths, even withthe reinforcement of TaSi2. The residual strength of ZST5 is lowerthan that of ZS through the entire content range, whereas the residualstrengths of ZST10 and ZST15 are higherwhen the temperature changesto below1600 °C. The residual strengths of ZS are 45.03 MPa at 1000 °C,43.17 MPa at 1200 °C, and 53.11 MPa at 1400 °C, whereas that of ZST15reaches 74.28, 90.07, and 83.11 MPa, respectively. At 1600 °C, the resid-ual strength of ZS is 96.69 MPa, whereas those of ZST10 and ZST15 are91.48 and 73.52 MPa, respectively. Fig. 2(b) shows the residual fracturetoughness of the ceramics after water quenching. The TaSi2 reinforce-ment does have an apparent improvement on the residual fracturetoughness, especially for ZST10 and ZST15. ZST10 and ZST15 exhibithigher fracture toughness values than ZS and ZST5 ceramics at all the in-vestigated temperatures. For ZST10 and ZST15, fracture toughness alsodegrades slightly, except for a sharp degradation at 1400 °C. In addition,the ZS and ZST5 results indicate that the influence of quenching temper-ature difference on fracture toughness is unclear. The fracture tough-ness of ZS is 2.72 MPa·m1/2 at 1000 °C and 2.01 MPa·m1/2 at 1600 °C,whereas that of ZST15 reaches 8.91 MPa·m1/2 and 7.89 MPa·m1/2,respectively.

The microstructure of the samples after the quenching test pro-vides an insight on the important details of thermal shock behavior.The literature confirms that thermal shock resistance, especially re-sidual strength, is susceptible to flaws that emerge during machiningor thermal shock test [21]. During a thermal shock test, a huge tem-perature difference occurs between the surface and the interior.Given the inconsistency of thermal expansion in different grains ordefects, such as pores or grain boundary, micro cracks occur easily.Stress concentration forms small crack tips, which in turn drive thecrack growth. These micro cracks have a negative effect on residualstrength. Fig. 3 shows the SEM micrograph cross-section of the fourdifferent amounts of TaSi2-containing ZrB2–SiC-based composite

Fig. 1. SEM images of the polished surfaces: a, ZS; b, ZST15.

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ceramics after thermal shock at 1600 °C. The cross-section micro-structure of ZS displays an obvious change compared with that ofthe ZSTx composite ceramics, which may be attributed to the differentoxidation resistances to ZSTx ceramics (this part will be discussed in thefollowing section). Fig. 3(a) shows that ZS has a borosilicate layer with~5 μm thickness, a ~17 μm crystalline zirconia layer, and a ~20 μmSiC-depleted region. The obvious micro-cracks that appear in the ZrO2

layer are caused by the volume change and the thermal stress duringwater quenching. Some micro-cracks are also formed in the unoxidizedlayer, indicated with a dotted line in the inset of Fig. 3(a). The formationof cracks implies that thermal stress concentration occurred at the flawtip and that the volume expansion in the substrate, produced by the tem-perature difference during water quenching, is severe enough to causecrack formation.

In contrast to the cross section of ZS, according to Fig. 3(b) to (d),the ZSTxmaterials have a compact substrate, and the typical layers dam-aged by oxidation duringwater quenching are unclear. In Fig. 3(b), somemicro-cracks (denoted by the ellipse) can be found between the grainboundaries, which are thought to cause the severity of ZST5 degradationafter water quenching. According to Griffith's brittle material fracturetheory, concentrated stress (σm) can be expressed in terms of the aver-age strength of the material (σ), flaw size (C), and curvature radius (ζ)of the flaw tip, as shown in Eq. (1).

σm ¼ 1þ 2

ffiffiffiCζ

s !σ ð1Þ

The curvature radius (ζ) of the flaw tip is much smaller than theflaw size (C). Thus, fractures form fast in the ceramic component ata lower average stress that is significantly lower than the strengthof the ceramics [21]. The grain boundaries are believed to be criticalin resisting crack propagation. Considering the stress concentration,

Fig. 2. Residual strength and fracture toughness vs. thermal shock temperature difference fortoughness.

defects such as pores or micro-cracks are easily formed at the grainboundaries. Therefore, cracks or pores, as a crack source, can propa-gate easily to reduce the strength as the ceramics withstand a violenttemperature change in ZST5. For the ZSTx materials, the relative den-sity increases with the addition of TaSi2. When a combination of rela-tive density (Table 2) and SEM analysis is employed, no indication ofporosity is found in ZST10 or ZST15, which means that a large tem-perature differences cause less damage to ZS10 and ZS15. Previousstudies have shown similar results [22,23].

Figs. 4 to 6 show the surfacemicrographs of ZS and ZSTx (5 wt.% and15 wt.% TaSi2) specimens thermally shocked at 1000 °C to 1600 °C. Sur-face changes, as a result of oxidation, are observed at 1000–1600 °C. Be-cause ZST10 and ZST15 show analogous surface changes from 1000 °Cto 1600 °C, this study only reports the micrographs of ZST15. Fig. 4shows the micrographs of a ZrB2–SiC sample shocked at 1000 °C to1600 °C. As Fig. 4(a) shows, the specimen is covered by a noncompactzirconia layer, which was confirmed by EDS to be due to the reactionof zirconium diboride to zirconia that is generated by quenching at1000 °C. The obvious micro-cracks found in the ZrO2 particles are at-tributed to both the volume change and the thermal stress duringwater quenching, which can penetrate into the substrate during crackpropagation. In addition, the borosilicate phase is seldom observed onthe surface, which may be due to the time limitation for the oxidationof specimens at 1000 °C. With the increase in thermal shock tempera-ture, the microstructure changes, as shown in Fig. 4(b) to (d). The mag-nified surface micrograph in the inset of Fig. 4(b) shows that zirconiagrains do not closely adhere to one another and do not completelyseal the surface of the sample, thus forming holes and porosity. Afterthermal shocking at 1400 and 1600 °C, the surface of the specimen ismainly covered by a continuous borosilicate layer, on which dendriticzirconia grains are embedded (Fig. 4c and d). This layer completelyseals the sample surface and impedes crack formation through itscrack-healing effect. The formation of borosilicate is consistent with

ZrB2–SiC-based composites with different TaSi2 contents: a, residual strength; b, fracture

Fig. 3. SEM images of the cross section for ZS, ZST5, ZST10, and ZST15 after thermal shock at 1600 °C.

Fig. 4. SEM images showing the surface of ZS after thermal shock at 1000, 1200, 1400, and 1600 °C. a, 1000 °C; b, 1200 °C; c, 1400 °C; d, 1600 °C.

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Fig. 5. SEM images showing surface of the ZST5 after thermal shock at 1000 °C, 1200 °C, 1400 °C and 1600 °C. a, 1000 °C; b, 1200 °C; c, 1400 °C; d, 1600 °C.

Fig. 6. SEM images showing the surface of ZST15 after thermal shock at 1000 °C, 1200 °C, 1400 °C and 1600 °C. a, 1000 °C; b, 1200 °C; c, 1400 °C; d, 1600 °C.

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the reported oxidation studies of ZrB2 system composites. Previousstudies have shown that ZrB2 starts to oxidize to boria and zirconia at~800 °C and that SiC oxidizes to silica at ~1100 °C [24]. Significantamounts of silica combinewith the available boria and yield borosilicateglass. Upon cooling to room temperature, borosilicate forms an amor-phous layer, which was reported to have a sealing effect on the cracks,decreasing its sensitivity. Considering the crack-healing property ofthe borosilicate phase, the strength of ZS increases accordingly at 1400and 1600 °C. The crack formation and porosity observed during thetest at 1000 and 1200 °C correspond to flexural strength degradation.

Fig. 5 shows the surface micrographs of ZS5 thermally shocked at1000 °C to 1600 °C. The surface microstructures of the oxidized spec-imen quenched at the temperature difference of 1000 and 1200 °Care similar, as shown in Fig. 4(a) and (b). The zirconia grains closelyadhere to one another, and the micro-crack size and density in thesurface decrease compared with the ZS specimen thermally shockedat 1000 and 1200 °C. The elemental analysis confirms the average dis-tribution of Ta in the oxide layer and the presence of a small quantityof Si. This finding is consistent with Diletta Sciti's research. Thesurface can form a borosilicate liquid layer that undergoes Ta enrich-ment due to the oxidation of TaSi2 at intermediate temperatures(~1400 °C), which is protective against oxidative damage to materials[25]. The formation of a liquid layer prevents crack formation andseals the flaws on the surface during thermal shock. Given the smallamount of TaSi2 in ZST5 and the evaporation of borosilicate duringthermal shock, the liquid phase on the surface after thermal shockis invisible. The surface, nonetheless, contains a compact oxidizedlayer, and only little cracks are formed compared with ZS. As the tem-perature difference increases to 1400 and 1600 °C, a borosilicate layeris formed. However, numerous micro-cracks easily develop on thesurface of the specimen, as shown in Fig. 5(c). The crack density andsize increase because of severe thermal stress during quenching,which is believed to weaken flexural strength. When the temperatureincreases to 1600 °C, bubbles form in the borosilicate, resulting in theoxidation of SiC to SiO(g) and the evaporation of the liquid phase.Bubbles tend to burst because of the vapor pressure exceeding the

Fig. 7. X-ray diffraction patterns of ZST15 at different thermal shock

ambient pressure, causing a large hole in the amorphous phase. Nota-bly, bubble formation can damage the continuous liquid layer, andlarge numbers of pores are formed in the borosilicate layer.

As shown in Fig. 6(a), the surface of the ZST15 becomes muchdenser, and little cracks develop on the surface compared with ZS andZST5. In contrast to ZST5 at 1200 °C, the surface of the ZST15 specimenis covered with a continuous borosilicate phase. Through EDS analysis,the white particles are confirmed to be mainly composed of oxygenand zirconium elements, and that Ta exists in the borosilicate layer.The discrepancy between ZST5 and ZST15 in the range of 1000 °C to1200 °C results from the much greater TaSi2 addition in ZST15 and thesignificant oxidation of TaSi2 occurring at a much lower temperaturethan SiC. Some cracks can also be observed on the surface, the size anddensity of which are much smaller than that of ZST5. Given the compactsurface and crack-healing effect of the borosilicate phase, the strength ofZST15 shows higher values than the others at 1000 °C to 1200 °C. As thethermal shock tests are performed at 1400 °C, bubbles formdue to vaporevaporation. As temperature increases to 1600 °C, the evaporation of theliquid phase intensifies, and SiC is oxidized to SiO [15]. As the tempera-ture increases to the range of 1400 °C to 1600 °C, the surface of the spec-imen is covered by an incontinuous amorphous phase layer as a result ofbubble formation. Thus, the retained strength of the specimen aftershocking at 1400 and 1600 °C decreases correspondingly.

As pointed out earlier in this paper, the fracture toughness valuesobtained for ZS remain relatively constant throughout the range of1000 °C to 1600 °C. However, the values for ZST10 and ZST15 sharplydecrease at ~1400 °C and relatively increase thereafter. During contrac-tion, the fracture toughness of ZST5 decreases slightly at ~1400 °C andthen increases at 1600 °C. Furthermore, much higher KIC values weremeasured for the ZrB2-based ceramics with the addition of TaSi2. Theenhancement of fracture toughness of the ZSTx materials may be asso-ciatedwith twomajor effects in the structure ofmaterials. (1) The cohe-sive structure of ZSTx materials, as shown in Fig. 3. Hu [26] confirmedthat the solid solution process probably decreases the boride grainboundary activation energy, thereby contributing to the formationof coherent structures of grain boundaries. A denser and cohesive

temperatures: a, untreated; b, 1000 °C; c, 1400 °C; d, 1600 °C.

Fig. 8. Weight gain vs. thermal shock temperature difference for the ZrB2–SiC-basedcomposites with different TaSi2 contents.

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structure can prevent crack formation and propagation. As Eq. (1)implies, high stress at the crack tip will decrease the fracture resistanceof materials. No indication of crack propagation is found in the interiorof ZST10 and ZST15. Therefore, the ZrB2-based ceramics added withTaSi2 retained higher fracture toughness than ZS throughout the tem-perature range. (2) Y. W. Mai revealed that crack-tip blunting resultingfrom plastic deformation in the materials increased the fracture tough-ness after thermal shock [27]. The addition of TaSi2 can increase fracturetoughness of ZrB2-based ceramics at room temperature because ofthe soft fracture at the crack tip. However, the consumption of TaSi2 at~1400 °C due to oxidation indicates that the toughening of TaSi2will weaken with an increase in temperature, especially at 1400 °C. Astemperature increases at 1600 °C, much more ZrO2 particles (Fig. 7)are oxidized from ZrB2 that are embedded in the borosilicate phase,which is attributed to the formation of an oxide layer. The compressivestress zone beneath the surface oxide layer (illustrated in Fig. 10) favor-ably inhibits crack initiation and propagation during quenching prior tofinal failure. This result is beneficial for increasing the residual fracturetoughness because the thermal stress is counteracted partially by theappearance of a compressive stress layer [21].

The XRD spectra in the 2θ range from 20° to 90° of the polishedsurface of the SPS ZrB2–SiC–TaSi2 ceramics are shown in Fig. 7(a).Cubic SiC, hexagonal ZrB2, and hexagonal TaSi2 are detected, alongwith a series of reflections, which are attributed to the solid solution

Fig. 9. Schematic diagram of the oxygen tr

that Ta dissolves into the ZrB2 lattice. Compared with pure ZrB2, thesepeaks are shifted toward higher angles, which indicate a contractionof the diboride unit cell. According to Vegard's rule, only Ta can enterthe diboride lattice in the ZrB2–SiC–TaB2 system because Ta5+ andZr4+ have similar sizes (0.073 nm vs. 0.08 nm), which is consistentwith the study of Sciti [25]. According to the EDS analysis in Fig. 1(a),a Zr–Ta–B solid solution phase forms. Fig. 7(b) and (c) shows the XRDpattern of ZST15 after 1000 and 1400 °C thermal shock. ZrB2 is partlyoxidized to ZrO2, and the Zr–Ta–B and Zr–Ta–O solid solution form.This phenomenon is due to the substitution of Ta at the Zr sites inZrB2 and ZrO2. As shown in Fig. 7(d), most of the ZrB2molecules are ox-idized to ZrO2 at 1600 °C to form a compressive stress zone beneath thesurface oxide layer, which promotes toughening at 1600 °C.

Fig. 8 shows the change in sample weight in the thermal shocktemperature range. At each temperature, the weight change of thesampleswith TaSi2 addition is less than thatwith ZS. Theweight changebecomes smaller with increasing TaSi2 content. At 1600 °C, ZST15 indi-cated a weight change of 0.68%, whereas ZS changed by 1.6%. Hence,TaSi2 addition can greatly improve the oxidation resistance of ZrB2–SiC-based ceramics.

The cross-section of the samples after thermal shock distinguishesthe microstructure and oxidation resistance discrepancy betweenZrB2–SiC ceramics and TaSi2-containing ZrB2 composite ceramics.According to the interpretation described above, Fig. 3(a) revealsfour distinct layers: (1) the silica-rich outer layer, which is thoughtto protect the ceramic from being oxidized; (2) the thin layer ofcrystalline zirconia containing little SiO2 glass wherein ZrO2 exhibitsvisible oriented growth, which suggests the large volume expansionupon conversion from ZrB2 to ZrO2; (3) the SiC-depleted region locat-ed underneath the ZrO2–SiO2 layer, which has a porous structure,with SiC partially or entirely removed; and (4) the unoxidized layer.This layer-like structure is similar to that reported in other studies[15,28,29]. Notably, small cracks were found in the sample after ther-mal shock.

As shown in Fig. 3(d), the micro-cracks and the layer-like structurecannot be observed. Furthermore, the EDS map results indicate thatthe element content of the oxide surface clearly changes. Hence, thelayer-like arrangement of the elements is retained.

The oxygen transfer mechanism determines the thickness of theoxide layer. For the dense ZrB2–SiC–TaSi2 composites, the ZrB2, SiC,and TaSi2 particles as well as the grain boundaries are the possibleroutes for oxygen transport. The presence of impurities at the grainboundaries is known to enhance oxygen transport. Thus, oxygentransport through the grain boundaries is faster than that throughthe particles. Oxygen reaches ZrB2, SiC, and TaSi2 particles through

ansport in ZrB2–SiC–TaSi2 composites.

Fig. 11. Gibbs-free energy change for reactions (2) to (4), (6), and (8).

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the grain boundaries and spreads to the surroundings, as shown sche-matically in Fig. 9.

As analyzed above, oxygen content decreases from the surface inward.The sample surface has three oxide layers, including the silica-rich outerlayer, the completely oxidized layer, and the semi-oxidized layer, asshown in Fig. 10.

Silica-rich outer layer: The SiO2 glass phase covers the entire sur-face. Different-sized bubbles are found in the glass phase. Large bub-bles are formed because of the coalescence of gaseous products (SiO,CO, and B2O3). The passive oxidation of SiC occurs in the oxygen-richlayer, as shown in Eq. (2). Active oxidation occurs in the oxygen-poorlayer, as shown in Eqs. (3) and (4). SiO2 glass vaporizes at tempera-tures higher than 1400 °C, as shown in Eq. (5).

SiC sð Þ þ 32O2 gð Þ→ CO gð Þ þ SiO2 lð Þ ð2Þ

SiCðsÞ þ O2ðgÞ→ COðgÞ þ SiOðgÞ ð3Þ

SiCðsÞ þ 2SiO2ðlÞ→ COðgÞ þ 3SiOðgÞ ð4Þ

SiO2ðlÞ→SiO2ðgÞ ð5Þ

ZrB2 is oxidized according to Eq. (6). B2O3 has a high vapor pres-sure and an unusually lowmelting point (450 °C). Therefore, B2O3 va-porizes quickly at high temperature, according to Eq. (7).

ZrB2 sð Þ þ 52O2 gð Þ ¼ ZrO2 sð Þ þ B2O3 lð Þ ð6Þ

B2O3ðlÞ→B2O3ðgÞ ð7Þ

TaSi2 is oxidized as follows:

TaSi2 sð Þ þ 134

O2 gð Þ→12Ta2O5 sð Þ þ 2SiO2 lð Þ: ð8Þ

As shown in Fig. 11, the Gibbs-free energy change, ΔG, for reac-tions (2) to (4), (6), and (8) were calculated according to the datacompiled by Yang Zhong, Fahrenholtz, and Opila [30–32]. Analysisshows that reactions (2), (3), (6), and (8) have a negative Gibbs-freeenergy change and that of reaction (4) is positive. Thus, all the reactionswith negative Gibbs-free energy change can occur under thermal shockcondition in the range of 1000 °C to 1600 °C. Furthermore, reactions (8)and (6) have higher reaction-driving force than reactions (2) and (3),suggesting that TaSi2 and ZrB2 may oxidize at a lower temperature,from the thermodynamic viewpoint. This result is consistent with themicrograph results that ZrB2 oxidizes to ZrO2 and TaSi2 oxidizes to

Fig. 10. Schematic diagram of the cross-section of the

form an amorphous liquid layer at a lower temperature. As with the ex-patiation above, the addition of TaSi2 improves the oxidation resistanceof ZrB2–SiC-based ceramics.

Oxidation resistance is improved by two possibilities: Ta2O5–SiO2

liquid phase with a higher viscosity and phase separation. The bub-bles tend to burst when the vapor pressure exceeds the ambient pres-sure. After the bubble bursts, the ceramic is possibly exposed in theair at high temperature. This phenomenon clearly reduces oxidationresistance properties [33]. The EDS result indicates that Ta is presentin this layer, as shown in Fig. 6(d). This result may be due to thehigher viscosity of the Ta2O5–SiO2 liquid phase, which can inhibitoxygen transport [15,16,32]. Phase separation is also indicated inFig. 6(d). A previous study [16] shows that phase separation playsan important role against oxidation. Therefore, the Ta2O5–SiO2 liquidphase is more protective than the SiO2 layer in ZS.

Adding Ta also results in the substitution of Ta at the Zr site inZrO2. The doping reaction is given in standard Kroger–Vink notation:

Ta2O5 þ V••O→

2ZrO2 2Ta•Zr þ 5OO ð9Þ

The concentration of oxygen vacancies in ZrO2 is reduced. A pre-vious study has shown that a lower concentration of oxygen vacan-cies decreases oxygen transport, thus reducing the oxidation rate ofTa-containing UHTC materials [15].

Completely oxidized layer: This layer is located in the oxygen-richregion and contains crystalline zirconia, tantalic oxide, and little silicate.

scale for ZST15 after thermal shock at 1600 °C.

515S. Wang et al. / Int. Journal of Refractory Metals and Hard Materials 41 (2013) 507–516

SiC is oxidized to SiO2 and CO through passive oxidation, as shown inEq. (2). ZrB2 is oxidized to ZrO2 and B2O3, as shown in Eq. (6). Glassphase fills the defects and boundaries. The gas dissolves in the glassphase and reaches the surface through the defects and boundaries be-tween the grains. TaSi2 particles are oxidized to Ta2O5. Ta+5 and Zr+4

have similar ionic sizes. Hence, the interaction between Ta2O5 and ZrO2

produces the Zr–Ta–O solid solution.Semi-oxidized layer: This layer is in an oxygen-poor region with

low oxygen pressure. SiC is oxidized to gaseous SiO and CO by activeoxidation, as shown in Eqs. (3) and (4). Part of the ZrB2 and TaSi2

Fig. 12. The compositional maps for Zr and O of the polished cross section of ZS, ZST5, ZST10,and (f); ZST15: (g) and (h).

particles is oxidized. Unchanged region: Oxygen cannot be trans-ferred in the unchanged region. Thus, oxidization does not occur inthis region.

Fig. 12 shows the compositional maps for Zr and O of the polishedcross section of ZS, ZST5, ZST10, and ZST15 after thermal shock at1600 °C. As Fig. 12(b), (d), (e), and (f) shows, the thickness values ofthe outer oxidized layer are estimated to be 32, 20, 12, and 10 μm, re-spectively. This result is consistent with the conclusion drawn fromFig. 8, which indicates that the addition of TaSi2 has a positive effecton the oxidation resistance of ZrB2.

and ZST15 after thermal shock at 1600 °C; ZS: (a) and (b); ZST5: (c) and (d); ZST10: (e)

516 S. Wang et al. / Int. Journal of Refractory Metals and Hard Materials 41 (2013) 507–516

4. Conclusions

ZrB2–SiC composite ceramics with different TaSi2 contents werefabricated by SPS. Adding TaSi2 particles significantly improves theinitial fracture toughness of the samples but simultaneously reducesthe initial strength at room temperature. The relative density of thesamples is also evidently improved. The addition of TaSi2 remarkablyenhances the thermal shock behavior of ZrB2–SiC ceramics. Theaddition demonstrates outstanding oxidation resistance and resid-ual mechanical properties, especially residual fracture toughness,at thermal shock temperatures between 1000 and 1600 °C. When theresidual fracture toughness of ZS is 2.72 MPa·m1/2 at 1000 °C and2.01 MPa·m1/2 at 1600 °C, the value for ZST15 reaches 8.91 MPa·m1/2

and 7.89 MPa·m1/2, respectively. The surface microstructure of TaSi2-added ZrB2 after thermal shock retains a denser surface than that ofZrB2–SiC, which is beneficial in obtaining a higher residual strengthafter thermal shock at 1000 °C to 1400 °C. When the temperature in-creases above 1400 °C, the Ta-borosilicate layer becomes incontinuousbecause of the evaporation of liquid and the vapor formed from oxida-tion. The residual fracture toughness increases by increasing the TaSi2content at each temperature. The toughening mechanism of ZSTx isbased on the contribution of the addition of a TaSi2 soft phase and acohesive structure resulting from the solid solution. An oxidation layeris notably generated after the specimen thermal shock. From the ther-modynamics viewpoint, the improved oxidation resistance of ZSTx isattributed to the addition of TaSi2, allowing the sample to develop aTa-borosilicate glass with higher viscosity and a decreased concentra-tion of oxygen vacancies in ZrO2.

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