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Scripta Materialia 61 (2009) 1075–1078

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Room-temperature dislocation activity during mechanicaldeformation of polycrystalline ultra-high-temperature ceramics

Dipankar Ghosh,a Ghatu Subhasha,* and Gerald R. Bourneb

aDepartment of Mechanical and Aerospace Engineering, University of Florida, Gainesville, FL 32611, USAbDepartment of Materials Science and Engineering, Major Analytical Instrumentation Center, University of Florida, Gainesville,

FL 32611, USA

Received 30 June 2009; accepted 28 August 2009Available online 2 September 2009

Ceramics such as ZrB2 and HfB2 are potential candidates for ultra-high-temperature applications. Their electrical conductivity val-ues are comparable to those of metals. Such unusual electrical properties arise from the presence of metallic bonds in their crystal struc-ture. We argue that the metallicity in chemical bonding is also reflected in their room-temperature mechanical deformation, which wasinvestigated through indentation-induced slip bands and the resulting dislocation activity. These observations were rationalized on thebasis of metallic character of their dislocation core structure.� 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Ceramics; Mechanical deformation; Transmission electron microscopy; Dislocations; Bonding

In recent years, diborides of transition metals,such as zirconium (Zr) and hafnium (Hf), and their com-posites have emerged as potential candidates for high-temperature aerospace applications because of theirextremely high melting temperature (>3000 �C), high-temperature strength retention, good oxidation resistanceat elevated temperatures, etc. [1]. Understanding themechanical response in this class of ultra-high-tempera-ture ceramics (UHTCs) is essential for the design of struc-tural components made out of these materials. Recentindentation [2,3] and scratch [2,4] studies by the authorson a polycrystalline zirconium diboride–silicon carbide(ZrB2–5 wt.% SiC) composite have revealed readilydetectable metal-like plastic deformation features calledslip-lines. Figure 1(a) and (b) are optical micrographs ofroom-temperature static indentation imprints on poly-crystalline ZrB2 and hafnium diboride (HfB2) ceramics,respectively, where intense slip-line or surface-step forma-tion in the vicinity of the indented regions are seen. Suchmacroscopic slip-steps are clear evidence of dislocation-induced plastic deformation and dislocation mobility atroom temperature in these ceramics. Owing to the extre-mely high melting temperature of these UHTCs, suchroom-temperature dislocation mobility is unexpected[5]. However, dislocation studies using transmission elec-tron microscopy (TEM) have been extremely limited in

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* Corresponding author. Tel.: +1 352 392 7005; e-mail: subhash@ufl.edu

ceramics such as ZrB2 and HfB2 [6], particularly in poly-crystalline UHTCs. In this work, which also employsTEM, we have investigated the room-temperature dislo-cation activity responsible for the observed randomly ori-ented slip-lines in ZrB2 ceramics and identified the slipsystems activated during the mechanical deformation.In addition, we have attempted to build an understandingon the origin of room-temperature dislocation activityand mobility in ZrB2 ceramics from crystal structureand chemical bonding perspectives.

TEM specimens were extracted from scratch-inducedplastically deformed grooves in a polycrystalline ZrB2–5wt.% SiC composite due to availability of large regionscontaining slip-lines [4]. The scratch grooves were cre-ated using a Berkovich nanoindenter at room tempera-ture. A site-specific cross-sectional focused ion beam(FIB) technique (FEI Strata DB235) was utilized to ex-tract thin TEM specimens from scratch grooves. Speci-mens were then analyzed in a JEOL 200CX microscope.

The TEM investigations revealed high density of dis-locations within the deformed ZrB2 grains in the regionsdirectly beneath the scratch groove. Figure 2(a) shows abright-field TEM image at the [0 0 0 1] zone axis, reveal-ing a large area of dense dislocation activity andclustering within the deformed ZrB2 grains. Dislocationbands in three different orientations were clearly visible,suggesting slip on multiple planes. Similar multiple andintersecting slip bands were also observed at a macro-

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Figure 3. (a) A bright-field TEM micrograph at ½�12�10� zone axis and(b) the corresponding selected area electron diffraction pattern.

Figure 1. Optical micrographs revealing slip-lines formed in thevicinity of indented regions of (a) ZrB2 and (b) HfB2 ceramics.

1076 D. Ghosh et al. / Scripta Materialia 61 (2009) 1075–1078

scopic scale on the deformed surfaces (see Fig. 2(b)).These slip bands, observed from the TEM and scanningelectron microscopy (SEM) images, appear to be at anorientation of approximately 60o to one another. Suchintersecting slip bands and their threefold symmetryindicated prismatic or pyramidal slip in ZrB2 duringmechanical deformation. To confirm the nature of slip,these three sets of dislocations were imaged in two-beamcondition with three prismatic planes, ð1 0�10Þ, ð01�10Þ,and ð�1100Þ. For each of these planes, one of the threesets of dislocation bands was out of contrast, suggestingprismatic slip within ZrB2 grains. The h1 1�20i typedirections are the closest-pack directions containedwithin the f10�10g type prismatic planes and are ori-ented 90o from the poles of these planes. For example,the pole of the ð1�100Þ plane is perpendicular to the½11�20� direction. Therefore, it was inferred that Burgersvectors of these dislocations are parallel to the h11�20idirections. The possibility of [0 0 0 1] being the Burgersvector for these sets of dislocations (as c/a ratio in ZrB2

is 1.114) was excluded as they were clearly visible at the

Figure 2. (a) A bright-field TEM micrograph revealing dense disloca-tion activity on prismatic planes and (b) multiple sets of intersectingslip-lines formed on the surface due to scratch load, resemblingpatterns observed in (a).

[0 0 0 1] zone axis (see Fig. 2(a)). Thus, we inferf10�10gh11�20i slip in this ceramic.

The TEM investigations also revealed another slipsystem less commonly reported in hcp materials [21].Figure 3(a) and (b) show a bright-field image of disloca-tions at the ½�12�10� zone axis and the corresponding se-lected area electron diffraction pattern, respectively.Since these dislocations were visible at the ½�12�10� zoneaxis, their Burgers vector cannot be parallel to this ori-entation, for they would be invisible. This zone axis con-tains (0 0 0 1), ð10�10Þ and ð10�11Þ reciprocal latticevectors (g), corresponding to basal, prismatic and pyra-midal planes, respectively. Contrast experiments re-vealed that these dislocations (as seen in Fig. 3(a))remained in contrast when imaged with ð0001Þ andð10�11Þ reflecting planes but were completely invisiblewith ð10�10Þ plane. This excluded the possibility off0001gh11�20i slip (basal slip) in ZrB2. Since the onlyother low-index direction perpendicular to the pole ofð10�10Þ plane is ½0001�, the above contrast experimentssuggested that the Burgers vector is parallel to ½0001�.For further confirmation, the same set of dislocationswas imaged with a ð11�20Þ plane (in ½�1100� zone) be-cause the f1 1�20g type of planes (secondary prismatictype) only contain the ½000 1� Burgers vector. As thesedislocations went completely out of contrast, it was con-cluded that Burgers vectors are parallel to the ½0001�direction. This Burgers vector has never been reportedin the open literature for ZrB2 ceramics. Although thespecific slip-plane was not identified, for dislocationswith Burgers vector parallel to the ½0001� orientation,the only possible close-pack slip-plane is of prismatictype, either f10�10g or f1 1�20g. The f1 0�10g primaryprismatic planes have a higher planer density (0.727)compared to that of the f11�20g secondary prismaticplanes (0.682). Thus, f10�10g½0001� is the most proba-ble slip system compared to f11�20g½0001� in ZrB2.However, due to nearly comparable packing density ofthe primary and secondary prismatic planes, slip mayalso occur in the f11�2 0g system. Nevertheless, the cur-rent TEM investigations clearly revealed prismatic slipas the only identified room-temperature plastic deforma-tion mechanism in polycrystalline ZrB2. The above

D. Ghosh et al. / Scripta Materialia 61 (2009) 1075–1078 1077

TEM studies are also consistent with the work by Hagg-erty and Lee in ZrB2 single crystals [6].

There is also considerable evidence [7,8] supportingprismatic slip being favored over basal slip in ZrB2.Xuan et al. [7] performed Vickers microhardness mea-surements in ZrB2 single crystals, on basal ð0001Þ,and prismatic ð10�10Þ and ð11�20Þ, planes, from roomtemperature to 1000 �C. It was observed that over theentire temperature range hardness was same on bothtypes of prismatic planes but was lower than on the ba-sal plane. A similar result was also obtained in the workof Nakano et al. [8]. When a basal plane is indented,only possible slip-planes are pyramidal planes. On theother hand, indentation on the ð10�10Þ and ð11�20Þplanes can cause slip on other prismatic planes fromthe f10�1 0g and f11�20g families as well as on pyrami-dal planes. Therefore, it is clear that in ZrB2 prismaticslip is more favorable. The work of Xuan et al. [7] andNakano et al. [8] also suggested that dislocation motionis probably equally favorable on both the f10�10g andf11�20g prismatic planes. The current TEM investiga-tion clearly revealed that prismatic slip was the onlyidentified mechanism responsible for room-temperatureplastic deformation in polycrystalline ZrB2.

Having determined the nature of dislocations, in thefollowing we have attempted to rationalize room-temper-ature dislocation activity and mobility based on the prop-erties, crystal structure and chemical bonding present inthese ceramics. The nature of chemical bonding plays adominant role in governing the properties of metals andceramics. In general, the non-metallic character of bond-ing (mainly covalent) in structural ceramics is reflectedwell in their properties through high elastic modulus,chemical stability at room- and high-temperatures, andextremely low electrical conductivity, and thus corre-sponding high electrical resistivity in comparison tometals [9]. In contrast, UHTCs have a unique combina-tion of properties that are characteristics of both metalsand ceramics [1]. As mentioned before, their high melting

(a)

(b)

a

M (Zr, Hf)

B

0001[ ]1123[ ]

1120[ ]

c

M-M bonding

B-B bonding

M-M bonding

M-B bonding

M-B bonding

1.00E-15

1.00E-12

1.00E-09

1.00E-06

1.00E-03

1.00E+00

1.00E+03

1.00E+06

1.00E+09

0 100 200 300 400 500 600

Elastic modulus (GPa)

Ele

ctri

cal c

ondu

ctiv

ity

(S/m

)

1.00E-02

1.00E+01

1.00E+04

1.00E+07

1.00E+10

1.00E+13

1.00E+16

1.00E+19

1.00E+22

1.00E+25

Ele

ctri

cal r

esis

tivi

ty ( µ

Ohm

-cm

)HfB2

TiB2

ZrB2UHTCs

ZrO2

Si3N4 Al2O3

Structural ceramicsSiC

Au Zr Ti Hf

Metals

Cu

Ag Al

Figure 4. (a) Electrical conductivity (filled symbols) and electricalresistivity (open symbols) of selected metals and ceramics and (b)schematic of ZrB2 and HfB2 crystal structure.

temperature (>3000 �C), high elastic modulus(>400 GPa) and superior thermo-chemical stability aretypical of ceramics. In addition, UHTCs have surpris-ingly high electrical conductivity (>106 S m�1) and thuslow electrical resistivity (<10�5 X cm), comparable tothose of metals. These values are in contrast to typicalstructural ceramics such as aluminum oxide (Al2O3),silicon carbide (SiC) and silicon nitride (Si3N4), as illus-trated in Figure 4(a) [1,10–18]. The high electrical con-ductivity indicates significant metallicity in the chemicalbonding of UHTCs.

ZrB2 and HfB2 have AlB2-type (P6/mmm spacegroup) hexagonal crystal structures [1] where the unitcell contains a six-member graphite-like ring or a netof boron (B) atoms in two dimensions which alternateswith close-packed layers of metal (M) atoms (seeFig. 4(b)). Vajeeston et al. [19] studied the electronicstructure and nature of chemical bonding in AlB2-typeceramics using the self-consistent tight-binding linearmuffin-tin orbital method. Zhang et al. [20] utilized thefirst-principles total-energy plane-wave pseudopotentialmethod and investigated the electronic structures andelastic properties of ZrB2 and HfB2 ceramics. Thesestudies revealed the existence of three types of chemicalbonds in this crystal structure: (i) B–B, (ii) M–B, and (iii)M–M (M = Zr, Hf) [19,20]. While the B–B bonds arepurely covalent, the M–B bonds are a mixture of cova-lent and ionic characters. However, due to small differ-ences in electronegativity values (1.33, 1.30, and 2.04for Zr, Hf, and B, respectively), ionicity is less than8%. In contrast, the M–M bonds are predominantlymetallic due to nearly free-electron non-bonding d-orbi-tal states which contribute to significant metallicity inthese ceramics [19,20]. Both covalent and metallic bondsare known to exhibit strong influence on properties ofmaterials. Here we argue that such metallic characterof chemical bonding in ZrB2 and HfB2 ceramics is wellreflected in their metal-like mechanical deformationbehavior, as observed in indentation and scratch studiesat room temperature [2–4].

For a given crystal structure, dislocation mobility isstrongly influenced by chemical bonding, especially atroom temperature, because chemical bonds constantlydistort, break and reform during the dislocation motion[21,22]. Gilman [22,23] correlated the optical band gapwith the glide activation energy in covalently bondedsolids and concluded that it is the local or non-localnature of chemical bonding that determines mobilityat room temperature. In a similar fashion, dislocationmobility can be qualitatively correlated with the localor non-local character of chemical bonding in ZrB2

and in other AlB2-type UHTCs as having a mixture ofcovalent and metallic bonds [19,20]. Similar to theoptical band gap, the electrical conductivity can alsoreflect the local or non-local nature in chemical bonding.While the exact mechanism of high electrical conductiv-ity in ZrB2 and HfB2 ceramics is not known, Zhang et al.[20] suggested that their unusually high conductivenature is related to the Zr–Zr and Hf–Hf metallicbonding. We argue that such bonding characteristicscan also play a significant role in the room-temperatureplastic deformation and the associated dislocationmobility in these ceramics. One of the reasons why such

1078 D. Ghosh et al. / Scripta Materialia 61 (2009) 1075–1078

room-temperature macroscopic plastic deformationfeatures are not readily observed in other structuralceramics (e.g. Al2O3, ZrO2, SiC, and Si3N4) is that theypossess strong covalent and ionic bonding instead ofdirect metal–metal bonds [24–27]. Such bonding makesdislocation mobility extremely difficult, particularly atroom temperature, where availability of thermal energyis low. Dislocation mobility is strongly influenced by thewidth of the dislocation core, which depends on the typeof chemical bonding. While metallic bonding results in awide dislocation core, covalent bonding makes it narrow[24]. The energy barrier for dislocation mobility of awide dislocation core, consisting of non-directionalmetallic bonds, will be lower compared to that of anarrow core structure made up of highly directionalcovalent bonds [24]. A classic example is cubic-SiC,which has the required number of independent slip sys-tems to exhibit room-temperature plasticity but, due toits pure covalent bonding, has too narrow a dislocationcore and thus dislocations are immobile at room temper-ature [24].

An accurate description of dislocation core structureand its motion in ZrB2 (or HfB2) crystal structure is com-plex, and requires a fully atomistic approach. However,for a prismatic slip in ZrB2, the dislocation core structurenot only involves B–B and Zr–B bonds but also Zr–Zrbonds. The existence of metallic bonds can make the dis-location core relatively wider compared to a core made upof only B–B or Zr–B bonds. Thus, the Zr–Zr bonds are as-sumed to play a crucial role in enhancing the dislocationmobility at room temperature by lowering the chemicalbarrier. Due to this non-localized nature of Zr–Zr bondsand the associated increased dislocation mobility duringmechanical deformation, dislocations are not restrictedto a small material volume (typical of ceramics) but,rather, extend over a relatively larger region, allowingthe dislocation activity to be readily visible even at roomtemperature at distances further away from the regionof load application, as seen in Figure 1.

In summary, ZrB2 and HfB2 ceramics were observedto exhibit metal-like slip-line patterns at room tempera-ture when deformed by indentations and scratch loads.This deformation pattern indicates significant room-temperature dislocation mobility in these ceramics. In-depth TEM analysis in a ZrB2–SiC composite revealedthat dislocations were indeed generated within theZrB2 phase as a result of mechanical deformation atroom temperature. Prismatic slip was identified as themajor plastic deformation mechanism in polycrystallineZrB2 ceramic. Finally, a possible role of Zr–Zr metallicbonds in room-temperature dislocation mobility wasdiscussed.

The authors thank the Major Analytical Instru-mentation Center at the University of Florida for provid-ing access to SEM, TEM, and FIB.

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