Deformation Microstructure in (001) Single Crystal Strontium Titanate by Vickers Indentation

Deformation Microstructure in (001) Single Crystal Strontium Titanateby Vickers Indentation

Kai-Hsun Yang, New-Jin Ho, and Hong-Yang Lu*,w

Department of Materials Science and Centre for Nanoscience, National Sun Yat-Sen University,Kaohsiung 80424, Taiwan

Recent interests on the plastic deformation of strontium titanate(SrTiO3) are derived from its unusual ductile-to-brittle-to-duc-tile transition (DBDT). The transition is divided into threeregimes (A, B, and C) corresponding to the temperature rangeof 1131053 K (1601 to 7801C), 1053 to B1503 K (7801 toB12301C), and B15031873 K (B12301 to 16001C), discov-ered by Sigle and colleagues in the MPI-Stuttgart. We reportthe dislocation substructures in (001) single crystal SrTiO3deformed by Vickers indentation at room temperature, studiedby scanning and transmission electron microscopy. Dislocationdipoles of screw and edge character are observed and conrmedby insideoutside contrast using7g-vector by weak-beam darkeld imaging. They are formed by edge trapping, jog dragging,and cross slip pinching-off. Similar to dipole breaking off in de-formed sapphire (a-Al2O3) at 12001C and g-TiAl intermetallicat room temperature, the dipoles pinch off at one end, and emit astring of loops at trail. Two sets of slip systems {110}/110Sand {100}/011S are activated under both 100 g and 1 kg load.The suggestion is that plastic deformation has reached the stageII work hardening, which is characterized by multiplication ofdislocations through cross slip, interactions between disloca-tions, and operating of multiple slip systems.

I. Introduction

MOST ceramics are brittle and fractured by cleavage withoutappreciable plastic deformation (of a failure straineo0.1%) at room temperature, but with increasing tempera-ture they become more plastic at and above a transition tem-perature (TC). The transition temperature varies, e.g.,41273 K(10001C) for sapphire (single-crystal alumina (a-Al2O3)).

1 Suchmechanical behavior is commonly known as the brittle-to-ductile transition (BDT).14

Strontium titanate (SrTiO3) has the perovskite structure, andretains the cubic symmetry until below 105.5 K (167.51C).5 Notonly does it undergo a signicant plastic deformation by a com-pressive strain e5 7%9% at room temperature, but also exhib-its an unusual ductile-to-brittle-to-ductile transition (DBDT).68

Taking SrTiO3 as a structural analogue,9 the rheology of Earths

lower mantle,10 consisting predominantly of silicate perovskites,(Mg,Fe)SiO3, can be simulated and investigated.

Two types of BDT are commonly reported.11,12 Gradualtransition from brittle to ductile is more frequently observed,where the fracture stress increases over a temperature range of4373 K (1001C), e.g., in steels, BCC metals, intermetallics, and

Al2O3. On the other hand, transitions in Si and sapphire are verysharp and occur abruptly within 51101C. The transition is afunction of the strain rate (or the rate of the applied stress in-tensity factor) and structure sensitive dependent on the variabil-ity in the density of dislocation sources, and is temperaturedependent.11,13 Sharp transition is due to triggering of disloca-tion sources while gradual transition has dislocation sourcesreadily available at the crack tip. Not only does SrTiO3 exhibitunusual ductility at room temperature,68 its BDT to high-temperature ductile behavior over an extended temperaturerange of B1050 K (7771C) and B1500 K (12271C) also sug-gests the transition is the gradual type.3,4

Deformed microstructures in SrTiO3 have been reported.68,14

They were often taken to support the proposed DBDT models,e.g. dissociated collinear partials of smaller cores and highercore energy rendering larger Peierls stresses were taken to suggesttemperature-dependent core structure, and decreasing length ofmobile dislocation segments. However, the room temperatureplasticity for a traditionally brittle ceramic is yet satisfactorilyexplained.

The Vickers indentation technique is used to investigatehow dislocations are generated and migrated in SrTiO3 in theplastic regime (i.e., regime A in Brunner et al.14). We reportan analysis of dislocation substructure thus produced in (001)single crystal to understand how the plasticity occurs at roomtemperature.

II. Experimental Procedures

Single crystal SrTiO3 containing 0.7 wt% Nb-doping (approx-imately 0.02 mol/% Nb2O5), of the /001S orientation suppliedby MTI (Richmond, CA) were used in this study. The crystalswere microindented by two applied loads of 100 g and 1 kg usinga Vickers indenter (Nikon AVK-C2, Tokyo, Japan). Extrinsicdefects, e.g. charge-compensating cation vacancies of VSr

00 andVTi00 00 (using the Kroger-Vink notation) for Nb2O5-donor oxide,

may have contributed to plastic deformation by dislocation pin-ning, and kink migration. The indented samples were thenchemical-etched for 60 s in a solution of 70% HNO3 addedwith a few drops of HF. Microstructure was observed using thescanning electron microscope (SEM, SEM6330, JEOLt, To-kyo, Japan) equipped with a cathodoluminescence (CL) detector(MonoCL2, Gatant, Pleasanton, CA) operating at 20 kV, andthe transmission electron microscope (TEM, AEM3010,JEOLt) operating at 300 kV.

TEM foils were prepared either by the conventional tech-nique of cutting, polishing, dimple-grinding mechanically, andAr-ion beam thinning (PIPSt, Gatan) to electron transparency,or using a focus ion beam (FIB) apparatus (SMI3050, Seiko,Tokyo, Japan).

III. Results

A representative indentation crack is shown in Fig. 1 (for load-ing with 1 kg). A large c/d-ratio42.5 (where d is the diagonal of

T. Mitchellcontributing editor

*Member, The American Ceramic Society.This research is funded by the National Science Council of Taiwan through contract

NSC-95-2221-E110-033, and the Ministry of Education through the Centre of NanoScienceand Technology.

wAuthor to whom correspondence should be addressed. e-mail: [email protected]

Manuscript No. 25998. Received March 11, 2009; approved May 2, 2009.

Journal

J. Am. Ceram. Soc., 92 [10] 23452353 (2009)

DOI: 10.1111/j.1551-2916.2009.03189.x

r 2009 The American Ceramic Society

2345

the Vickers impression and c the crack length measured from thecenter of the impression) suggests15 the crack is the half-pennytype. Fracture toughness KIC0.89 MPa m1/2, similar to thosereported by Bernard et al.,15 was determined using the formula

KIC constantE=H1=2P=c3=2

where constant is 0.0154 (dimensionless), E/H is the ratio ofelastic modulus (225 GPa) to hardness (9.5 GPa), and P theapplied load (9.8 N).

(1) Slip Lines and Dislocation Etch Pits

Figure 2(a) shows radial cracks emanating from the indent im-pression (as indicated). Dislocation etch pits are aligned alongboth the /010S and /110S directions, making a 451 angle toeach other.15 The indication is that two slip planes intersect at451 as depicted in the inset of Fig. 2(a). A schematic illustrationof etch pits constituting slip lines is shown in juxtaposition,where the radial cracks lie parallel to /110S. Slip lines indicatethat the dislocations are generated predominantly around theindent. Dislocations are viewed end-on, where they lie in slipplane and constitute a slip band.

Dislocation multiplication by multiple cross glide is evidencedfrom the extending and lateral growth of slip bands (Fig. 2(a))along a major radial crack. A dislocation-free zone16,17 of sev-eral micrometers wide on both sides, straddled across the crack,where no etch pits are visible (indicated in Fig. 2(a)), is detectedalong the whole length of radial crack.

Strain eld and slip lines are clearly visible in CL images. Theareas of bright contrast indicate highly strained region producedby indentation (Fig. 3(a)), and the criss-cross dark lines extend-ing along the /110S and /100S directions are the slip lines.These slip lines running along two directions intersect witheach other. While the strain contrast is revealed unambiguouslyby CL, the slip lines are also signicantly better discerned(Fig. 3(b)) as compared with SEM secondary-electron image(SEI) in Fig. 2(a). This also conrms the observation shown inFigs. 2(a) and (b), where dislocation-etch pits are analyzed bySEM-SEI.

(2) Burgers Vectors and True Line Directions

Representative dislocations in the crack vicinity are shownin Fig. 4(a) by (g,3g) weak-beam dark eld (WBDF) imag-ing. All TEM images shown here are taken from along the [001]zone axis.

Slip planes are determined by b u5 slip plane normalfrom the dislocation Burgers vector (b) and the true disloca-tion line direction (u). Both b and u are determined by conven-tional contrast analysis; the former by adopting the invisibilitycriteria of g b5 0 or 2np from more than three noncoplanar

g-vectors, the latter by performing trace analysis for at leastthree projected dislocation images from the correspondingstereographic projection (using CaRIne

s

Crystallography,version 3.1).

Tables I and II summarize an analysis of dislocations shownin Figs. 4(a) and (b) for Burgers vectors and true line directions,and the type of dislocations determined from angles between band u. Both the edge dislocation B1 (with bB5 [110]) in Fig. 4(a)and the edge dislocation E1 (with bA5 [011]) in Fig. 4(b),are produced by indentation. Accordingly, both the primary{110}/110S6 and secondary {001}/110S slip systems are acti-vated by indentation with a 100 g load at room temperature.The angle between the faces of the pyramidal Vickers indenterwith a square base is 1361. Although the Schmid factor vanishesin pure compression along /001S, shear stresses generatedalong the interface between the indenter and work piece(i.e., /001S SrTiO3 single crystal) would have made the Sch-mid factor nonzero (m5 0.172) such that the slip system {001}/110S is activated.

Edge dislocation glides in its slip plane, as indicated, isevidenced from the upper left corner of Fig. 4(b). The intersec-tion of screw and edge dislocations would have resulted informing jogs in both screw dislocation D1 and edge dislocationE1

18 after they have moved passing each other. Figure 4(b)reveals the mutual interaction occurring at the intersection, asindicated in the inset. Sessile jogs can only move nonconser-vatively by climb.

(3) Dislocation Intersection and Screw Dipoles

Dislocations in the framed region of Fig. 5(a) are furtheranalyzed. Intersecting screw (A2) and edge (B2) dislocations,similar to those in Fig. 4(b) are again detected. Adopting theinsideoutside contrast technique, the features (as indicated byarrows) are determined to be dislocation dipoles,6 rather thandissociated partials,19 by imaging with 7g-vectors understrong-beam diffraction conditions. Results of (g,3g) WBDFimages are shown in Figs. 5(b) and (c) under g5 110(exhib-iting outside contrast and wider separation between dipoles) and1g5 110 (exhibiting inside contrast and narrower separationbetween dipoles), respectively. Although collinear partialsin sintered SrTiO3 ceramics

20 and in undoped bicrystals21 werereported, only dipoles6 are detected here in the indented /001Ssingle crystal.

Both edge (D1) and screw (D2 and D3) dipoles, and mixedtype (D4 and D5), as indicated in Fig. 5(b), are found.Screw dipole D2 pinching-off at its end has cross slippedand containing segments of a mixed nature. Most of the dipolesare close ended in the view, one (D1) of the open-ended dipolesis indicated in Fig. 5(c). Spherical loops pinched off fromthe close end of a screw dipole are detected (indicated inFig. 5(a)).

Under the applied stress, an edge dislocation gliding in itsslip plane has drawn out two screw dislocation lines froman edge jog, J1, as indicated in Fig. 5(d), and forms a screwdipole (D6, also indicated). The two dislocation arms (desig-nated D6s) joining at edge jog (J1) and forming a dipole, havethe same Burgers vector but mutually antiparallel line direc-tions. An edge dipole bowing out at another sessile jog (J2)generating an edge dipole trail is also observed, which is rela-tively short in length (e.g., comparing with edge dipole D1 andscrew dipole D3 in Fig. 5(b)). Edge dipole trails are ubiquitous inother areas of the deformed crystal, as indicated in Fig. 5(e).Dislocation entanglement forming subgrain boundaries isalso evidenced.

(4) Dislocation Intersection and Edge Dipoles

Dislocations with Burgers vector bB5/110S of a mixed char-acter, indicated by lled arrowheads (Fig. 6(a)), are found tocross slip in (110). Only the primary slip system {110}/110S isactivated, representing the initial stage of plastic deformation.

Fig. 1. A typical indentation crack of the half-penny shape having c/d-ratio 42.5 (SEM-SEI).

2346 Journal of the American Ceramic SocietyYang et al. Vol. 92, No. 10

The framed area imaged by WBDF is shown in Figs. 6(b) and(c), using 7g-vectors. Two edge dipoles (indicated by D7 andD8) appear in the eld of view. At one end of dipole D7 is itsbowed-out segments, of a screw character (D7s, as indicated inFig. 6(a)). A portion of dipole D7 shows a distinctive contrastfrom the rest of the dipole (Fig. 6(b)). The indication is that ei-ther or both dipole arms, of an edge character, have slipped intheir own glide plane. At the other end of D7 is a jog (designatedas J3). Nevertheless, the screw segment D8s (as indicated in

Fig. 6(c)) of edge dipole D8, indicated by lled arrowheads,remain separated under the applied stress. Again, subgrainboundaries are formed (Fig. 6(d)).

Dislocation dipoles are further decomposed into a stringof loops by pinching-off at dipole trail. Although appearedin both samples, this is more pronounced in crystals indentedby the higher load of 1 kg. Loops are broken up from one endof both the open-ended (Fig. 7(a)) and close-ended (Fig. 7(b))dipoles.

Fig. 2. Dislocation etch pits (b) in the vicinity of radial cracks (a), running at 451 to each other (SEM-SEI).

October 2009 Deformation of SrTiO3 by Vickers Indentation 2347

IV. Discussion

Dislocation substructure in Vickers-indented (001) single crystalSrTiO3 is represented by: (1) dislocations of edge, screw, andmixed character, (2) sessile jogs, (3) dipole trails, edge, screw,and mixed dipoles, (4) decomposition of such dipoles into loops,and (5) formation of subgrain boundaries. Together with twoslip systems {110}/110S and {100}/011S activated by the in-dentation stress, this suggests that plastic deformation in suchsingle crystals has the characteristic features of the stage II workhardening (Figs. 4(a), 5(a) and 6(b), and Table II). Deformationoccurs when the easy glide system /110S{110} is operative,as experimentally evidenced (Table II). It was activated atB900 mN for the (001) single crystals,22 but the stage I workhardening would have been obscured at the indentation stressesbecause the primary and secondary slip systems are both acti-vated. Therefore, a signicant plastic deformation would havetaken place at room temperature before unstable crack growthensues to result in brittle fracture eventually.

(1) Mechanisms of Dipole Formation

For ductile materials, represented by FCC metals, dislocationdipoles are a feature of the stage I hardening when dislocationglide only occurs in a single slip system. They are formed byedge trapping in this stage. Edge dislocations with Burgersvectors of the opposite sign,18,23,24 gliding passed each otheron parallel slip planes are trapped by their mutual (attractive)stress eld and form edge dipoles.23,24 Edge dipoles are alsoproduced by jog dragging when a gliding screw dislocation,pinned by a sessile edge jog of small length (y),18 draws outtwo edge-dislocation lines of the same Burgers vector whenapplied stresses (per unit length, tb) are unable to overcome

their mutual repulsion force (B0.25 Gb2/2p(1n)y, whereG: shear modulus, n: Poisson ratio). Jog dragging only occursafter sessile jogs have been produced from the intersectionbetween dislocations.

Edge dipoles may have been generated in deformed SrTiO3single crystals by three mechanisms. (i) Firstly, similar tosapphire,23,24 edge dipoles in deformed SrTiO3 (represented byD1 in Fig. 5(b)) are most likely produced from mutual trappinginitially in stage I hardening. Therefore, such edge dipoles ofopposite sign are open ended (e.g., D1).

Fig. 3. (a) Strain eld and (b) slip lines produced by indentation as re-vealed by SEM-CL image.

Fig. 4. Dislocations in the crack vicinity shown by (g,3g) weak-beamdark eld (WBDF) imaging whose Burgers vectors (b) and true line di-rections (u) are determined [(a) in Table I and (b) in Table II], jogs (asframed) are formed at the intersection between screw (D1) and edge (E1)dislocations (TEM).

Table I. Determination of Burgers Vectors (b) Adoptingg b5 0 Invisibility Criteria from WBDF Images at Three

Noncoplanar g-Vectors

Z

g b

b

001 113 102 112112

g 200 110 020 110 110 211 121 211 020 111 110

A1 X O O O O X O O O X O 7[011]B1 X O O O O X O O O X O 7[011]C1 O O X O O O X O X X O 7[101]

Dislocations are invisible when g b5 0. X, invisible; O, visible.

2348 Journal of the American Ceramic SocietyYang et al. Vol. 92, No. 10

(ii) Secondly, edge dipoles are formed by a mechanismdescribable by cross slip and pinching-off proposed by Tetel-man25 for deformed single crystal Fe-3 wt% Si. Suchmechanismfacilitated by a long jog produced by cross slip is best illus-trated schematically in Figs. 8(a)(c). Double cross slip providessources for long jogs, which are the edge component laidin the cross-slip plane. Two dislocations with Burgers vector bof the same sign but inclined to each other are gliding ontwo parallel planes (Fig. 8(a)). Part of their lengths reori-ented by an applied stress, have resulted in two parallel edgesegments of opposite line directions (Fig. 8(b)). For small joglength (y), only large stresses are able to separate the twodislocations that they stay attracted mutually. If one end ofone of the dislocation segments can cross slip to join the other,this creates a superjog (of a minimum length of the separationbetween two parallel planes), and the joint pinches off readilyresults in an edge dipole (Fig. 8(c)). Such dipoles generatedat jogs from cross slip are usually close ended and longerin length, e.g., D7 from long jog J3 in Fig. 6(c). It is likely that

dipole D8 is in an intermediate stage of cross slip and pinching-off, where the two dislocations are cross gliding at one end(Fig. 8(b)).

(iii) Thirdly, although edge dipoles can also be producedfrom bowing out at an edge jog of a screw dislocation, theyare usually rather short in length and form dipole trails.26,27

Sessile jogs are produced18 when screw and edge dislocations,or edge and edge dislocations intersect each other. Whentwo slip systems (Table II) are activated at room temperature,edge dipoles (Figs. 5(a) and (b)) are formed by jog drag-ging (e.g., at J2 in Fig. 5(b)). Such sessile jogs are formedin edge as well as screw dislocations with mutually perpendic-ular line directions (i.e., screw dislocations D1 and edge dislo-cations E1 shown in Fig. 4(b)). Nevertheless, two screwdislocation arms tagging behind an edge dipole trail are easilyannihilated, as evidenced in single crystal MgO during in situtensile test.27

Screw dislocations usually assumed to have a higher veloc-ity23 are easily annihilated by cross slip. Screw dipoles are there-fore unlikely to be formed by trapping (i.e., mechanism (i)), thisleaves two possible mechanisms. They are either produced bycross slip and pinching-off (mechanism (ii)25), or a direct resultof edge jogs on screw dislocations (mechanism (iii)). Suchdipoles are bowed out at a sessile jog from edge dislocations(e.g., D6 in Fig. 5(d)), which have been intersected by screw dis-locations. Therefore, screw dipoles are usually the close-endedtype terminating at a sessile jog. On the other hand, edge dipolesare both open ended (D1 in Fig. 5(c)) when formed by edgetrapping in stage I and close ended at super jog J3 (D7 inFig. 6(b)) when formed by Tetelmans mechanism25 (Figs.8(a)(c)) in stage II hardening.

Table II. Slips Systems with Slip Plane NormalDetermined from Dislocation Burgers Vectors (b) and

True Line Directions (u)

Dislocation True direction ut Burger vector b Slip plane Dislocation type

A1 uA15 [001] 7[ 011] 100 Mixed (451)B1 uB15 [144] 7[ 011] (011) Mixed (101)C1 uC15 [001] 7[101] 010 Mixed (451)

Fig. 5. (a) Dislocations in the framed region, forming jogs; weak-beam dark eld (WBDF) images using (b) g5 110, (c) 1g5110 shows inside andoutside contrast conrming that they are dislocation dipoles, other features include (d) screw dipole D6 formed by a gliding edge dislocation in its slipplane, and (e) edge dipole trails (TEM).

October 2009 Deformation of SrTiO3 by Vickers Indentation 2349

Fig. 6. (a) Cross-slipped dislocations, (b) and (c) insideoutside contrast shown by weak-beam dark eld (WBDF) images of edge dipoles using7g-vectors, respectively, which are produced by cross slip and pinching-off. (d) Sub-grain boundaries are formed (BF) image (TEM).

2350 Journal of the American Ceramic SocietyYang et al. Vol. 92, No. 10

(2) Formation of Dislocation Loops by Cross Slip

Similar loops were reported for sapphire23,24 deformed by com-pression at high temperatures (@12001C). It was modeled on thebasis of the Rayleigh instability,28,29 describing the morpholog-ical change of solid rods or pore channels. However, conserva-tive climb in sapphire involving pipe diffusion23,24 is unlikely tooccur at room temperature. This may only occur if atom diffu-sion leading to unstable uctuation is signicantly enhanced inSrTiO3 by the applied stress.

Veyssie`re and Gre`gori30 also proposed a mechanism of dipoletruncation by cross-slip annihilation for the prismatic loopsproduced from unjogged dipoles in Al-rich g-TiAl deformed atroom temperature. A schematic illustration based on the orig-inal drawings30 is shown in Figs. 9(a)(d).

At a local position in open-ended dipoles (e.g., D2 and D3in Fig. 5(b)), reorientation occurs to a screw character due tothe applied stress, which could then cross glide. The dipolesare divided into two close-ended segments each terminatedat a sessile jog in the slip plane. When such cross glidehappens in several places along the dipole, an equal numberof loops are produced. It may occur to dipoles of an edge and a

mixed character. Edge dipoles may also acquire kinks to facil-itate cross slip. Kinks are formed in edge dipoles when inter-secting with edge dislocations, or nucleated upon moving acrossPeierls barrier.18

The process relies on cross slip of the screw portion when thedipoles are anchored at both ends (by e.g., extrinsic defects orimpurities). Further, several cross-slip positions (i.e., the numberof loops formed as shown in Figs. 7(a) and (b)) must nucleatesimultaneously along the dipole. This necessitates that the uc-tuation of dipole orientation make the dipole possess corre-spondingly an equal number of positions with a screw characterfor cross slip to ensue. Again, the driving force for dipole re-orientation must come from the applied stress when deformed atroom temperature.

Price31 suggested a mechanism for the prismatic loopsemitted from screw dipoles based on the observation of Zndeformed at room temperature. Several segments along a screwdipole must cross glide simultaneously under the applied stress.The cross-slipped screw segments are annihilated and thatresults in pinching-off to a string of prismatic loops. This pro-cess is shown schematically in Figs. 10(a)(c), which eventuallyleaves an edge dislocation with two large jogs (Fig. 10(d)) but itis not observed here.

All mechanisms proposed before for sapphire, g-TiAl inter-metallic, and Zn for pinching-off loops require stress uctuationin a long-range stress eld produced by indentation along dipolelines a priori for producing such loops, which is likely to comefrom the applied stress.

(3) Plasticity at Room Temperature

The slip families and resulting numbers of independent slip sys-tems from FCC, BCC, and HCP metals32 and SrTiO3 are listedin Table III. It is derived experimentally from dislocation anal-ysis that all together ve systems are activated in SrTiO3:(010)[101], (100)[011], (001)[110], (101)[101], and (110)[110].

The von Mises criterion requires ve independent slipsystems33 operative for signicant plasticity, but this is a neces-sary not sufcient condition.32 The ability of each grain in poly-crystals deformed simultaneously by an arbitrary amountof strain along the ve systems and the slip systems operativecooperatively known as slip exibility,32 is further required forthe ductility in polycrystalline solids to be realized. Althoughonly two or three of such slip systems were usually observed,34,35

signicant plasticity in polycrystalline SrTiO3 ceramics can beexpected.

The present results are consistent with Brunner and col-leagues6 in the sense that plastic deformation at room temper-ature has also occurred in the indented samples. Although theplastic strain is not determined experimentally, and dislocationdissociation affecting its mobility is not detected here, the plas-ticity is demonstrated unambiguously from the microstructurefeatures characteristic of the stage II work hardening (Figs. 4(a),5(a) and 6(b), and Table II). Only that shear stresses introducedby Vickers indentation has also induced the secondary slipfamily {001}/110S. Its contribution to the deformation strain isnot known at the present time because the deformation strainwas not determined. Nevertheless, plasticity would be further

Fig. 7. Dislocation loops decomposed from (a) open-ended and(b) closed-ended dipoles (TEM).

Fig. 8. Schematic illustration for a mechanism of edge dipole formation proposed by Tetelman25: (a) two dislocations of same b gliding in parallelplanes, (b) reorientation of a portion by applied stress leading to two parallel edge segments of opposite u, and (c) Cross slip resulting in joining of twodislocations, pinching-off jog forms an edge dipole.

October 2009 Deformation of SrTiO3 by Vickers Indentation 2351

enhanced with the additional, secondary slip systems induced bythe applied load.

V. Conclusions

Besides {110}/110S reported for compressed samples,6 anadditional, secondary slip system of {001}/110S is identiedin (001) SrTiO3 single crystal deformed at room temperature byVickers indentation. Both screw and edge dipoles are observed,

Fig. 10. Prices model31 of loop formation from (a) a screw dipole, (b) cross slip of several segments in the dipole, (c) and (d) annihilation of cross-slipped segments resulting in pinching-off a string of loops, and leaving the initial edge dislocation with two long jogs.

Fig. 9. Amechanism of dipole truncation by cross-slip annihilation producing prismatic loops from two unjogged dipoles based on Veyssie`re and Gre`gori30:(a) dipole suffers local reorientation along screw components, (b) annihilation of screw portions in the cross-slip plane resulting in two closed-ended dipoles,(c) a few segments along dipole experience reorientation similar to (a), (d) annihilation of screw components in the cross-slip plane producing an equal numberof prismatic loops.

Table III. Independent Slip Systems in FCC, BCC, and HCPMetals, and Those Determined for (001) Single Crystal SrTiO3

from this Study

Crystal structure Slip family Physically distinct Independent

FCC {111} /110S 12 5BCC {110} /111S 12 5HCP {0001} /1120S 3 2SrTiO3 at roomtemperature

{110} /110S{001} /110S

12 5

2352 Journal of the American Ceramic SocietyYang et al. Vol. 92, No. 10

which decompose subsequently into loops under the indentationstress. Screw dipoles stemmed from edge dislocations are gen-erated by bowing out at edge jogs, resulting from dislocationintersection. Edge dipoles are formed by edge trapping as inhigh-temperature deformed sapphire, and by cross slip andpinching-off as in deformed SiFe single crystal. They also ap-pear in dipole trails, which are resulted from dragging of jogsproduced by dislocation intersection.

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October 2009 Deformation of SrTiO3 by Vickers Indentation 2353