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Int. Journal of Refractory Metals & Hard Materials 27 (2009) 367–375

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Int. Journal of Refractory Metals & Hard Materials

journal homepage: www.elsevier .com/locate / IJRMHM

The synthesis and consolidation of hard materials by spark plasma sintering

Dustin M. Hulbert, Dongtao Jiang, Dina V. Dudina, Amiya K. Mukherjee *

Department of Chemical Engineering and Materials Science, University of California, One Shields Avenue, Davis, CA 95616, USA

a r t i c l e i n f o

Article history:Received 26 March 2008Accepted 5 September 2008

Keywords:Spark plasma sinteringBoron carbideHafnium diborideSilicon carbideFunctionally graded materials

0263-4368/$ - see front matter � 2008 Elsevier Ltd. Adoi:10.1016/j.ijrmhm.2008.09.011

* Corresponding author. Tel.: +1 530 752 1776; faxE-mail address: [email protected] (A.K. M

a b s t r a c t

Dense ceramic matrix composites consisting of boron carbide, titanium diboride, hafnium diboride weresynthesized by spark plasma sintering (SPS). Specifically, a functionally graded boron carbide–aluminumcomposite with a precipitous microstructural gradient was created by a unique SPS die/punch off-settechnique. Additionally, pseudo eutectic titanium diboride–boron carbide composites were synthesizedby SPS using mechanically alloyed elemental powders. Lastly, hafnium diboride–20 vol.% silicon carbidecomposites were created from high energy ball milled precursor powders. This work demonstrates theSPS equipment’s promising potential to synthesize bulk, hard, refractory composites.

� 2008 Elsevier Ltd. All rights reserved.

1. Introduction

Composites such as boron carbide–aluminum, titanium dibo-ride–boron carbide and hafnium diboride–silicon carbide areattractive materials due to their hardness and refractory proper-ties. Unfortunately, the very properties that make these materialsattractive for design purposes make them unattractive for process-ing reasons. Specifically, these materials are traditionally very dif-ficult to consolidate [1–6]. Fortunately, spark plasma sintering(SPS) has been shown to be a very effective means of consolidatingthese borides and carbides [7–10].

SPS uses moderate uni-axial pressures (typically less than100 MPa) and an on–off DC pulsing current to sinter powders. Dur-ing this on–off DC pulsing, there are a number of proposed mech-anisms to account for the enhanced sintering behavior. Many, butnot all, of these mechanisms assume the presence of a momentar-ily generated spark plasma between particles and include: sparkimpact pressure [11,12], plasma cleaning of particle surfaces[13,14], joule heating [12–14], local melting and/or evaporation[11], particle surface activation, electromigration [15,16] and fieldassisted diffusion [14,16,17]. Of the mechanisms listed above, Jouleheating, local melting and evaporation (especially in metallic sys-tems) and electromigration have been relatively well established.The other mechanisms, especially those that invoke the presenceof plasma, are highly hypothetical in nature. Other than sintering,SPS technology has been used to join materials, grow crystals, facil-itate advanced chemical reactions and even form materials into

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: +1 530 752 9554.ukherjee).

different shapes [11,19,20]. For more information regarding theSPS process, Munir et al. [16] provide an excellent review.

The purpose of this paper is not to discuss the mechanisms ofthe SPS process, but rather to show its utility in processing a vari-ety of refractory and hard materials. The specific advantages, dis-advantages and background regarding each system are discussedbriefly below.

1.1. Boron carbide–aluminum composites

Boron carbide (BxC, x P 4) is an attractive material for manyreasons. It has low density (2.51 g/cm3), excellent chemical resis-tance and it is extremely hard [2]. Unfortunately, monolithic B4Cis highly brittle, especially when shock loaded at high strain rates[18]. It is known that combining the B4C with a metal can mitigatethe problems associated with brittleness. By far, the most popularmetal used in conjunction with B4C for this purpose is aluminum[6,19,20]. Aluminum is lightweight, readily available, cost-effectiveand wets B4C well at elevated temperatures [19]. Molten alumi-num has been shown to form a variety of binary and tertiaryphases when in contact with B4C such as Al3BC, AlB10 and Al4C3

[6,19,20]. Al3BC is the most commonly observed phase through avariety of melt infiltration temperatures and helps bond the metalto the carbide matrix.

In the past many different techniques have been employed tocreate functionally graded materials (FGMs). The technologies usedto create FGMs have developed considerably over the last 35 years[21,22]. Kieback et al. [23] provide an excellent review on the var-ious processing techniques employed to create FGMs. The idea ofusing SPS to create FGMs is a relatively new idea and little workhas been done. Feng et al. [24] used SPS to create a Ti–TiB2–B

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368 D.M. Hulbert et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 367–375

FGM using a step-wise method. However, problems with delami-nation due to dimensional instability between layers can be aproblem with step-wise FGMs. In this study, SPS was used to createFGMs with a continuous microstructure and precipitous propertygradients.

1.2. Titanium diboride–boron carbide composites

Titanium diboride has several attractive properties. It has highhardness, a high melting temperature (3225 �C), low density(4.5 g/cm3) and good thermal (96 W/m/K) and electrical(22 � 106 X cm) conductivity [8,25]. Moreover, titanium diboridehas been proposed as a promising second phase for tailoring themechanical and functional properties of boron carbide [26–30].Pairing boron carbide and titanium diboride in a composite canbe beneficial for the design of thermoelectric materials [29] andcomposites with increased fracture toughness and bendingstrength relative to pure boron carbide [26–28,30,31]. Titaniumdiboride was also suggested as a possible grain growth inhibitor[27] for boron carbide. Several attempts have been made to devel-op boron carbide–titanium diboride composites by two main ap-proaches. The first tact involves sintering boron carbide andtitanium diboride powders directly [29,31]. The second methoduses in situ reactive sintering of boron carbide with additions ofTi, TiO2 [31] or TiO2 and carbon [28,30]. Generally speaking, thein situ reactive sintering method provides the best results in termsof final density.

Building on this idea, the present authors suggest a new synthe-sis route for boron carbide–titanium diboride composites based onthe in situ formation of both phases during SPS. This method pro-vides more flexible microstructural control of the final composite.As confirmation of this approach, a study conducted by Anselmi-Tamburini et al. [10] showed significant improvement in the finaldensity of hafnium diboride while sintering the material from ele-mental powders using SPS technology.

1.3. Hafnium diboride–silicon carbide composites

Ultra-high-temperature-ceramics (UHTCs) are a class of cera-mic materials that exhibit extremely high melting temperatures(>3000 �C) [25,32,33]. While there are several material systemsthat can be classified as UHTCs, HfB2 and ZrB2 based systems seemto be the most common chemistries studied [1,9,10,25,32,33]. Dueto their high melting temperatures and relative chemical inertnessat elevated temperatures UHTCs are often used as crucibles, elec-trodes in arc furnaces and as furnace linings.

Recently, HfB2 and ZrB2 have been attracting attention as poten-tial thermal protection materials for hypersonic aircraft. Oxidationcan be an issue for both pure HfB2 and ZrB2 at temperatures as lowas 1200 �C. This oxidation is due to the formation of volatile phaseslike B2O3 [34]. Fortunately, it is known that the addition of 20–30vol.% of SiC greatly mitigates the oxidation issue through the for-mation of a borosilicate glass surface layer [34,35]. In the presentwork, the authors synthesized a near fully dense HfB2–20 vol.%SiC composite using HEBMed starting powders in conjunction withSPS.

2. Experimental procedures

2.1. Powder preparation

2.1.1. Boron carbide–aluminum compositesAmorphous boron (95–97% pure, Atlantic Equipment Engineers,

USA) and amorphous carbon (99.99% pure, Alpha Aesar, USA) pow-ders with a mean particle size of 100 and 42 nm, respectively were

mixed in B4C stoichiometric ratios in �38 g batches to produce thestarting powder. Mixing was accomplished by combining the pow-ders with approximately 2 L of high purity ethanol then ultrasoni-cating the slurry for 60 min. The slurry was then dried while beingstirred on a hot plate. The dried slurry was then crushed and sievedto 150 lm before being loaded into graphite SPS dies.

2.1.2. Titanium diboride–boron carbide compositesTitanium (99.5%, �325 mesh, Alfa Aesar), amorphous boron

(95–97%, submicron, Fluka, Sigma–Aldrich) and carbon black(99.99%, 42 nm, Alfa Aesar) powders were used as the raw materi-als. The contents of the elemental powders in the mixtures werecalculated according to 23 vol.% TiB2 and 77 vol.% B4.5C in the com-posite product. This ratio of TiB2 to B4.5C is known to be a pseudo-eutectic composition. It is well known that boron carbide has abroad non-stoichiometry range (9–20 at.% C). The boron carbidesynthesized in this work is close to the congruently melting com-position [2].

Mechanical milling was conducted in a SPEX 8000 mill (SPEXCertiPrep, USA) with a tungsten carbide ball and a tungsten carbidevial. The O-ring seals were used to firmly close the vial in order toprevent any ingress of air during milling. The powder mixture loadin the milling vial was 2.5 and 5.0 g. The weight of the tungstencarbide ball was 10 g. The milling time was 4, 8 and 16 h. Themechanically milled powders were cold pressed at a pressure of10 MPa, the green density of the samples being 50–55% of the the-oretical density of the two-phase boron carbide–titanium diboridecomposite.

2.1.3. Hafnium diboride–silicon carbideThe elemental powder of B metal (99%, -325 mesh, CERAC) and

Hf (99.8%, -325 mesh, CERAC) together with b-SiC (>97%, 20–30 nm, Nanostructured & Amorphous Materials Inc.) were high en-ergy ball milled (HEBM) for 16 h.

2.2. Spark plasma sintering

To make FGM, 3.2 g powder charges were loaded in an offsetmanner as shown in Fig. 1a. In order to prevent the powder fromreacting with the graphite surface of the die and punches graphitefoil was used. Prior to SPS the specimens were cold pressed in thedie using a uniaxial pressure of 48 MPa. The dies used were38.0 mm in height and had an outside diameter (O.D.) of44.4 mm. The punches were 25.4 mm in height and had an O.D.of 19.1 mm. The tolerance between the die and punch was lessthan 25 lm. The samples were then subjected to SPS using the fol-lowing parameters using a Dr. Sinter 1050 (Sumitomo Coal MiningCo., Japan) SPS apparatus.

� Heated from room temperature to 600 �C in 3 min.� Held at 600 �C for 3 min.� Heated from 600 �C to 1660 �C in 2 min.� Power shut off.

Pressure was held at approximately 53 MPa during sinteringand rapidly removed immediately following power shut down.The direct current pulsing duty cycle was set at 12 pulses on and2 off. Temperature was monitored using an optical pyrometer fo-cused on a near through hole (Fig. 1a).

Consolidation of titanium diboride–boron carbide compositeswas performed in the same SPS equipment as above. The heatingrate was 200 �C/min from room temperature up to 600 �C and thenthe temperature was increased with a rate of 100–200 �C/min untila maximum temperature was reached. As before, the direct currentpulsing cycle was set at 12 pulses on and 2 off. The applied voltageat the maximum temperature was 4.3–5.5 V. The maximum

Fig. 1. (a) A sketch showing the SPS die and punch set up used to create the B4C/FGM. (b) A photograph showing the gradient in color and intensity of light emitting from thedie and punches.

D.M. Hulbert et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 367–375 369

current reached 1300 A when the sample was heated up to 1200 �Cand 1600–2000 A when the temperature increased up to 1700 �C.The holding time varied from 2 to 5 min at the maximum temper-ature. Graphite dies and punches were used. The weight of thepowder loaded in the die was 0.85 g. The powders were loaded intothe dies as before. A unidirectional pressure of 50–100 MPa wasapplied during the SPS process. Temperature was monitored usingan optical pyrometer focused on a near-through hole in the wall ofthe die.

Spark plasma sintering of hafnium diboride–silicon carbide wasconducted with a load of 80 MPa, a heating rate of 200 �C/min, andhold at 1800 �C for 8 min.

2.3. Melt infiltration

Following SPS the partially porous compacts were infiltratedwith 99.999% pure molten aluminum (Electronics Space ProductsInternational, USA). The boron carbide was placed in a graphitecrucible with the porous side facing up. Aluminum shot(�20 mm3 per piece) was then placed on the boron carbide. Infil-tration was accomplished in a vacuum of 10�3 to 10�4 Pa. The spec-imen was heated from room temperature to 500 �C at a heatingrate of 60 �C/min and held between 500 and 600 �C for 30 min. Thiswas done to remove any absorbed moisture and other volatilecompounds as well as to stabilize the vacuum. Heating then con-tinued at a rate of 60 �C/min to1180 �C where temperature washeld constant for 10 min. Furnace power was then shut off andthe specimen was allowed to cool in vacuum.

2.4. Characterization

Archimedes’ method was employed to measure the density ofthe sintered samples with deionized water as the immersion li-quid. Due to the cylindrical geometry of the specimens the densi-ties of the newly formed FGMs were calculated by dividing massover volume. The newly formed FGM were sectioned using a dia-mond saw and polished to a 1 lm diamond finish. Fracture sur-faces were obtained by cleaving a partial diamond saw cut inmode I fracture. Macroscopic images were obtained using a digitalcamera (Olympus, USA) coupled to a stereoscope. X-ray diffraction(XRD) phase analysis of the powder mixtures and the sinteredsamples was performed with a XDS 2000 X-ray diffractometerusing Cu Ka radiation (Scintag Inc., USA). Microstructural observa-tion (including grain size approximation), backscatter electron(BSE) images and energy dispersive spectroscopy (EDS) were car-

ried out using an FEI XL30-SFEG (FEI, USA) field emission high-res-olution scanning electron microscope (SEM). High-resolutionimaging using secondary electrons was done at 5–10 kV whileEDS and BSE imaging was performed at 20 kV.

Hardness and toughness measurement were done on a WilsonTukon (American Chain & Cable Company Inc., USA) hardness tes-ter with a Vickers diamond indenter using a 2.5 kg load with adwell time of 15 s. The fracture toughness of the fully dense re-gions of the FGM was calculated with the Antis equation usingan elastic modulus value of 410 GPa [2,36]. The elastic modulusof the composite was calculated according to the rule of mixturesbased on the Voigt model. The values of 465 and 565 GPa for theelastic modulus of boron carbide and titanium diboride, respec-tively, were taken for the calculation.

3. Results and discussion


Following mixing, drying, crushing and sieving the amorphousB and C powder mixture had a very fine and fluffy texture. Duringthe SPS process, the die and punches showed a very noticeable dif-ference in color and intensity of light as seen in Fig. 1b. One punchand side of the die was a dull red color with the opposite side of thedie and opposing punch being white-hot in color. Rapid heatingand cooling is necessary in order to preserve the temperature gra-dient in the specimen and the surrounding graphite. It was foundthat a hold time of even 1 min destroyed any microstructure gradi-ent; this result is consistent with previous modeling results [14].The density of the porous compacts after SPS varies somewhatbut was typically found to be near 70% of the theoretical value of2.51 g/cm3. After SPS one side of the compact was fully dense withthe opposite side being highly porous.

One problem in creating Al–B4C cermets is the delicate balanc-ing act of finding favorable temperatures and times to facilitatecomplete infiltration without growing other phases such as Al4C3.The conditions used in this study are the same as those found inwork by Halverson [37] and are designed to maximize Al infiltra-tion while minimizing the formation of other phases. Followingmelt infiltration the density of FGMs varied slightly but was foundto be near 90% of the theoretical value based on the rule of mix-tures. Most of the porosity is found in a region of closed poreporosity near the dense, B4C rich side of the FGM.

X-ray diffraction data (Fig. 2) of the starting powder andthe FGM compact before and after Al infiltration yielded several

Fig. 2. XRD pattern of the starting powder, the FGM after SPS and the specimen following melt infiltration at 1180 �C for 10 min.


important results. Firstly, the starting powder is amorphous. Thisfact coupled with extremely fast heating rates in the SPS, explainthe submicron nature of the B4C grains in the more porous regionsas seen in Fig. 3. B4C grains in the dense region are approximatelyone order of magnitude larger. From analyzing the fracture surfaceof the pre and post infiltrated specimens the smallest B4C grainswere equiaxed and approximately 500 nm or less while the largestgrains appear to be approximately 5 lm or larger.

Also apparent from the XRD pattern of the pre-infiltrated B4C isthe presence of microstructural defects, specifically twining. This isevident by the presence of the broad halos between the doubletsapproximately centered on the 2h values of 23�, 36� and 64.5�. Ans-elmi-Tamburini et al. [38,39] found similar defects both throughTEM investigation and through XRD modeling studies. Such defectscould be from the rapid thermal cycling common in SPS processing.

Fig. 3. The left image shows the dense side of the FGM consisting of 100% B4C with grainon the right shows the porous B4C region prior to melt infiltration. Grains are equiaxed

Fig. 4. A BSE montage image of the macrostructure of the FGM. The lighter regions

After infiltration the defect concentration appears to decrease asapparent through the reduction in the halos previously mentionedand the increase in intensity of the main B4C peaks at 24�. Thisreduction in defect concentration is most likely due to the anneal-ing affect of the melt infiltration process.

The microstructural evolution and extent of infiltration of Al inthe FGM can be seen in Fig. 4. This image is a back scattered elec-tron (BSE) microscopy montage image of an entire cross section ofthe B4C–Al cermet. The lighter regions in the BSE image correspondto areas with more Al while the darker regions are B4C rich. FromFig. 4, the depth of penetration of the molten Al can be approxi-mated to be between 3 and 4 mm.

One convincing piece of evidence that points to the effective-ness of this processing route to produce quality FGMs can be foundin Fig. 5. Here we see an increase in hardness in the pre-infiltrated

s on the order of 10 lm. Arrows point to possible twin defects. The SEM micrographin nature and submicron in size.

contain more Al than the material on the far right is fully dense boron carbide.

Hardness vs. Position of Functionally Graded B C 4

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Fig. 5. The hardness profile of the B4C–Al cermet before and after Al melt infiltration.


specimen hardness from 1.74 to 32.3 GPa. The difference in hard-ness after infiltration is slightly less, varying from 8.36 to32.38 GPa. The elevated hardness in the post infiltrated specimenopposite the dense side is most likely due to the reinforcing effectthe aluminum has on preventing pore collapse. To the authors’knowledge, hardness gradients of this magnitude have not beenachieved before in a continuous or non-layered FGM. The fracturetoughness of the dense side has found to be 3.3 MPa/m�1/2.

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XRD Patterns of TiB2 - B4C

Fig. 6. XRD patterns of the Ti–B–C powder mixtures mechanic

3.2. Titanium diboride–boron carbide

XRD patterns of the Ti–B–C mechanically milled mixtures areshown in Fig. 6. The reflections observed in the 17–21� range aredue to the presence of some crystalline boron as an admixture inthe amorphous boron powder. The presence of tungsten carbideis due to contamination from the milling media and the vial walls.As was expected, the amount of tungsten carbide increases with

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ally milled for various time intervals (powder load 5.0 g).


increasing milling time. Additionally, cubic boron oxide appears inthe mixture upon milling for prolonged periods of time. As isapparent from the XRD patterns, neither the titanium diboride,or the boron carbide phase forms during the milling process.

The powder composition studied presently can be described asa stoichiometric mixture of titanium and boron taken to producetitanium diboride. However, the solid state reaction is highly di-luted by the boron and carbon black used to form the boron car-bide phase. Thus, despite titanium diboride’s tendency to formfrom a self-propagating high-temperature synthesis (SHS), such areaction is not likely to be ignited during milling.

For boron carbide–titanium diboride composites formed fromelements, Halverson et al. [40] reported a series of experimentsand thermodynamic calculations, establishing the necessaryrequirements for a SHS reaction. The SHS mode of the reaction forthe composition studied in the present work is possible only whenthe mixture is initially heated up to 927 �C [40]. According to Takacsand McHenry [41], the ball temperature in a SPEX mill does not ex-ceed 100 �C provided that no exothermic reaction takes place duringmilling. Presumably, the local temperatures during HEBM can bemuch higher; however, these local temperatures do not result inan overall heating of the mixture to the requisite SHS temperature.

The goal during SPS consolidation was to create fully reacted,fully dense TiB2–B4C composites. To that end, Anselmi-Tamburiniet al. [38] showed that in mixtures of carbon black and amorphousboron, the formation of boron carbide was nearly complete aftersintering in the SPS at 1200 �C. However, densities higher than90% could only be achieved after sintering at 1700 �C. In that study,a thermocouple was used to monitor the temperature of the SPSprocess, while in the present study a pyrometer was used to mea-sure temperature. Thus, one can assume that the temperatures re-ported in this study are underestimating the actual sinteringtemperature by approximately 100 �C [10]. Given this information,the Ti–B–C mechanically milled powders were sintered in the SPSat 1600 and 1700 �C.

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XRD Patterns of TiB2 - B4C S

Fig. 7. XRD patterns of the boron carbide–titanium diboride composite sintered at 160identical.

XRD patterns of the samples sintered at 1600 and 1700 �C showthe presence of the two phases in the composite (Fig. 7). In thesample sintered at 1700 �C boron carbide is well crystallized. It isworth mentioning that tungsten carbide phase is absent on theXRD patterns of the samples sintered at 1600–1700 �C (Fig. 7)though it is present in the milled powders (Fig. 6). It is known thatat temperatures higher than 1100 �C boron carbide reacts withtungsten carbide to form tungsten boride W2B5 and carbonthrough the following reaction [42]:

B4CþWC!W2B5 þ C

Since the initial stoichiometry of boron carbide was shifted fromB4C to a boron-enriched composition, any residual carbon is not ex-pected to remain in the sintered compacts. The reason tungstenboride phase does not appear on the XRD patterns is due to the for-mation of (Ti, W)B2 solid solutions during sintering. These com-pounds have the crystallographic structure of titanium diborideand have virtually identical lattice parameters [43,44]. This is dueto the small difference in the atomic radii of Ti and W.

The densities of the sample sintered at 1600 �C, with 4 h of mill-ing was 90.6%. The densities of the specimens sintered at 1700 �Cwere all over 95%. The densities for the 1700 �C specimens milledat 4, 8 and 16 h were 95.3%, 96.2% and >99%, respectively. TheSEM images of the cross-sections of the composites are shown inFig. 8. The correspondence of the bright phase on the SEM micro-graphs to titanium diboride was confirmed by the BSE imaging.

As the density data indicates, it was found that increasing themilling time helped improve the final density of the composites.This result is consistent with those of Heian et al. [45]. Throughmore aggressive milling schedules the elemental Ti becomes moredispersed in the elemental C and B. This reduces the size of the TiB2

agglomerates. Large TiB2 agglomerates were found to contain adisproportionate amount of porosity relative to more finely dis-persed TiB2 and B4C phases. Fig. 8 reveals the fracture surface ofthe nearly fully dense material, which consists of titanium diboride

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egrees)

1600 ºC *

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PS Compacts

0 �C (milled for 4 h) and 1700 �C (milled for 16 h). Note the patterns are virtually

Fig. 8. SEM images of the fracture surfaces of the TiB2–B4C composites. The left hand image is of the specimen sintered at 1600 �C with powder milled for 4 h. The adjacentmicrograph shows a nearly fully dense specimen sintered at 1700 �C with powder milled for 16 h.

XRD Pattern of SPS Consolidated HfB2 - SiC Composite

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Fig. 9. XRD Pattern of HfB2–SiC compact after SPS at 1800 �C for 8 min.

Fig. 10. SEM fracture surface of the fully dense HfB2–SiC compact.


grains 1–2 lm in size uniformly distributed among the boron car-bide grains of 5–7 lm. So, prolonged mechanical milling wasadvantageous for the development of a more uniform microstruc-ture. However, the higher sintering temperature resulted in somegrowth of boron carbide grains.

Interestingly, the hardness of the composites produced in thisstudy is lower than the value calculated from the rule of mixtures.According to Gusev [46], the composition studied in this work isclose to the pseudobinary eutectic B4.5C–TiB2. The reduced hard-ness of the eutectic compositions of ceramic compounds was ob-served earlier by other researchers [47]. However, such eutecticswere found to exhibit elevated resistance to brittle fracture. In-deed, from indentation measurements, the sample sintered to fulldensity in this study showed a relatively remarkable fracturetoughness of 5.9 MPa m1/2 as compared to monolithic boron car-bide, which has a fracture toughness ranging from 2.9 to3.7 MPa m1/2 [2]. The 5.9 MPa m1/2 value obtained here is in goodagreement with the results obtained by Skorokhod et al. [28] forhot-pressed B4C–TiB2.

3.3. Hafnium diboride–silicon carbide

XRD analysis revealed that the powder mixture after high en-ergy ball milling remains as before the milling process, i.e. the for-

mation of HfB2 did not occur. Fig. 9 shows the XRD pattern of thecomposite after SPS at 1800 �C for 8 min. It is clear that reactionformation of the HfB2 happened during sintering, which is in


agreement with [10]. The presence of Hf2O3 is believed to be due tothe high energy ball milling that introduced oxygen into the pow-der system. Both transgranular and intergranular facture modeswere observed for the composite, as shown in Fig. 10. The sinteredcomposite is fully dense. The composite has a hardness of 20.5 ±1.5 GPa. High temperature mechanical properties of the compositeare being investigated.

4. Conclusions

The utility of the SPS processes to synthesize and consolidateultra hard & ultra high temperature materials and compositeshas been shown. By combining SPS with HEBM type processes itis possible to create hard, refractory ceramic materials with un-ique, highly tailorable microstructures. While not nanocrystalline,all of the ceramic materials in this study exhibited relatively fine,equiaxed grains. More system specific conclusions are given below.


Functionally graded boron carbide with precipitous propertyand microstructural gradients has been synthesized using SPS. Thisprocessing route results in a material with very promising proper-ties and interesting microstructural features. During SPS the amor-phous powders react and partially consolidate forming a densitygradient. X-ray diffraction data reveals that after SPS the materialis in a stressed state, most likely due to residual stresses fromthe rapid heating and quenching (on the order of 500 �C/min).The defect concentrations appear to decrease after infiltration per-haps simply due to annealing effects or perhaps due to the forma-tion of small amounts of Al3BC, which is also apparent in the XRDpattern.

4.2. Titanium diboride–boron carbide composites

The present study showed that the combination of HEBM andSPS is a promising route for preparation of fine-grained in-situ com-posites in the boron carbide–titanium diboride system. Neither B4Cnor TiB2 was formed during milling. Additionally, no phases otherthan B4C or TiB2 were detected in the samples sintered at 1600–1700 �C. The density of the sintered composites increased with sin-tering temperature. Density also improved with increased millingtime. Densities higher than 95% were obtained at SPS temperaturesas high as 1700 �C. The nearly fully dense material obtained in thisstudy consisted of equiaxed titanium diboride grains 1–2 lm in sizeuniformly distributed among equiaxed boron carbide grains of 5–7 lm. The sintered fully dense composite showed a somewhatremarkable fracture toughness of 5.9 MPa m1/2.

4.3. Hafnium diboride–silicon carbide composites

Fully dense composite was successfully sintered at 1800 �C for8 min by using SPS method. Both transgranular and intergranularfacture modes were observed for the composite. The compositehas a hardness of 20.5 ± 1.5 GPa.

Acknowledgements

This work is supported by the Office of Naval Research (GrantsN00014-03-1-0148 and N00014-07-1-0745) with Dr. LarryKabacoff as program manager and the Army Research Office (Grant# W911NF-04-1-0348) with Dr. Sheldon Cytron as programmanager. The authors are also grateful to Drs. Umberto Anselmi-Tamburini and Cosan Unuvar for experimental assistance andhelpful discussion.

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