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Materials Science and Engineering A 552 (2012) 125– 133

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A

jo ur n al hom epage: www.elsev ier .com/ locate /msea

park plasma sintering of silicon carbide and multi-walled carbon nanotubeeinforced zirconium diboride ceramic composite

ovindaraajan B. Yadhukulakrishnana, Arif Rahmana, Sriharsha Karumuria, Margaret M. Stackpooleb,. Kaan Kalkana, Raman P. Singha, Sandip P. Harimkara,∗

School of Mechanical and Aerospace Engineering Oklahoma State University, Stillwater, OK 74078, United StatesELORET Corporation, Moffett Field, CA 94035, United States

r t i c l e i n f o

rticle history:eceived 26 March 2012eceived in revised form 1 May 2012ccepted 8 May 2012vailable online 16 May 2012

a b s t r a c t

In this paper spark plasma sintering (SPS) of silicon carbide and multi-walled carbon nanotube reinforcedzirconium diboride ultra-high temperature ceramic matrix composites is reported. Systematic investiga-tions on the effect of reinforcement type (SiC and CNTs) and content (10–40 vol.% SiC and 2–6 vol.% CNTs)on densification behavior, microstructure development, and mechanical properties (microhardness, bi-axial flexural strength, and indentation fracture toughness) are presented. With the similar SPS processing

◦

eywords:interingeramicsirconium diborideilicon carbidearbon nanotubes

parameters (1900 C, 70 MPa pressure, and 15 min soaking time), near-full densification (>99% relativedensity) was achieved with 10–40% SiC (in ZrB2–SiC) and 4–6% CNT (in ZrB2–CNT) reinforced composites.The SiC and CNT reinforcement further improved the indentation fracture toughness of the compositesthrough a range of toughening mechanisms, including particle shearing, crack deflection at the particle-matrix interface, and grain pull-outs for ZrB2–SiC composites, and CNT pull-outs and crack deflection inZrB2–CNT composites.

. Introduction

Space vehicles re-entering the earth’s atmosphere experienceery high temperatures due to aerodynamic heating. Ultra-highemperature ceramics (UHTC) with melting point higher than200 ◦C are promising materials for thermal protection systemsf such space vehicles re-entering the earth’s atmosphere [1–7].mong several UHTC systems (TaC, HfC, NbC, ZrC, HfB2, ZrB2, andiB2), ZrB2 based ceramics are particularly important for thermalrotection systems. While the UHTCs exhibit very high meltingoints, the thermal conductivities of these materials are very low,aking dissipation of heat from the surface very difficult upon

hermal exposure. This often results in higher surface temper-tures and steep thermal gradients in the components, leadingo accelerated high temperature oxidation and fracture. Hence,ctual application of these materials for critical components such asharp leading edges in future generation hypersonic re-entry vehi-les would require significant improvements in their thermal andechanical properties, including thermal shock resistance, oxida-

ion resistance, and fracture toughness [8–11].Reinforcements are often added to ZrB2 to improve its den-

ification during sintering and also its mechanical, thermal, and

∗ Corresponding author. Tel.: +1 405 744 5830.E-mail address: sandip.harimkar@okstate.edu (S.P. Harimkar).

921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved.ttp://dx.doi.org/10.1016/j.msea.2012.05.020

© 2012 Elsevier B.V. All rights reserved.

oxidation properties. Several investigations have indicated that SiCreinforcement improves the strength and fracture toughness ofUHTC composites [12–16]. It also improves the oxidation prop-erties by the formation of layers of glassy SiO2 on the surface ofthe ceramic that provides passive oxidation protection at elevatedtemperatures [17]. Multi-walled carbon nanotubes (CNTs) exhibitexcellent mechanical and thermal properties making them attrac-tive candidates as nano-scale reinforcements in ceramic matrices[18–21]. CNTs have been successfully used as reinforcements instructural ceramics like alumina and silicon nitride [22–24]. CNTsimprove the fracture toughness of the composites through a rangeof toughening mechanisms like CNT pull-outs, crack bridging, andcrack deflection. While the toughening effects of CNT reinforce-ment in structural ceramics are now well established, these effectsare not well investigated for UHTC ceramics. Tian et al. fabri-cated ZrB2–SiC composites with and without CNT by hot pressing(1900 ◦C for 1 h). They reported about 15% increase in fracturetoughness with the CNT reinforcement [25]. However, no signif-icant improvement in hardness, strength, or thermal conductivitycould be achieved. This insignificant improvement in mechani-cal and thermal properties was attributed to partial degradationof CNTs and strong interfacial reactions between the CNTs and

ceramic matrix during hot pressing. Therefore there is a criti-cal need to study the survival and reinforcement effect of CNTsin UHTCs. Furthermore, the conventional hot pressing results insignificant grain growth during sintering. Recently spark plasma

1 Science and Engineering A 552 (2012) 125– 133

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26 G.B. Yadhukulakrishnan et al. / Materials

intering (SPS) is attracting significant interests for the rapid con-olidation of materials, including metals, alloys, and ceramics26–29]. The SPS process involves sintering under simultaneousnfluence of uniaxial pressure and pulsed direct current. While theature of exact mechanisms of SPS sintering are still being debated,he joule heating at the particle contacts and/or sparks at the gapsetween particles are often cited. These unique mechanisms allowintering at significantly lower temperature and in shorter sinter-ng time compared to conventional hot pressing. These attributesf the process are also attractive for the rapid sintering of UHTC-ased composites without significant grain growth and interfacialeactions. In this paper, SPS of ZrB2 based UHTC composites rein-orced with SiC and CNTs is reported. Systematic investigations onhe effect of reinforcement on densification behavior and mechan-cal properties (fracture toughness, hardness, and flexure strength)f the composites are presented.

. Experimental procedure

Commercially available ZrB2 powder (1–2 �m diameter; Alfaesar, USA), SiC powder (2 �m diameter; American Elements, USA),nd multiwalled CNTs (outer diameter of 30–50 nm; inner diam-ter of 5–15 nm; number of walls ∼23–67, length of 10–20 �m,anostructured and Amorphous Materials Inc, USA) were used as

tarting materials in this investigation. ZrB2, ZrB2–SiC (with 10, 20,0, and 40 vol.% of SiC), and ZrB2–CNT (with 2, 4, and 6 vol.% of CNT)omposite powder mixtures were prepared by dry ball milling for

min at 500 rpm with ball to powder weight ratio of 5:2. Tung-ten carbide (WC) jar and tungsten carbide balls were used forhe ball milling. The ball milling parameters were carefully cho-en to prevent any structural damage to the CNTs. All the ZrB2,rB2–SiC, and ZrB2–CNT composite powders were SPS sinteredt 1900 ◦C with uniaxial pressure of 70 MPa and soaking time of5 min under inert argon atmosphere (SPS system: Thermal Tech-ology LLC, USA). A heating rate of 100 ◦C/min was used during SPSrocessing. Graphite dies and punches were used for sintering disc-haped samples of 20 mm diameter and 2 mm thickness. The punchisplacement was continuously monitored during the SPS sinter-

ng cycle to investigate densification behavior. Bulk densities of theintered samples were measured using Archimedes’ principle. Theule of mixtures was followed to calculate the theoretical densitiesf the composites. X-ray diffraction (XRD) analysis was carried outn the sintered samples using Philips Norelco X-ray diffractome-er operating with Cu K� radiation (� = 1.54178 A). The diffractionngle (2�) was varied between 20 and 90◦ at a step increment of.02◦ with a count time of 1 s. Raman spectroscopy was carried outn fracture surfaces of ZrB2–CNT composites using WITec alpha300

Raman system (532 nm laser excitation, 0.8 mW laser power, and0 �m spot size) to identify the presence and structural quality ofNTs in the composites. The hardness of the samples was deter-ined using Vicker’s micro-indentation (Clark Instruments, USA;odel: CM-700AT) operated with normal force of 9.8 N and holding

ime of 15 s. The indentation fracture toughness of the samples wasalculated using diagonal crack lengths produced at the indentationorners from the microindentation tests. The fracture toughness KICs given by:

IC = 0.016(

E

H

)1/2 P

c3/2, (1)

here E is the Young’s modulus of the composites obtained usingule of mixture (Young’s moduli of ZrB2, SiC, and CNT are 500,75, 1000 GPa, respectively), H is the Vickers hardness (GPa), P

s the applied load (N), and c is the diagonal crack length (m)30]. The fracture toughness values are averaged for three sam-les, with five indents per sample, for each composition. The flexuretrength for the samples (20 mm diameter and 2 mm thickness) was

Fig. 1. Variation of relative density of (a) ZrB2–SiC, and (b) ZrB2–CNT compositeswith reinforcement content.

determined by a ring-on-ring (RoR) test method (Instron 5567,ASTM C1499-05). The support and loading ring diameters were15 mm and 5 mm, respectively, with a displacement controlledloading rate of 0.5 mm/min. The flexure strength �RoR is given by:

�RoR = 3P

2�t2

((1 − �)

(a2 − r2

)2R2

+ (1 + �) lna

r

), (2)

where P is the applied load (N), � is the Poisson’s ratio of the com-posites obtained using rule of mixture (the Poisson’s ratios of ZrB2,SiC and CNT are 0.15, 0.19 and 0.165, respectively), a is the radiusof the support ring, r is the radius of the load ring (m), and R andt are the radius and thickness (m) of the sample [31]. The fracturesurfaces of the samples were analyzed using SEM to investigatemicrostructure development and fracture behavior. Table 1 liststhe relative density and mechanical properties of the spark plasmasintered ZrB2, ZrB2–SiC, ZrB2–CNT samples.

3. Results and discussion

3.1. Relative density and microstructure

The variation of relative density of ZrB2–SiC and ZrB2–CNT com-posites with reinforcement content (SiC and CNT) is presented inFig. 1. The relative density of monolithic ZrB2 is also indicated.

For the given SPS processing parameters (sintering temperatureof 1900 ◦C, soaking time of 15 min, and pressure of 70 MPa), therelative density of monolithic ZrB2 was about 85%. For the simi-lar processing parameters, the ZrB2–SiC and ZrB2–CNT composites

G.B. Yadhukulakrishnan et al. / Materials Science and Engineering A 552 (2012) 125– 133 127

Table 1Relative density and mechanical properties of the spark plasma sintered ZrB2, ZrB2–SiC, ZrB2–CNT samples.

Sample Relative density (%) Hardness (GPa) Fracture toughness (MPa m1/2) Flexural strength (MPa)

ZrB2 84.8 16.64 ± 0.90 1.51 ± 0.02 162 ± 31ZrB2 + 10%SiC 99.1 19.38 ± 0.13 2.21 ± 0.25 553 ± 07ZrB2 + 20%SiC 99.6 20.80 ± 0.46 2.66 ± 0.34 301 ± 51ZrB2 + 30%SiC 99.7 21.44 ± 1.27 1.92 ± 0.32 431 ± 39ZrB2 + 40%SiC 99.7 22.71 ± 0.19 2.31 ± 0.23 410 ± 17

eactcdit4wot

Zpo

ZrB2 + 2%CNT 95.3 14.17 ± 0.38

ZrB2 + 4%CNT 99.3 16.39 ± 1.95

ZrB2 + 6%CNT 99.3 15.18 ± 0.40

xhibited much higher relative densities (>95%). In general, the rel-tive density of the composites increased with the reinforcementontent. The data on relative density showed relatively larger scat-er at lower reinforcement levels (10% SiC and 2% CNT). For ZrB2–SiComposites, the samples reached near-full densification (relativeensity >99%) at all reinforcement levels (10–40% SiC). While sim-

lar trend of densification was observed for ZrB2–CNT composites,he near-full densification was achieved at reinforcement levels of–6% CNTs. The relative density of ZrB2–CNT composite reinforcedith 2% CNTs was found to be about 95%. Clearly, very small amount

f CNT reinforcement is sufficient to achieve full densification inhese composites.

To understand the densification behavior of ZrB2–SiC andrB2–CNT composites during SPS processing, the data on punch dis-lacement was recorded during sintering. Fig. 2 shows the variationf punch displacement with sintering time during initial heating

2520151050-1

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Fig. 2. Ram displacement, temperature, and pressure profiles during heating and soak

1.53 ± 0.07 151 ± 072.62 ± 0.68 315 ± 653.53 ± 0.47 274 ± 20

and soaking stages of sintering cycles. The sintering temperatureand applied pressure during these sintering stages are also indi-cated in the figure. Note that the sintering cycle was initiated withsimultaneous increase in temperature (100 ◦C/min) and pressure(10 MPa/min). The sintering pressure of 70 MPa was reached in first8 min, and the sintering temperature of 1900 ◦C was reached infirst 17 min of the sintering cycle. The samples were soaked for15 min at the given sintering temperature of 1900 ◦C and pressureof 70 MPa (total heating and soaking time of 32 min). As showedin Fig. 2, the punch displacement showed three distinct trendsduring SPS sintering: initial increase corresponding to densifica-tion with pressure (stage I), intermediate decrease corresponding

to thermal expansion (stage II), and final increase due to densi-fication (stage III). As the early stages of sintering are dominatedby rearrangement of particles and plastic deformations, the stageI punch displacement (densification) seems to be primarily due

300

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ing stages of SPS sintering cycles for (a) ZrB2–SiC, and (b) ZrB2–CNT composites.

128 G.B. Yadhukulakrishnan et al. / Materials Science and Engineering A 552 (2012) 125– 133

Fig. 3. SEM micrographs from the fracture surfaces of (a) ZrB2, (b) ZrB2 + 10% SiC, (c) ZrB2 + 20% SiC, and (d) ZrB2 + 40% SiC composites.

Fig. 4. SEM micrographs from the fracture surface of (a) ZrB2, (b) ZrB2 + 2% CNT, (c) ZrB2 + 4% CNT, and (d) ZrB2 + 6% CNT composites.

Science and Engineering A 552 (2012) 125– 133 129

tsewdfridippfppeddrITto(c∼paicZmcttpfitdwFZtsd(ssttgtFn(siswZnesfiCw

G.B. Yadhukulakrishnan et al. / Materials

o these effects. The extent of punch displacement during thistage I is very small (<0.25 mm) for pure ZrB2 samples. In gen-ral, the extent of punch displacement during this stage increasesith increasing reinforcement content (both SiC and CNTs). Punchisplacements of ∼1.25 mm for ZrB2–SiC (40% SiC) and ∼0.5 mmor ZrB2–CNT (6% CNTs) composites were observed. Clearly, theeinforcing phases (SiC and CNTs) facilitate the densification dur-ng stage I. While direct observation of exact mechanisms of suchensification behavior is difficult, it seems that effects are primar-

ly due to size difference between matrix phase and reinforcinghase (for ZrB2–SiC and ZrB2–CNT composites), and lubricatingroperties of CNTs (for ZrB2–CNT composites). These size dif-erence and lubricating effects in composites seem to facilitatearticle re-arrangement and better densification during stage I. Theunch displacement during stage II decreases indicating thermalxpansion of packed/locked powder mixtures. Finally, the punchisplacement increases in early stage III indicating later stages ofensification. The transition points between stages II and III cor-espond to punch displacement minima. For pure ZrB2, this stageI–III transition corresponds to starting of soaking stage (t = 17 min;

= 1900 ◦C). For ZrB2–SiC and ZrB2–CNT composite samples, thisransition point shifts to left (earlier times) in the heating stagef the SPS sintering cycle (indicated by dotted line). For ZrB2–SiC10–40% SiC) composite samples, this stage II–III transition pointorresponds to about t = 11–15 min (corresponding temperature of1300–1700 ◦C in heating cycle). For ZrB2–CNT (2–6% CNT) com-osite samples, this stage II–III transition point corresponds tobout t = 12–15 min (corresponding temperature of ∼1400–1700 ◦Cn heating cycle). This shift of stage II–III transition point indi-ates that final densification is favored at lower temperatures forrB2–SiC and ZrB2–CNT composite samples. The punch displace-ent increases in early stage III (t = 18–19 min) without significant

hange during remaining stage III (soaking period correspondingo t > 20 min). Clearly, the punch displacement data indicate thathe densification is favored for the ZrB2–SiC and ZrB2–CNT com-osite samples in early stage I (particle re-distribution stage) andnal stage III (densification starting at earlier times and at loweremperatures). The experimental data indicating higher relativeensity for ZrB2–SiC and ZrB2–CNT composites are in agreementith these general observations made from punch displacements.

ig. 3 shows the SEM micrographs from the fracture surfaces ofrB2 and ZrB2–SiC composites. The pure ZrB2 exhibits well dis-ributed porosity with the grain size equivalent to starting particleize (Fig. 3(a)). The fracture surfaces from ZrB2–SiC samples showense faceted grain structure with well distributed SiC particlesFig. 3(b–d)). The microstructural observations from the fractureurfaces are in general agreement with the measured relative den-ity values for these ZrB2 and ZrB2–SiC composite samples. Whilehe distributed larger grains with grain size about <5 �m appear inhe microstructures of the ZrB2–SiC composites, significant grainrowth was not observed. The SEM micrographs from the frac-ure surfaces of ZrB2 and ZrB2–CNT composites are compared inig. 4. The ZrB2–CNT composites with 2% CNTs shows intercon-ected porosity on the fracture surface with slightly larger grains2–3 �m) than starting particle size (Fig. 4(b)). The micrograph alsohows regions of necking between adjacent particles. The compos-tes with higher CNT content (4 and 6%) showed dense fractureurface with distributed networks of CNTs. Significant grain growthas observed in these ZrB2–CNT composites (4 and 6% CNTs) with

rB2 grain size in the range of 5–10 �m. The distribution of CNTetworks both inside the grains and at the grain boundaries can beasily seen. The fracture surface shows faceted grain structure with

hear-band like fracture features on the ZrB2 grains. A high magni-cation micrograph clearly shows the CNT networks and protrudedNTs in the ZrB2–CNT composite sample (Fig. 5). Clearly, the CNTsere retained in the composites processed using SPS. However,

Fig. 5. A high magnification SEM micrograph from the fracture surface ofZrB2–2 vol.% CNT composite showing CNT networks at the grain boundaries andinside the grain.

the SiC appear to be much more effective than CNTs in restrictingthe grain growth during SPS processing. The rapid heating rate andshorter soaking time involved in SPS ensured insignificant graingrowth and better densification, and eliminated the risk of CNTdegradation or strong interfacial reactions between the CNTs andceramic matrix as encountered by Tian et al. [25].

3.2. X-ray diffraction (XRD) and Raman spectroscopy analysis

XRD patterns from SPS sintered ZrB2, ZrB2–SiC, and ZrB2–CNTcomposites are presented in Fig. 6. For ZrB2–SiC composites, all thecharacteristic peaks of ZrB2 and SiC were identified. No additionalpeaks were observed suggesting sintering of two phase mixtureswithout any undesirable interfacial reactions. For ZrB2–CNT com-posites, all the characteristic peaks of ZrB2 were identified. Notethat characteristic (0 0 2) peak (2� = 26.2◦) corresponding to CNTscould not be found in the XRD patterns from the ZrB2–CNT com-posites. This is probably due to very low volume fractions of CNTsin the ZrB2 ceramic matrix.

In order to understand the chemical and structural changes inthe CNTs during the SPS process, we performed a detailed char-acterization of the CNT networks in the ZrB2–CNT compositesusing Raman spectroscopy. The MWCNTs used in this work aretoo large in diameter to observe Radial Breathing Mode (RBM) fea-tures and hence, we focused on the characteristic Raman peaks ofsp2 hybridized carbon namely, D (∼1350 cm−1), G (∼1580 cm−1), D′

(1620 cm−1), and G′ (∼2700 cm−1). The G and D are associated withthe in-plane stretching of C–C and breathing modes, respectively,whereas G′ and D′ correspond to their respective second orderphonon counterparts [32–34]. The D and D′ are not observed in theperfect graphite and they become active in the presence of defectsand therefore these peaks are indicative of disorder in the sp2 car-bon materials. Fig. 7 depicts the Raman spectra of ZrB2–CNT (2, 4,and 6%) composites compared with ZrB2, pristine MWCNTs, andsintered MWCNTs. ZrB2 has no active Raman modes [35] and thusthe spectra show the characteristic peaks of CNTs in the compos-ites. It is evident from the spectra that the D, G, and G′ bands have

shifted to higher energies in case of sintered CNTs and ZrB2–CNTcomposites. This high energy shift can be owed to: (i) decreasein the average distance between defects [32–34]; (ii) the residualcompressive stress on the CNT network imposed by the ceramic

130 G.B. Yadhukulakrishnan et al. / Materials Science and Engineering A 552 (2012) 125– 133

9080706050403020

a

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ig. 6. XRD patterns from spark plasma sintered (a) ZrB2–SiC and (b) ZrB2–CNTomposites.

atrix [36–39] that evolves during the cooling step (i.e., thermalontraction of ZrB2 matrix surrounding CNTs).

The ratio of intensities of D to G (ID/IG) is another way to quan-

ify the density of defects or the degree of chemical damage inhe CNTs [32,34]. From Fig. 7, it is obvious that the ID/IG has notaried significantly compared with pristine CNTs implying there iso significant defect formation during the SPS process. Indeed, the

300027001600140012001000

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G

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Dl

ig. 7. Raman spectra from ZrB2–CNT composites showing in the frequency ranges000–1750 cm−1 and 2450–3000 cm−1.

Fig. 8. Hardness of (a) ZrB2–SiC and (b) ZrB2–CNT composites as a function of rein-forcement content.

persistence of the G′ band during the SPS process indicates struc-tural consistency of sp2 carbon [34]. Our highest ID/IG value is0.61 (6% CNT). For this value of ID/IG, the G peak was found at1565 cm−1 for ion-bombarded glassy carbon [34]. On the otherhand, our G peak is found at 1589 cm−1 (being significantly higherthan 1565 cm−1) that suggests the high energy shift in our G peak ismainly due to stress. This argument is also supported by the Ramanspectrum of sintered CNTs in Fig. 7, where ID/IG is 0.81 but G peakis only at 1585 cm−1(in agreement with [34]). In addition, the shiftin the peaks is higher in case of 4 and 6% CNT than 2% CNT compos-ites. We explain this observation by higher stress formation in the4 and 6% CNT cases due to denser ceramic matrix formation thatinhibits stress relaxation. Therefore, our Raman findings corrob-orate that the structure of CNTs remain intact with insignificantdefect formation during SPS process. Further, we infer develop-ment of compressive stress in the CNTs possibly due to thermalcontraction of the ceramic matrix during cooling.

3.3. Mechanical properties

Fig. 8 shows the variation of hardness of ZrB2–SiC and ZrB2–CNTcomposites with the reinforcement content. The average hard-ness of SPS sintered, with relative density of ∼85%, was found tobe 16.64 GPa in this investigation. The average hardness of theZrB2–SiC composites increased with increasing SiC reinforcementcontent, reaching ∼22.71 GPa for composites with 40% SiC. The

higher hardness for the ZrB2–SiC composites compared to ZrB2seems to be due to combination of effects such as higher relativedensity and presence of high volume fraction of harder rein-forcement (SiC) particles. As all the ZrB2–SiC composites (10–40%

G.B. Yadhukulakrishnan et al. / Materials Science and Engineering A 552 (2012) 125– 133 131

0

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toughening mechanisms (SiC grain pull-outs and crack deflec-

ig. 9. Flexural strength of (a) ZrB2–SiC, and (b) ZrB2–CNT composites as a functionf reinforcement content.

iC) showed comparable relative density (>99%) and grain size2–5 �m), the increasing hardness with SiC content seems pri-

arily due to composite strengthening. On the other hand, thetrengthening effect of CNTs was not evident in ZrB2–CNT com-osites even though the CNT reinforcement resulted in significant

mprovement in the densification (relative density >95%). Note thatNTs are distributed in the ZrB2 matrix in the form of dispersedetworks instead of individual CNTs. Furthermore, the ZrB2–CNTomposites also showed significant grain growth. The combinationf these competing effects (densification, non-uniform CNT distri-ution, and grain growth) seems to have resulted in insignificantr no strengthening in these ZrB2–CNT composites. The SPS sin-ered ZrB2 samples had biaxial flexure strength of 162 MPa (Fig. 9).he flexural strength of the ZrB2–SiC and ZrB2–CNT compositesas also improved with increasing reinforcement content (SiC andNT). The improvements in flexural strength are more pronouncedith the SiC reinforcement with the average flexural strength in

he range of ∼300–550 MPa. The ZrB2–SiC composites with 10% SiCxhibited highest average flexural strength of 553 MPa. The aver-ge flexural strength for ZrB2–CNT composites was in the rangef 150–315 MPa, with composites with 4% CNTs exhibiting highestverage flexural strength of ∼315 MPa.

Significant improvement in fracture toughness was observedn the ZrB2–SiC and ZrB2–CNT composites compared to ZrB2Fig. 10). The average fracture toughness of SPS sintered ZrB

2as 1.5 MPa m1/2. This is relatively lower than fracture tough-ess (3–4 MPa m1/2) reported for dense monolithic ZrB2 (relativeensity >99%). The observed lower fracture toughness of ZrB2

Fig. 10. Fracture toughness of (a) ZrB2–SiC, and (b) ZrB2–CNT composites as a func-tion of reinforcement content.

sintered in this investigation seems primarily due to lower rel-ative density (∼85%). While the SiC reinforcement resulted inthe improvement in fracture toughness of ZrB2, the average frac-ture toughness values of all the ZrB2–SiC composites remainedin the range of 2–2.7 MPa m1/2. Among ZrB2–SiC composites, thecomposites reinforced with 20% SiC exhibited highest fracture of2.66 ± 0.34 MPa m1/2. For ZrB2–CNT composites, the fracture tough-ness also increased with increasing CNT reinforcement content andwas in the range of 1.5–3.5 MPa m1/2. The ZrB2–CNT compositesreinforced with 6% CNTs exhibited highest fracture toughness of3.53 ± 0.47 MPa m1/2. Fig. 11 shows the morphologies of indenta-tion cracks propagated through the SPS sintered ZrB2, ZrB2–SiC, andZrB2–CNT composites. In ZrB2 samples, the crack front was rela-tively straight without significant bending or deflection (Fig. 11(a)).However, the crack seems to propagate along boundaries of smallergrains (1–2 �m). The fracture surface of ZrB2 samples indicatedpredominantly intergranular fracture features with distributedporosity (Fig. 3(a)). For ZrB2–SiC composites, the indentation crackinteracts with the reinforced SiC particles causing shear of the finerparticles and crack getting deflected at the coarser particles. Thefracture surface of the ZrB2–SiC composites also shows predomi-nantly intergranular fracture features with distributed depressionsdue to grain pull-outs (Fig. 3(b–c)). The relatively higher frac-ture toughness of these ZrB2–SiC composites compared to ZrB2samples seems to be due to better densification and additional

tion). For ZrB2–CNT composites, the propagating crack encountersthe networks of CNTs and goes around the networks (Fig. 11(c)).The fracture surface shows both intergranular (faceted surface

132 G.B. Yadhukulakrishnan et al. / Materials Scienc

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Kiricsi, P. Arato, Compos. Sci. Technol. 65 (5) (2005) 727–733.

ig. 11. High magnification crack propagation and toughening mechanism of (a)onolithic ZrB2, (b) ZrB2–20% SiC, and (c) ZrB2–4% CNT.

eatures) and transgranular (shear-bands or river-marks type fastracture features) features with voids corresponding to pulledut CNT networks (Fig. 4(c)). The high magnification image fromrB2–CNT composites clearly showed pulled out CNTs and CNT net-

orks both at the grain boundaries and inside the grains. As the

rB2–CNT composites with very low level of reinforcement (up to% CNTs) exhibited higher fracture toughness than that for ZrB2–SiC

[

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e and Engineering A 552 (2012) 125– 133

composites (with up to 40% SiC), it seems that CNTs are more effec-tive in toughening the ZrB2-based UHTCs. Better dispersion of CNTsin the UHTC matrix is further expected to improve the tougheningeffects.

4. Conclusions

Spark plasma sintering was successfully used to sinterfine-grained ZrB2–SiC and ZrB2–CNT composites to near-full den-sification (>99% relative density). Detailed analysis of punchdisplacement during SPS sintering indicated that the densificationis favored for the ZrB2–SiC and ZrB2–CNT composite samples inearly stage I (particle re-distribution stage) and final stage III (densi-fication starting at earlier times and at lower temperatures). Whilethe fine grains (<2 �m) were retained in ZrB2–SiC composites, somegrain growth with grain size in the range 5–10 �m was observedfor ZrB2–CNT composites. Raman spectroscopy analysis and highmagnification SEM micrographs indicated that the multi-walledCNTs were retained in ZrB2–CNT composite samples sintered athigh sintering temperature of 1900 ◦C. The SiC and CNT reinforce-ment further improved the indentation fracture toughness of thecomposites through a range of toughening mechanisms, includingparticle shearing, crack deflection at the particle-matrix interface,and grain pull-outs for ZrB2–SiC composites, and CNT pull-outs andcrack deflection in ZrB2–CNT composites.

Acknowledgments

SPH acknowledges support from the Oklahoma Space GrantConsortium through NASA-EPSCoR Research Initiation Grants pro-gram. RPS acknowledges partial student support through the C.F.Colcord Chair Professorship endowment.

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