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BiFeO3 Ceramics: Processing, Electrical, and Electromechanical Properties

Tadej Rojac,‡,† Andreja Bencan,‡ Barbara Malic,‡,* Goknur Tutuncu,§,* Jacob L. Jones, ¶ ,*

John E. Daniels,k and Dragan Damjanovic‡‡,*‡Electronic Ceramics Department, Jozef Stefan Institute, Ljubljana 1000, Slovenia

§Department of Materials Science and Engineering, University of Florida, Gainesville 32611, Florida

¶ Department of Materials Science and Engineering, North Carolina State University, Raleigh 27695, North Carolina

kSchool of Materials Science and Engineering, University of New South Wales, NSW 2052, Australia

‡‡Ceramics Laboratory, Swiss Federal Institute of Technology – EPFL, Lausanne 1015, Switzerland

Dedicated to Prof. Dr. Marija Kosec, our Maricka, who left us after a long struggle in her tenacious spirit in December 2012.

Bismuth ferrite (BiFeO3), a perovskite material, rich in proper-ties and with wide functionality, has had a marked impact onthe field of multiferroics, as evidenced by the hundreds of arti-

cles published annually over the past 10 years. Studies fromthe very early stages and particularly those on polycrystallineBiFeO3 ceramics have been faced with difficulties in the prepa-ration of the perovskite free of secondary phases. In thisreview, we begin by summarizing the major processing issuesand clarifying the thermodynamic and kinetic origins of theformation and stabilization of the frequently observed second-ary, nonperovskite phases, such as Bi25FeO39 and Bi2Fe4O9.The second part then focuses on the electrical and electrome-chanical properties of BiFeO3, including the electrical conduc-tivity, dielectric permittivity, high-field polarization, and strainresponse, as well as the weak-field piezoelectric properties. Weattempt to establish a link between these properties andaddress, in particular, the macroscopic response of the ceram-ics under an external field in terms of the dynamic interaction

between the pinning centers (e.g., charged defects) and theferroelectric/ferroelastic domain walls.

I. Introduction

B ISMUTH ferrite (BiFeO3) has been subject of intensiveresearch with a large number of articles (over 3000)

published in the past 10 years. Possessing a rare combinationof both (anti)ferromagnetic and ferroelectric properties, coex-

isting at room temperature, and intercoupling between theseproperties, the ferrite has had an important impact in thefield of multiferroics.1 – 3 The great interest in BiFeO3 was ini-

tiated by the study of Wang et al. in 2003.1

The authorsreported a large remanent polarization in epitaxial BiFeO3

thin films, that is, 50 – 60 lC/cm2, which was an order of magnitude higher than the best value reported for bulk BiFe-O3 (6.1 lC/cm2) at that time and, thus naturally, thoughwrongly, associated with the epitaxial strain. Soon after,another group measured a comparable remanent polarizationin single crystals of BiFeO3, that is, ~ 100 lC/cm2 along thepseudocubic (pc) [111]pc direction, revealing the true intrinsicorigin of the large spontaneous polarization of the ferrite.4

The very large spontaneous polarization coupled with thepossibility of manipulating the magnetic ordering with anelectric field have triggered a series of studies on ferroelec-tric/ferroelastic switching in BiFeO3 thin films.5 – 7 Parallelstudies on polycrystalline bulk BiFeO3 ceramics were, how-

ever, faced with more severe problems related to the process-ing and high electrical conductivity of this material,providing additional difficulties in properly understandingthe switching behavior of the ferrite in bulk form.2

Another functional property of BiFeO3 that has been widelyexplored is piezoelectricity. There are several reasons thatplace the ferrite among potentially valuable piezoelectric cera-mic materials. Firstly, BiFeO3 is a lead-free compound and itcontains bismuth, an element whose electronic structure is sim-ilar to Pb,8 but is surprisingly harmless when compared tomost heavy metals, many of which are toxic.9 Secondly, owingto its high Curie temperature (T C = 825°C),3,10 it has beenconsidered for high-temperature piezoelectric applications.11 – 15

Thirdly, its rhombohedral R3c structure in solid solutionswith other perovskites allows the creation of morphotropic

phase boundaries (MPBs) at which the piezoelectric coeffi-cients exhibit a maximum. Recently, all of these reasons initi-ated a large number of studies on chemically modified

D. J. Green—contributing editor

Manuscript No. 34489. Received January 30, 2014; approved April 4, 2014.*Member, The American Ceramic Society.†Author to whom correspondence should be addressed. e-mail: [email protected]

J. Am. Ceram. Soc., 97 [7] 1993–2011 (2014)

DOI: 10.1111/jace.12982

© 2014 The American Ceramic Society

Journal

Feature



BiFeO3 ceramics. Among the most interesting are the BiFe-O3 – PbTiO3 (BFPT)14,16 and BiFeO3 – BaTiO3 (BFBT)15,17,18

systems, which provide both enhanced piezoelectricity and ahigh T C at the MPB, the latter exceeding that of Pb(Zr,Ti)O3

(PZT) (T C ~ 650°C for BFPT, T C ~ 600°C for BFBT andT C ~ 350°C for PZT at the MPB). In addition, a number of other BiFeO3-based lead-free compositions are presently thesubject of intensive studies, including BiFeO3 – REFeO3

(RE = La, Nd, Sm, Gd, Dy),19 – 21 BiFeO3 – AETiO3

(AE = Mg, Ca, Sr),22 – 26 BiFeO3 – Bi0.5K0.5TiO327, and BiFe-

O3 –

Bi(Zn0.5Ti0.5)O3.

28

The piezoelectric properties of manyof these ceramic systems have not yet been characterized sys-tematically.

The processing of single-phase BiFeO3 ceramics is diffi-cult; however, significant progress has been made recently,particularly in relation to the identification of the origins of the frequently formed nonperovskite, secondary phases. Inaddition, the complex relationship between processing anddefects, on one hand, and the high- and weak-field electricaland electromechanical properties, on the other, has beenaddressed to some extent. Up to now, a lot of these newfindings, in particular those relating to processing anddomain-switching behavior, have not been considered suffi-ciently or are even ignored in the literature. Along with theaim of presenting a detailed overview of the past and recent

results on BiFeO3 ceramics, this absence or poor coverageof some important topics was one of the motivations thatled us to prepare a comprehensive article, which alsoincludes new data.

The review comprises two topics on BiFeO3 that are themost controversial and have encountered significant researchbarriers, that is, the processing and the electrical/electrome-chanical properties. The study does not cover magnetic prop-erties and structural aspects. Whereas the crystal structure of BiFeO3 has been reviewed in a study by Catalan and Scott,3

and referred to in earlier publications,10,29,30 the data on themicrostructure and the domain structure are only reportedoccasionally, without general relations to, for example, thepiezoelectric properties of the undoped BiFeO3 bulk materi-als and ceramics in particular.

The section on processing covers the major issuesencountered in the synthesis of BiFeO3: (i) the thermody-namic instability and ternary phase diagrams with impurityoxides; (ii) the reaction kinetics of the Bi2O3 – Fe2O3 system;and (iii) the sublimation/evaporation of bismuth oxide dur-ing high-temperature treatments. By critically analyzingthese interrelated issues we reveal the multiple origins of theoften-formed secondary phases and their persistence in sin-tered BiFeO3 ceramics. The processing section ends by illus-trating the major problems accompanying the conventionalsolid-state processing of BiFeO3, along with general guide-lines to be considered when undertaking the synthesis of theferrite.

BiFeO3 is known as a perovskite with a high electricalconductivity, which is probably the major concern for the

applications of the ferrite and its chemical modifications. Thefirst part of the section on properties is devoted to this issuewhere we discuss the impact of the electrical conductivity onthe “low-frequency” (mHz-to-MHz region) dielectric permit-tivity of BiFeO3 and address some of the inconsistenciesfound in the literature. Based on the available literature dataand our own studies, the character of dielectric relaxation,the type of conductivity, as well as the nature of point andelectronic defects are discussed.

Considering the unfortunate combination of the high elec-trical conductivity and the high coercive field of BiFeO3, it isnot surprising that difficulties in studying domain switchingand piezoelectricity persist. We address the switching behav-ior of the ferrite through an analysis of the polarization-(P – E ) and strain-electric-field (S – E ) hysteresis loops, supported

by synchrotron X-ray diffraction analysis. In addition, wereview and discuss the direct longitudinal piezoelectric

response of BiFeO3 ceramics through its dependence on theexternal driving variables, that is, the amplitude and the fre-quency of the stress.

II. Processing Issues

(1) Thermodynamic Stability, Impurities, and Chemical CompatibilityThe Bi2O3 – Fe2O3 phase diagram, recently revised by Maitreet al.31 and Palai et al.,32 indicates three equilibrium phases

(in order from Fe- to Bi-rich side): the orthorhombicBi2Fe4O9 [space group (s.g.) Pbam, ICSD #20067], the rhom-bohedral perovskite BiFeO3 (s.g. R3c, ICSD #15299), whichdecomposes peritectically to Bi2Fe4O9 and a liquid phase at~ 935°C, and cubic Bi25FeO39 (s.g. I23, ICSD #62719), thelatter exhibiting a peritectic decomposition to Bi2O3 and aliquid phase at ~ 790°C.

The instability of BiFeO3 at elevated temperatures, whichwas only indicated in one of the early Bi2O3 – Fe2O3 phase dia-grams,33,34 was recognized as early as the 1960s when attemptsto determine its Curie temperature were compromised by thereported slow decomposition of the ferrite at T > 700°C, thatis, at temperatures well below its peritectic decomposition.34 – 36

Different decomposition products were reported, includingBi2O3, Fe2O3, and Bi2Fe4O9.35,37 It was suggested that the

decomposition reaction is irreversible,36

taking place attemperatures before the onset of the Bi2O3 sublimation.35

Experimental evidence for the high-temperature instabilityof BiFeO3 was later reported by several other authors. Owingto the difficulty in preparing single-phase perovskite by a ther-mal treatment of a Bi2O3 – Fe2O3 powder mixture, the authorsgenerally refer to the “metastability” of BiFeO3.32,38 – 40 Mor-ozov et al.38 showed that apparently pure BiFeO3 can be pre-pared by reacting a Bi2O3 – Fe2O3 mixture at 850°C for shorttime, that is, 5 – 10 min; however, upon further isothermalannealing at the same temperature for 2 h, the BiFeO3 startedto decompose into Bi25FeO39 and Bi2Fe4O9, revealing moredirectly the thermodynamic instability of the BiFeO3. Thesame kind of decomposition was later confirmed by annealinga sol – gel-derived BiFeO3 at 600°C for 65 h.40 The decomposi-

tion was also observed in BiFeO3 single crystals, confirmingthe intrinsic nature of the instability.32

The thermodynamic instability of BiFeO3 has been clari-fied only recently as part of the experimental and theoreticalstudies by Selbach et al.41 Using the thermodynamic datapublished by Phapale et al.,42 they calculated the Gibbs freeenergy for the equilibrium reaction:

1=49Bi25FeO39 þ 12=49Bi2Fe4O9 ! BiFeO3 (1)

The results of the calculation for the temperature range500 – 1200 K are shown in Fig. 1(b). The calculations showedthat the Bi25FeO39 and Bi2Fe4O9 phases are, though onlyweakly, more thermodynamically stable than the BiFeO3 in

the temperature range 447°C –

767°C, that is, the DrG°m is

slightly positive in this temperature interval, implying thatthe equilibrium reaction (1) is spontaneously shifted to theleft, that is, toward the Bi25FeO39 (Bi-rich) and Bi2Fe4O9

(Bi-poor) phases. This is in agreement with the experimentalresults: a presynthesized BiFeO3 with a small initial amountof the two Bi-rich and Bi-poor secondary phases [Fig. 1(a),25°C] decomposed into these phases within the temperatureinstability range at 700°C [Fig. 1(a), 700°C], whereas the Bi-FeO3 phase reappeared once the sample was annealed abovethe instability range [Fig. 1(a), 775°C and 850°C]. The samebehavior was observed when annealing a Bi2O3 – Fe2O3 mix-ture, instead of BiFeO3, in agreement with earlier stud-ies.38,40

It should be noted that the decomposition reaction

described so far should not be confused with that related tothe formation of BiFeO3 from constituent oxides:

1994 Journal of the American Ceramic Society—Rojac et al. Vol. 97, No. 7



Bi2O3 þ Fe2O3 ! 2BiFeO3 (2)

the calculated enthalpy of which is approximately 70 kJ/mol at 25°C.41 Thus, while reaction (2) will take place at ele-vated temperatures, resulting in the formation of BiFeO3, thecreated BiFeO3 will tend to decompose according to reaction(1) if the sample is annealed between 447°C and 767°C[Fig. 1(b)].

In principle, the equilibrium reaction (1) implies that thecoexistence of the three phases would only be possible at447°C and 767°C, where DrG°

m = 0 [Fig. 1(b)]. The often-reported coexistence of these three phases in a broader tem-perature range [Fig. 1(a)] may be explained by the small driv-ing force (DrG

°m) of the forward and backward reactions,

that is, DrG°m is slightly positive within the instability region

[447°C – 767°C, Fig. 1(b)] and slightly negative in the stableregion at temperatures close to 767°C. This small drivingforce may slow down the reactions from BiFeO3 to the twosecondary phases and backwards, resulting in an apparentthree-phase coexistence.41

In addition to the above-mentioned reasons, Valantet al.43 showed that the coexistence of the perovskite andthe two Bi-rich and Bi-poor phases may be explained ther-modynamically by introducing a third component in theBi2O3 – Fe2O3 system. In terms of the Gibbs phase rule, this

effectively increases the number of independent variables.They showed that the third component, which could be animpurity added in a very small quantity (<1 wt%), couldresult in a large amount (several tens of vol%) of thermody-namically stabilized Bi25FeO39-like and Bi2Fe4O9-like phasesin cases when the impurity is soluble in either Bi25FeO39

and/or Bi2Fe4O9.For the sake of clarity, we explain in Fig. 2 the example

of SiO2 as the third component (impurity).43 In this case itwas assumed and subsequently confirmed by X-ray diffrac-tion (XRD) analysis that the SiO2 will react to formBi12SiO20, that is, a phase belonging to the family of so-called sillenite phases, which are isostructural with c-Bi2O3.44

The next reasonable assumption was that the Bi12SiO20 willexhibit complete or limited solid solubility with the isostruc-tural Bi25FeO39 phase [Fig. 2(b)]. According to the proposedphase diagram [Fig. 2(b)], the incorporation of SiO2 into theBi2O3 – Fe2O3 system would shift the composition into athree-phase region and, thus, three phases will result in equi-librium conditions: BiFeO3, Bi2Fe4O9, and the sillenite phase[see the three-phase field (triangle) marked with an arrow ina small portion of the phase diagram in Fig. 2(b)]. Owing tothe exceptionally elongated and narrow BiFeO3 – Bi2Fe4O9 –

sillenite ternary phase field, which is determined by the spe-cific positions of the phases in the diagram, it was shown by

(a) (b)

Fig. 1. Thermodynamic instability of BiFeO3. (a) X-ray diffraction (XRD) patterns of BiFeO3 at room temperature (25°C) and after isothermalannealing at 700°C, 775°C, and 850°C, showing firstly the decomposition of BiFeO3 into Bi25FeO39 and Bi2Fe4O9 at 700°C, followed by thedisappearance of these secondary phases (i.e., the reappearance of BiFeO3) at T ≥ 775°C. (b) Calculated temperature dependence of the Gibbsfree energy (DrG°

m) of the equilibrium reaction between BiFeO3 and the Bi- and Fe-rich phases [data are plotted for DrH °m = – 1.98 kJ/mol andDrS °m = – 3.62 J/molK; see Ref. (41)]. Note the slightly positive DrG

°m in the region 447°C< T < 776°C, implying that the Bi25FeO39 and

Bi2Fe4O9 are, though weakly, more thermodynamically stable than BiFeO3 in this temperature interval, in agreement with the experimental datashown in panel (a). Reprinted with permission from Ref. [41] Copyright 2009 American Chemical Society.

(a) (b)

Fig. 2. Influence of impurities on the phase composition of a reacted Bi2O3 – Fe2O3 mixture (example of SiO2 as impurity). (a) X-ray diffraction(XRD) patterns of Bi2O3 – Fe2O3 mixture, annealed at 800°C for 5 h, without (0.0%) and with additions of 0.1wt% and 0.5wt% SiO2. (b)Proposed phase relations in the ternary Bi2O3 – Fe2O3 – AOx system, where AOx is SiO2. According to the proposed phase diagram (panel b), a

small amount of SiO2 added to the Bi2O3 –

Fe2O3 mixture should result in a large amount of secondary phases at equilibrium, as confirmedexperimentally (see panel a). Reprinted with permission from Ref. [43] Copyright 2007 American Chemical Society.

July 2014 BiFeO3 Ceramics: Processing and Properties 1995



calculations that a small amount of introduced SiO2 will gen-erate a large amount of the Bi-rich and Bi-poor secondaryphases upon reacting the mixture. The calculations were con-firmed experimentally by deliberately adding 0.1 and 0.5 wt% of SiO2 to a stoichiometric Bi2O3 – Fe2O3 mixture; uponannealing, the additions indeed resulted in large amounts of the secondary phases, which were easily detected by XRD[Fig. 2(a)].

In addition to SiO2, ternary phase diagrams were also pro-posed for other impurities: (i) Al2O3, which, according to

energy-dispersive X-ray spectroscopy analysis (EDXS), incor-porates into the Bi2Fe4O9 phase; and (ii) TiO2, which, inaddition to incorporating into the Bi2Fe4O9, partially substi-tutes for the Fe3+ in BiFeO3.43 Both oxides, added in smallquantities (up to 0.5 wt%), led to the compositional degrada-tion of the BiFeO3 upon annealing, just like in the case of SiO2 [Fig. 2(a)]. Thus, the strong influence of impurities onthe equilibrium phase composition has its origin in thespecific phase relations intrinsic to the Bi2O3 – Fe2O3 – AOx

ternary systems, where AOx represents an impurity oxidethat is more soluble in either Bi25FeO39 and/or Bi2Fe4O9

than in BiFeO3.Though not directly discussed, the effect of impurities on

BiFeO3 should be considered with care, taking into accountthe long list of oxides, including SiO2, ZnO, GeO2, PbO,

Ga2O3, ZrO2, Nb2O5, and MnO, that can form the sillenitephase by reacting with Bi2O3,44 – 47 possibly resulting in thecompositional degradation of the ferrite (analogous to thecase of SiO2, Fig. 2). Therefore, monitoring the type andamount of impurities present in the starting Bi2O3 and Fe2O3

powders is extremely important; as a matter of fact, phase-pure BiFeO3 can be successfully synthesized using startingoxides with ultrahigh purity, that is, >99.999%.43 Anothersource of contamination during the processing of BiFeO3,which is rarely considered in the literature, is the milling pro-cedure. For example, prolonged milling (>10 h) using con-ventional yttria-stabilized zirconia (YSZ) milling media andvials may cause substantial wearing of the milling balls andcontamination of the powder with ZrO2;48 the ZrO2 maythen react upon annealing the contaminated powder to form

the sillenite phase.46

Even though all the impurity oxides may not necessarilyhave such a strong effect as SiO2 or Al2O3, the sensitivity of BiFeO3 to impurities should also be considered in view of the doping. Namely, several authors report on the persistenceof the secondary phases in the ceramics even at low dopinglevels (1 mol%).49 – 51 For example, the study on TiO2,43 oneof the often-reported dopants for BiFeO3,52 – 56 suggests thatonce the maximum solubility of TiO2 in BiFeO3 is exceeded,which may happen even locally in the ceramics due to aninhomogeneous distribution of the dopant, the compositionof the system will fall within the narrow three-phase field[see, as an example, Fig. 2(b)], resulting in the appearance of secondary phases in large quantities. Therefore, when choos-ing a dopant for BiFeO3, one has to be careful particularly

when the dopant tends to be incorporated into the Bi25FeO39

and/or Bi2Fe4O9 since this may stabilize these secondaryphases. Extended discussions on this topic were presented byBernardo et al.55 for the case of the B-site W6+, Nb5+, andTi4+ doping cations. Other problems related to doping BiFe-O3 involve the segregation of the dopant at the grain bound-aries and the suppression of grain growth by doping.56

Another problem that is the direct result of the effect of impurities on the equilibrium phase composition is the chem-ical incompatibility between the BiFeO3 and substrates. SiO2

and Al2O3 are the main constituents of refractory materialsthat are commonly in contact with the ferrite during sinter-ing. This may lead to a compositional degradation of theBiFeO3 at such contacts. In fact, it was proposed that theoften-observed decomposition of BiFeO3 in its paraelectric

phase is due to the chemical incompatibility between the fer-rite and the materials in contact with the sample during

high-temperature treatments,57 rather than being intrinsic tothe BiFeO3.58

For illustrative purposes we show an example of the reac-tion between BiFeO3 and an Al2O3 substrate in Fig. 3. Aftersintering a BiFeO3 pellet on Al2O3 at 900°C an interfacereaction was observed, that is, a marked region appeared onthe crucible, below the original position of the pellet[Fig. 3(b)]. According to XRD analysis, this interface reac-tion coincides with the formation of large quantities of Bi25FeO39 and Bi2Fe4O9 secondary phases [Fig. 3(a), see

900°C and compare with 760°C, where the interface reactionwas not observed). The decomposition of the BiFeO3 is thusactivated at the contact with the Al2O3 and cannot beexplained by considering the intrinsic thermodynamic insta-bility of BiFeO3 as 900°C is well above the instability range[see Fig. 1(b)]. The sublimation/evaporation of Bi2O3 itself isnot likely either since, in this case, one would expect the for-mation of only the Bi-poor phase, that is, Bi2Fe4O9, or even-tually Fe2O3 (see next section and the Bi2O3 – Fe2O3 phasediagram31,32). Similar reactions were reported between BiFe-O3 and a quartz substrate,59 in agreement with the degradingeffect of SiO2 (Fig. 2).

(2) Reaction Kinetics, Sillenite Phase, and

Sublimation/EvaporationIn addition to the thermodynamics discussed in the previoussection, significant progress has also been made in the reactionkinetics of the Bi2O3 – Fe2O3 system. Studies on Bi2O3 – Fe2O3

diffusion couples, performed at 650°C, revealed that bismuthdiffuses several micrometers inside the Fe2O3, while there waslittle backward diffusion of iron into the Bi2O3.60 Assumingcoupled diffusion of the Bi3+ and O2 ions, which preservesthe electroneutrality, and considering that the oxygen diffu-sion in complex oxides is generally faster than the cation diffu-sion,61,62 we can reasonably assume that the formation of BiFeO3 is probably controlled by the diffusion of Bi3+ towardthe Fe2O3. The low diffusion rate of iron is consistent with theavailable diffusion data, calculated for 700°C: the tracer diffu-sion coefficient of Fe3+ in Fe2O3 (D700°C ~ 2.8 9 10 – 25 m2/s)

is five orders of magnitude lower than that of Bi3+

in Bi2O3

(D700°C ~ 6.8 9 10 – 20 m2/s).63,64

The faster bismuth diffusion explains the commonlyobserved microstructural features in BiFeO3 ceramics, that is,an Fe-rich region inside a BiFeO3 grain with the Bi-rich sille-nite phase close to the BiFeO3 grain boundary [Fig 4(a)].60

(a) (b)

(b1)

(b2)

Fig. 3. Chemical instability of BiFeO3 in contact with Al2O3

substrate. (a) X-ray diffraction (XRD) patterns of Bi2O3 – Fe2O3

mixture annealed at 760°C for 6 h and 900°C for 10 h and (b)corresponding photographs of the pellet annealed at 900°C,illustrating the interface reaction between BiFeO3 and the Al2O3

substrate. The pellet was placed on a curved Al2O3 substrate tominimize the BiFeO3/Al2O3 contact. These contacts, at which the

interface reaction was initiated, are marked with arrows in the photo(b2).




This picture may be understood as a kinetically stabilizedphase coexistence, in which incomplete diffusion of the bis-muth through the BiFeO3 grain resulted in an unreacted Fe-rich core. Thus, the initial particle size of the Fe2O3 powder,which determines the diffusion distance, is an importantparameter that should be controlled to avoid unreacted

phases in this system. For comparison, a similar diffusionmechanism governs the formation of BaTiO3 from a mixtureof BaCO3 and TiO2, whereby the growth of BaTiO3 is con-trolled by the diffusion of barium ions through the perovskitelayer.65,66

Based on the observed formation of the sillenite, BiFeO3

and Bi2Fe4O9 phases at the interface of the annealed Bi2O3 –

Fe2O3 diffusion couple, Bernardo et al.60 proposed a reactionmechanism, which is sketched in Fig. 4(b). At the beginningof the reaction, the sillenite is the first phase formed at theBi2O3 – Fe2O3 contact surface, which is consistent with earlierstudies,38,67 and the BiFeO3 phase develops toward the inte-rior of the Fe2O3 particles. In parallel with the BiFeO3, theBi2Fe4O9 phase is also formed, resulting in a diffusion/reac-tion phase sequence of the type: Bi25FeO39/BiFeO3/Bi2Fe4O9/

Fe2O3 (Fig. 4b). As discussed by the authors, it is often thecase that large, regularly shaped Bi2Fe4O9 grains persist inthe final ceramics [see also schematics in Fig. 4(b)]. This wasexplained by the stable growth of these crystals once formed;however, we note that a similar effect could also occur dueto the presence of impurities (see previous section). In addi-tion, one has to consider both the thermodynamics and thekinetics, for example, the Bi2Fe4O9 phase that is kineticallystabilized can grow if the reaction is driven at a temperaturewithin the intrinsic thermodynamic instability region of BiFe-O3 (Fig. 1).

One of the characteristics of the synthesis of BiFeO3 is theeasy formation of the sillenite Bi25FeO39 phase or its iso-structural phases stabilized by the impurity oxides (see theexample of Bi12SiO20 in Fig. 2). Several authors showed inde-

pendently that the Bi25FeO39 is the first reaction productformed upon reacting the Bi2O – Fe2O3 1:1 mixture at temper-atures as low as 400°C – 500°C.38,60,67 In terms of processingthe BiFeO3, this represents a serious problem that is rarelyconsidered in the literature. For example, if this phase doesnot react completely with the Fe-rich counterparts (Fe2O3 orBi2Fe4O9), which may be due to the thermodynamic orkinetic reasons discussed earlier, it can melt at temperatureseven lower than that of pure Bi2O3 (T m(Bi2O3) ~ 830°C andT p(Bi25FeO39) ~ 790°C31,32, where T m and T p refer to melt-ing and peritectic temperature, respectively). The incongruentmelting through peritectic decomposition of the sillenitephase may then lead to (i) uncontrolled Bi2O3 losses associ-ated with the increased vapor pressure of the bismuth oxideabove the melt compared to the expected lower vapor pres-

sure above solid BiFeO3,43,57 and (ii) to segregation of theliquid phase.68

A typical example of the Bi-rich liquid-phase segregationis illustrated in Fig. 5(b). After sintering the ceramics at760°C, that is, below the Bi25FeO39 peritectic point (T < T p),the sillenite phase appeared in the form of lm-sized inclu-sions, well distributed in the BiFeO3 matrix [see the brightspots indicated by an arrow in Fig. 5(a)]. Once the ceramics

were annealed at 820°C, that is, at T > T p, the small silleniteinclusions started to segregate [see the inset of Fig. 5(b)],forming larger inclusions of up to ~ 10 lm in size [Fig. 5(b),see larger bright regions]. Such de-mixing, which is driven bythe liquid-phase formation, spatially prevents the reactionbetween the Bi-rich phase and the Fe-rich counterparts,resulting unavoidably in multiphase ceramics. We note that alarge number of studies found in the literature report sinter-ing temperatures above 820°C; such ceramics may thuscontain nonnegligible concentrations of secondary phases.68

Another important issue is the sublimation or evaporationof the volatile Bi2O3 at elevated temperatures, which mayresult into substoichiometric BiFeO3. Our recent experiments,performed by annealing BiFeO3 ceramics for extended peri-ods (>10 h), suggest Bi2O3 loss at temperatures ≥820°C.68

The results are shown in Fig. 5 and are consistent with theexperiments and conclusions made by Thrall et al.67 Toenrich the atmosphere with bismuth oxide vapors and, thus,minimize the Bi2O3 loss from the BiFeO3 pellet, we embed-ded a presintered BiFeO3 pellet (820°C, 10 h) into a BiFeO3

packing powder. After postannealing at 820°C in the packingpowder, the sillenite phase was clearly detected [Fig. 5(b)]. Incontrast, this phase almost completely disappeared when thepellet was annealed at the same temperature (820°C), but inopen air [Fig. 5(c)], that is, without using the packingpowder. Such a result could be explained by the evaporationof Bi2O3 from the liquid phase that results from the meltingof the sillenite.43,68

A further increase of the annealing temperature to 880 °Cresulted in the formation of large (up to 50 lm) regularly

shaped Bi2Fe4O9 crystals [Fig. 5(d)]. Assuming equilibriumconditions and considering the Bi2O3 – Fe2O3 phase dia-gram,31,32 the system appears as though it shifted to a two-phase BiFeO3 – Bi2Fe4O9 region, corresponding to a Bi/Femolar ratio <1. Thus, the appearance of the Fe-rich phase isdue to the Bi2O3 loss, as was also observed by otherauthors,67,69 and can be represented by a reaction that wasproposed based on thermodynamic calculations:57

8BiFeO3ðsÞ ! 2Bi2Fe4O9ðsÞ þ 4BiOð gÞ þ O2ð gÞ (3)

where “s” and “ g” refer to the solid and vapor phases,respectively. We finally note that the formation of theBi2Fe4O9 crystals, like those shown in Fig. 5(d), may be eas-

ily confused with those stabilized kinetically and/or by impu-rities (see this and the previous section). Finding the origin

(a) (b)

Fig. 4. Reaction kinetics in the Bi2O3 – Fe2O3 system. (a) Scanning electron microscopy (SEM) image of a characteristic distribution of phaseswith different Bi/Fe molar ratios observed in the sintered ceramics, that is, a BiFeO3 grain (light-gray phase) with the Bi-rich sillenite phase atthe BiFeO3 grain boundary (bright phase) and an Fe-rich phase (Bi2Fe4O9 or Fe2O3; dark-gray phase) in the interior of the BiFeO3 grain. Thesample was prepared by a conventional mixed-oxide route at 750°C for 2 h. (b) Proposed reaction pathway mechanism for the solid-statesynthesis of BiFeO3 from Bi2O3 and Fe2O3, based upon diffusion-couple studies and experimental observations (see panel a). Reprinted withpermission from Ref. [60] Copyright 2011 Elsevier Ltd.




of this secondary phase in BiFeO3 is likely to be a nontrivialexercise.

(3) Processing MethodsOwing to the particular thermodynamics and kinetics, thepreparation of single-phase BiFeO3 is notoriously difficult.

Numerous and often unsuccessful attempts have been madeusing a variety of different processing methods. In additionto the conventional solid-state method, there are manyreports on hydrothermal synthesis, sol – gel, rapid liquid-phasesintering, mechanochemical synthesis, precipitation method,combustion synthesis, and high-pressure synthesis (due to thelarge number of publications for each method, references arenot given). It is beyond the scope of this study to review allof the proposed synthesis methods; instead, we will addresssome key problems related to the conventional solid-state

synthesis of BiFeO3, which may be general and applicable toother cases. In the second part of this section, we will thenintroduce our approach, that is, mechanochemical activation.

There are several problems related to the conventionalmulticalcination processing of BiFeO3. The following exam-ple from our own work refers to BiFeO3 prepared using thestandard ceramic procedure, including the premilling of

Bi2O3 and Fe2O3 powders in isopropanol using planetarymilling and YSZ milling media, the mixing of the two pow-ders and calcination, followed by a milling step. A typicalsintering curve of a homogenized Bi2O3 – Fe2O3 powder mix-ture with a micrometer particle size (median particle sized 50 = 1.1 lm) is shown in Fig. 6(a) (see “mixture”). To avoidparticle coarsening, which is a general approach, we calcinedthe mixture at three different temperatures between 650°Cand 700°C, that is, before the onset of the densification at800°C [Fig. 6(a), “mixture”]. For all the calcinations the

(a) (b)

Fig. 6. (a) Sintering curve of Bi2O3 – Fe2O3 mixture before (“mixture”) and after calcination at 700°C for 5 h with additional planetary milling(“calcined + milled”). The sintering curve of the mechanochemically activated Bi2O3 – Fe2O3 mixture (“mixture (activated)”) is added for

comparison. (b) X-ray diffraction (XRD) patterns of Bi2O3 –

Fe2O3 mixture calcined at 650°C (6 h), 680°C (4 h), and 700°C (4 h). The pattern of the mechanochemically activated mixture, which was reactive sintered at 760°C for 6 h (“A+RS”), is also added.

(a) (b)

(c) (d)

Fig. 5. Segregation of Bi-rich liquid phase and Bi2O3 sublimation/evaporation. Scanning electron microscopy (SEM) backscattered-electron(BE) images of BiFeO3 ceramics sintered at (a) 760°C for 6 h and (b, c, d) at 820°C for 10 h with an additional 10 h of postannealing at (b)820°C with the sample immersed in the packing powder, (c) 820°C in open air and (d) 880°C in open air (both (c) and (d) are denoted as“nonimmersed”). The inset of (b) shows the segregation of a Bi-rich liquid phase at 820 °C as a result of the melting of the Bi 25FeO39 phase withthe onset at the peritectic temperature (T p~ 790°C). The large, regularly shaped Bi2Fe4O9 crystals, identified after annealing at 880°C (panel d),are consistent with the Bi2O3 sublimation loss (see text for details). Reprinted with permission from Ref. [68] Copyright 2010 American Instituteof Physics.




resulting powders contained a significant amount of second-ary phases, which were easily detected by XRD [Fig. 6(b)].The result is not surprising, taking into account that thereaction is driven within the temperature region of the ther-modynamic instability of BiFeO3, that is, between 447°C and767°C [see Fig. 1(b)].

To achieve a higher reaction yield with a given powdermixture, one of the common approaches is to increase thediffusion rate of the reacting species by increasing the calci-nation temperature. In the case of BiFeO3, temperatures

higher than 767°C would be required to enter the tempera-ture region in which BiFeO3 is stable [see Fig. 1(b)]. Thereare, however, two key problems related to this approach.Firstly, the Bi25FeO39 phase starts to melt at T p ~ 790°C,which may cause the uncontrolled loss of Bi2O3 throughevaporation and/or segregation of the resulting liquid phase,as discussed in the previous section. In fact, the slight expan-sion of the pellet at ~ 790°C, observed on the sintering curveof the Bi2O3 – Fe2O3 mixture [Fig. 6(a), “mixture”], coincideswith the peritectic point of the Bi25FeO39. Secondly, as seenfrom the sintering curve [Fig. 6(a), “mixture”], the densifica-tion of the powder mixture sets in at 800°C. In principle, thislimits the calcination to below this temperature if we intendto avoid particle coarsening and, consequently, additionalmilling steps, which may increase the possibility of the con-

tamination and stabilization of secondary phases by impuri-ties. We also note that the onset of the densification of thepowder after the first calcination and milling appears at evenlower temperatures, that is, 600°C [Fig. 6(a), “calcined +milled”]. Thus, by considering these two arguments, the firstand subsequent calcination steps would necessarily requiretemperatures that fall within the range of the thermodynamicinstability of BiFeO3 or are close to the top-end of thistemperature interval.

To overcome the apparent incompatibility between thereaction and the densification of conventionally processed Bi-FeO3, we recently proposed the use of mechanochemical acti-vation (high-energy milling). This method is capable of providing highly sinterable and reactive precursors, allowingus to merge the two processes into a single step, known as

reactive sintering.70

We believe that reducing the number of processing steps, including milling, is indeed important forthe synthesis of BiFeO3, considering the degrading effect of impurities. The mechanochemical activation of the Bi2O3 –

Fe2O3 powder mixture resulted in an increased sinterability,that is, the sintering curve of the activated mixture [Fig. 6(a),“mixture (activated)”] is comparable to that of the calcinedand milled powder [Fig. 6(a), “calcined + milled”]; at thesame time, it also provided increased reactivity between theBi2O3 and the Fe2O3. By applying reactive sintering, we wereable to obtain BiFeO3 ceramics in one processing step with aminimum content of secondary phases [<1%; Fig. 6(b),“A+RS”] and a low level of impurities, originating from themilling media (W: 130 ppm, Co: 390 ppm). Figure 6(b)(“A+RS”) shows the case of 760°C, however, similar results

were obtained for higher sintering temperatures.68

Theobtained ceramics were able to withstand electric fields ashigh as 180 kV/cm, allowing us to study the domain-switch-ing behavior and piezoelectricity of BiFeO3. Those resultsare presented in Sections IV and V, while studies on thedielectric permittivity are presented in the following section.

III. Dielectric Permittivity, Electrical Conductivity, andDefects

A comprehensive review, including the most important refer-ences on the dielectric permittivity of BiFeO3 across a widefrequency range, has recently been published by Catalan andScott.3 The intrinsic GHz dielectric permittivity of BiFeO3 isactually small, that is, er

’ ~ 30, however, room-temperature

values as high as er’

~ 10000 were reported in the Hz-to-MHzfrequency range, both in ceramics71 and single crystals.72

This is due to the contribution of the electrical conductivity,which, in addition to directly affecting e’’ (true DC conduc-tivity), it may contribute both to low-frequency e’ and e’’

either through hopping or Maxwell – Wagner (interfacial)mechanisms.73,74 The latter was invoked to explain the fre-quency dispersions and high permittivity values of BiFeO3

ceramics and single crystals.3,71,72,75,76 The domain-wallcontribution to the permittivity may be equally important,but it is generally less often discussed.

It is tempting to assign the dissipation mechanism, for

example, hopping or Maxwell –

Wagner, based on the type of the observed dielectric relaxation, even when the covered fre-quency range is limited and two or more mechanisms behavesimilarly.74 Since this is often done for BiFeO3, we shall dis-cuss here an example where, qualitatively, different types of relaxations were observed in BiFeO3 that was processedusing the same procedure, but under slightly different pro-cessing conditions; in this case, the premilling of the startingBi2O3 and Fe2O3 powders (Fig. 7).

In the case shown in Fig. 7(a) both e’ and tand increaseabruptly below 1 Hz with no clear loss peaks. Consideringthe classification of Jonscher,73 we could assign this kind of behavior to the hopping conductivity. As in most experi-ments, however, only a part of the spectrum is measured andpossible loss peaks and divergence of the loss toward the DC

limit may become apparent only below the experimental fre-quency limit. Thus, other mechanisms, such as the Maxwell – Wagner, cannot be easily ruled out.74 A similar dispersion tothat in Fig. 7(a) was measured in the case shown inFig. 7(b), however, a clear peak in tand was observed at1 Hz. Even though a step-like increase in e’ with decreasingfrequency, characteristic for the Maxwell – Wagner relaxation(e.g., in CaCu3Ti4O12

77) and often reported for BiFe-O3,71,72,75,76,78 was not observed, the overall behavior is com-patible with this kind of dispersion mechanism. Finally, inthe case of Fig. 7(c), the tand increases with decreasing fre-quency significantly more than in the other two cases, reach-ing huge values (tand = 1100) in the low-frequency range [seealso Fig. 7(d)]. In this sample, the phase angle (d) betweenthe polarization and the electric field reached ~ 90° in the

low-frequency limit, suggesting purely conductive behavior(the current in phase with the voltage), in which case e’’ isexpected to diverge as the frequency approaches zero.73,74

The true DC conductivity apparently dominates the permit-tivity response of this sample.

Similar types of relaxations to those shown in Figs. 7(a)and (b) are frequently reported for BiFeO3, typically in theHz-to-MHz range (see, e.g., Refs. [76,78 – 84]). The proposeddispersion mechanisms in those cases are often speculative.

Although the origin of the diverse dielectric behaviorsshown in Fig. 7 is not clear and appears rather complex, thedata clearly suggest that the low-frequency dielectric disper-sion of BiFeO3 ceramics is strongly influenced by the pro-cessing conditions. Thus, significant qualitative andquantitative inconsistencies in the literature regarding the

dielectric dispersions in BiFeO3 are not surprising. In mostcases, the identification of the dominant mechanism is notstraightforward, as we illustrate here, and the dielectricresponse of BiFeO3 may be case sensitive and difficult tocontrol.

In a fashion similar to the low-frequency permittivity, theDC conductivity of BiFeO3 ceramics varies significantly fromcase to case. The as-reported specific DC conductivity of Bi-FeO3 ceramics and single crystals at room temperature arespread over several orders of magnitude, typically between~ 102 and ~ 1010 (Ohm m)1.3,56,72,85 – 87 The high conductiv-ity of BiFeO3 is commonly attributed either to the presenceof: (i) reduced Fe3+ (Fe2+ sites), and/or oxygen vacancies(see, e.g., Refs. [11,76,82,88 – 93]), (ii) secondary phases85,88 ora combination of these, whereby mechanism (i) found a con-

sensus over the largest number of studies (not all are refer-enced).




Recent theoretical studies using density-functional theory(DFT) suggested that BiFeO3 is actually a p-type semiconduc-tor.94,95 Under oxygen-rich conditions, typical for normalceramic processing (i.e., in air atmosphere), the creation of Bi and Fe vacancies is energetically favored over oxygenvacancies. Moreover, the cation vacancies form shallow

acceptor levels within the bandgap; for example, the first ion-ization transition for Bi vacancies (V Bi

X?V Bi’+h•) and Fe

vacancies (V FeX?V Fe’+h•) occurs 0.13 and 0.21 eV above the

valence-band edge, respectively. These defects are thus easilyionized, releasing holes and potentially leading to p-type con-duction. On the other hand and as expected, calculationsmade for oxygen-poor conditions predicted oxygen vacanciesas the most favorable intrinsic defects in BiFeO3 with, inter-estingly, the ionization transition (V O

X?V O

••+2e’) well ormoderately below the conduction-band edge, that is, 1.994 or0.6 eV.96 Thus, in comparison with the shallow acceptor lev-els (cation vacancies), the donor levels (oxygen vacancies) arerather deep, so that BiFeO3 samples treated in an oxygen-poor environment (leading to n-type conductivity) are, inprinciple, not expected to have as high conductivities as those

processed under ordinary oxygen-rich conditions ( p-type con-ductivity). For comparison, a reversed situation was foundexperimentally in BaTiO3, which is an oxygen-deficientn-type semiconductor and an insulator under oxygen-excessconditions with a weak p-type character.97 In the absenceof systematic experimental data on BiFeO3 it is useful tolook at other Fe-containing complex oxides. Interestingly,Sr-doped LaFeO3 was found experimentally to exhibit analo-gous behavior to that predicted for BiFeO3, with conductiv-ity that is orders of magnitude lower at low oxygen partialpressures [r(789°C) = 10 (Ohm m)1 at p(O2) = 1020] thanat high partial pressures [r(789°C) = 4000 (Ohm m)1 at p(O2) = 1].98

The theoretical predictions discussed above for the con-ductivity of BiFeO3 are supported by recent experimental

results. In addition to some sporadic studies on BiFeO3 thatsuggest p-type conductivity,78,99 – 102 it was shown more

directly that the conductivity of Ca-doped BiFeO3 could bereduced by several orders of magnitude if the ceramics weresintered in an oxygen-poor atmosphere (e.g., in N2), consis-tent with the p-type conduction behavior.103 Thus, Ca-dopedBiFeO3 is a p-type semiconductor with an activation energy(E a) of ~ 0.27 – 0.4 eV when sintered in oxygen or air, whereas

it is an ionic conductor with E a~ 0.82 –

1.04 eV when sinteredand cooled in nitrogen (Fig. 8). The semiconductivity wasattributed to the mixed Fe3+/Fe4+ valance state. This is inagreement with other studies, which indicate the absence of Fe2+ in BiFeO3-based ceramics,79,91 favoring p-type conduc-tivity.

Fig. 8. DC electric conductivity versus inverse of temperature(Arrhenius plot) for Ca-doped BiFeO3 ceramics sintered in air,oxygen (O2), and nitrogen (N2). When sintered in air or O2 theceramics behave as a p-type semiconductor with activation energy(E a) close to 0.3 eV. Sintering in N2 reduces significantly theconductivity with E a~ 0.8 eV, showing ionic conduction behavior.

Reprinted with permission from Ref. [103] Copyright 2012 AmericanChemical Society.

(a) (b)

(c) (d)

Fig. 7. Dielectric permittivity (e’ ) and dielectric losses (tand) versus field frequency for three types of BiFeO3 ceramics prepared by the reactivesintering of a mechanochemically activated Bi2O3 – Fe2O3 mixture. The three batches differ in the conditions of premilling of the initial Bi2O3 and

Fe2O3 powders: (a) no premilling and (b,c) premilling in a planetary mill at 200 min 1 of rotational frequency for 4 h using (b) isopropanol and(c) acetone. (d) e’ and tand of the three BiFeO3 ceramics plotted on a logarithmic scale. The median particle sizes (d 50) of the starting Bi2O3/Fe2O3 powders were 14.6/11.3 lm (no premilling), 1.6/0.6 lm (isopropanol milling) and 0.8/0.5 lm (acetone milling). All three ceramics werereactively sintered at 760°C for 6 h [details are given in Ref. (68)], resulting in relative densities in the range 92% – 95%. e’ in panel (c) is plottedin the frequency range 200 – 0.6 Hz; below 0.6 Hz, e ’ decreases and approaches zero as the phase angle reaches ~ 90° (not shown on the plot). Thereal part of the permittivity coefficient (e’ ) can be expressed as e’ = (Q0/V 0)(L/Ae0)cos(d) where Q0, V 0, L, A, e0, and d indicate the chargeamplitude, voltage amplitude, sample thickness, sample area, vacuum permittivity, and phase angle, respectively. According to this equation, if d = 90°, e’ becomes zero, while tand diverges (pure DC conductive behavior).




It appears, from recent studies, that the conductivity in Bi-FeO3-related systems can be significantly reduced by thermaltreatments in an oxygen-poor atmosphere; however, thisleaves an open question as to whether the Bi2O3 loss throughsublimation and/or evaporation may be controlled undersuch conditions. In fact, a lack of oxygen during annealingmay enhance the Bi2O3 loss [see Section II(2) and reaction(3)], an issue that has been initiated,69,104 but so far notstudied in detail. A related problem found in argon-annealedBiFeO3 is a poor microstructure with a high level of poros-

ity, which has its origin, apparently, in the loss of Bi2O3.

67

IV. Ferroelectric and Ferroelastic Domain-SwitchingBehavior

(1) High-Field Polarization Hysteresis LoopsStudies on the ferroelectric switching behavior of BiFeO3

have encountered considerable difficulties, which have theirorigin in the unfortunate combination of the high electricalconductivity and the high coercive field of this material, thelatter already predicted by Michel et al. in 1969.29 Due tothe high leakage current and low breakdown field, it is oftenthe case, even in very recent studies, that authors reportsubcoercive (unsaturated) polarization-electric-field (P – E )loops, which are often confused with the switching (satu-

rated) P –

E loops. This problem has recently been addressedby J. F. Scott,105 but has a tendency to persist. We note,however, that both P – E and strain-electric-field (S – E ) loopsmeasured under subcoercive field conditions, if done prop-erly, may contain very rich information about the domain-wall dynamics, dielectric, and piezoelectric behavior of thematerial, including nonlinearity, hysteresis, conductivity con-tributions, or coupled effects.106 These issues are discussedin Section V.

Some representative P – E loops of BiFeO3 from the litera-ture, reminiscent of those resulting from switching in otherferroelectric ceramics, are shown in Fig. 9.85,90,107 In somecases, such as those shown in Figs. 9(a) and (b), the loopsare pinched (see arrows), whereas in other cases [Fig. 9(c)and Refs. (12,108,109)] they appear more open. The three

loops, shown in Fig. 9, exhibit different degree of internalbias, manifested as a shift along the field axis. Despite thecomparable electric-field amplitude (150 kV/cm), the appar-ent remanent polarization (2Pr) in these cases varies by a fac-tor of three, that is, ~ 20 lC/cm2 < 2Pr < ~ 60 lC/cm2,whereas all the loops show a coercive field of E c ~ 75 kV/cm. The significant qualitative and quantitative differences inthe switching behavior among different samples call for asystematic study of the processing – properties relationship inBiFeO3.

Pinched or double ferroelectric hysteresis loops, like theones shown in Fig. 9(a) and (b) and also observed by otherauthors,110 – 112 indicate the presence of domain-wall pinningcenters, which affect the switching behavior and induce agingof the material’s properties.110 Pinched loops are commonlyobserved in acceptor-doped, often called “hard” ferroelectricmaterials, such as Mn and Fe-doped Pb(Zr,Ti)O3

(PZT)113,114 and Mn-doped BaTiO3.115,116 They were alsomeasured in undoped materials, for example, PZT,106, Ba-TiO3,117 and (K,Na)NbO3 (KNN).118 In all these cases, the

loop pinching is usually attributed to the pinning of thedomain walls by charged defects, which are shown orassumed to be acceptor – oxygen – vacancy (V O

••) defect com-plexes (e.g., FeZr,Ti’ – V O

•• in Fe-doped PZT or TiTi’ – V O•• in

reduced undoped BaTiO3119).

Considering the accepted view of “hardening” in theseacceptor-doped perovskites, one may assume that similardomain-wall pinning centers, that is, acceptor – V O

•• com-plexes, probably also operate in BiFeO3. We note, however,that this assumption, which was made previously in the liter-ature,68,110 was based on the macroscopic behavior of BiFe-O3, such as, loop pinching, depinching and aging, whichappears similar to that in other “hard” ferroelectrics (seelater in this section). In the absence of more direct proofs,other pinning mechanisms and/or pinning defects that may

be specific to BiFeO3 cannot be excluded and are discussedat the end of this section.

In terms of the processing – defects relationship, we firstexamine the cooling-rate effect, which may influence theorder/disorder state of the charged defects with respect to thepolarization within the domains, and, consequently, the stateof pinching of the P – E loop.

Several microscopic mechanisms have been proposed toexplain the pinning effect and aging behavior in PZT and Ba-TiO3.113,115 Given sufficient time, the acceptor – V O

•• defectcomplexes may arrange through the diffusion of oxygenvacancies into an ordered configuration by aligning along thespontaneous polarization within a domain;113,115 – 117,119 – 125 inthis ordered state, the complexes clamp the domain walls,giving rise to the observed pinched P – E loops. The domain

walls may be depinned and, thus, the loop may be depin-ched, by converting the arrangement of the defects from theordered to a disordered state. One of the ways to do this isto first heat the material above T C, that is, into the paraelec-tric phase where the defects are relaxed (disordered) due tothe absence of the spontaneous polarization. This is followedby rapid cooling, which provides “freezing” of the disordereddefect state at room temperature. The quenching experimentscan give indirect, but valuable information about the interac-tion between domain walls and pinning defects.106,114

(a) (b) (c)

Fig. 9. Selected polarization – electric-field (P – E ) hysteresis loops of BiFeO3 ceramics from the literature. The ceramics were prepared using (a)sol – gel powder processing [sintered at 800°C in O2 with high heating (30°C/min) and cooling rate (not specified)],90 (b) rapid liquid-phasesintering [sintered at 855°C with high heating (100°C/min) and cooling rate (~ 170°C/min)],85 and (c) solid-state reaction method [sintered at860°C with high cooling rate (quenching medium not reported)].107 Arrows in panels (a) and (b) indicate the pinched state of the loops.

Reprinted with permission from Ref. [90] (Copyright 2006 American Institute of Physics), Ref. [85] (Copyright 2006 Elsevier Ltd) and Ref. [107](Copyright 2005 American Institute of Physics).




Figure 10(a) shows the P – E loops of BiFeO3 ceramicsquenched into water from various temperatures.68 The open-ing of the loop is only achieved by rapid cooling fromT > T C, that is, from 900°C (T C = 825°C). In contrast, minorchanges in the P – E loops are observed upon quenching fromtemperatures T < T C, that is, 450°C and 760°C, as comparedto the initial, nonquenched state (non-Q). Note the signifi-cant reduction in the apparent coercive field and the increasein the remanent polarization after quenching from 900°C. Asdescribed above, these changes suggest disordering of the

pinning centers through quenching, which results in areduced pinning effect and facilitated polarization switchingand domain-wall motion. The fact that the P – E loops openafter quenching from T > T C is consistent with the assump-tion that the material has to be brought into the paraelectricphase, where the defects may be disordered due to theabsence of the spontaneous polarization.

The strong influence of the defects on the switching behav-ior of BiFeO3 is further confirmed by experiments conductedwith different cooling rates [Fig. 10(b)]. Reducing the coolingrate, from quenching in air to slow cooling within a furnace,while keeping the same target temperature (900°C), resultedin a progressive reappearance of the loop pinching and theassociated reduction of Pr, like that observed in Fe-doped“hard” PZT.106 The results are thus consistent with the grad-

ual reordering of the defects as the cooling rate is decreased.It is interesting to note that the loop of the sample quenchedin air [Fig. 10(b), “Q-air”] exhibited a stronger internal biasfield in comparison with the sample slowly cooled within thefurnace [Fig. 10(b), “slow”]; this might suggest that thedefects in the air-quenched sample partially ordered along apreferred direction, which is, in this case, the direction of theapplied electric field during the P – E measurements.113 Wealso note that the initial state before quenching [Fig. 10(a),“non-Q”] could not be completely recovered, even with 1°C/min of cooling rate [Fig. 10(b), “slow”], suggesting that thedefects encountered difficulties in rearranging back into theordered pinning positions or the presence of other irrevers-ible contributions to the polarization. The discrepancies inthe switching behavior of different BiFeO3 samples reported

in the literature, such as the different level of loop pinching(see Fig. 9), may be in part explained by the different coolingrates employed, which are often not reported and controlled.

Based on the pinched P – E loops and the associated depin-ching or de-aging by quenching, indicative of the presence of domain-wall pinning centers, an obvious question opens upas to which defects cause the pinning of domain walls in Bi-FeO3. Various possibilities, which we next discuss, have beenconsidered,68 however, the exact origin is still not clear.

Assuming the scenario of defect complexes, which isknown to play a major role in, for example, acceptor-dopedPZT and BaTiO3,119 possible pinning centers for the domainwalls in BiFeO3 would be the V Bi

’’’ – V O

•• complexes, createddue to the loss of Bi2O3 at elevated temperatures. Thesetypes of complexes, that is, A-site-vacancy – oxygen-vacancypairs, are believed to be responsible for the “hardening”behavior of undoped PZT.125

To evaluate the role of the defects created in BiFeO3 byBi2O3 loss, we performed a set of experiments in which we

controlled the atmosphere during annealing by using a pack-ing powder. The BiFeO3 was thus annealed with or withoutimmersing the ceramics into a packing powder, and the lossof Bi2O3 was estimated through the analysis of the phasecomposition of the ceramics by means of SEM. These resultswere discussed in Section II(2) (Fig. 5). In contrast to theceramics annealed at 820°C in a packing powder [Fig. 5(b)],the BiFeO3 annealed at temperatures ≥820°C in open air(without the packing powder) showed evidence of the Bi2O3

loss [Figs. 5(c) and (d); see also discussion in Section II(2)].The P – E loops of the ceramics annealed with and without

the packing powder are shown in Fig. 11. A pinched hystere-sis loop is observed after annealing the BiFeO3 in a packingpowder [Fig. 11(a), full line], whereas the loop opens up,showing a decreased coercive field and increased remanent

polarization, after annealing in air, that is, without coveringthe pellet with the packing powder [Fig. 11(b)]. Note also thesignificant difference between the loops of the two sampleswhen compared with the same electric-field amplitude(90 kV/cm): the polarization response of the ceramics treatedin the packing powder [Fig. 11(a), dashed line] is muchweaker than that of the sample annealed in air [Fig. 11(b),full red line]. This depinching of the loop by annealing theceramics in the absence of the packing powder is qualita-tively similar to that observed by quenching the ceramics[Fig. 10(a)] and it is thus indicative of a domain-wall depin-ning (details are reported in Ref. [68]). Although withoutdirect evidence, we could infer from these results that thedefects created in BiFeO3 as a result of the Bi2O3 loss, which,considering the analogous case of undoped PZT, could be

the V Bi’’’

– V O

••

complexes, do not cause domain-wall pinning.Instead, a depinning effect is seen as the Bi2O3 is lostthrough evaporation/sublimation. This behavior of BiFeO3

thus appears qualitatively different from that of PZT.In analogy with the TiTi’ – V O

•• complexes in BaTiO3, whichmay be formed during processing at low oxygen partial pres-sures,119 and considering the role of acceptor – V O

•• complexesin the hardening behavior of PZT and BaTiO3,113,115 – 117,119

the next possibility that might be considered for BiFeO3 are

(a) (b)

Fig. 10. Influence of quenching temperature and cooling rate on the domain-switching behavior of BiFeO3. (a) Polarization – electric-field (P – E )hysteresis loops of as-sintered (nonquenched) ceramics (non-Q) and after rapid quenching in water from 450°C, 760°C, and 900°C.68 The holdingtime at each temperature was 5 min. (b) P – E loops of ceramics annealed at 900°C for 5 min, followed by quenching in water (Q-water) and air

(Q-air), and by slow cooling within the furnace with 1°C/min of cooling rate (slow). Reprinted with permission from Ref. [68] Copyright 2010American Institute of Physics.




the FeFe’ – V O•• complexes, where FeFe’ is Fe2+.68 We note

that these complexes have been proposed to explain theswitching behavior of N2-annealed BiFeO3 thin films.126

Based on recent calculations and conductivity measurements(see Section III), however, Fe2+ does not appear likely in Bi-FeO3 processed under regular oxygen-rich sintering condi-tions (air atmosphere).79,91 Those studies rather suggest thepresence of Fe4+. The complex nature of the defects that cre-

ates such a strong pinning effect in the undoped BiFeO3 isstill to be clarified and requires further studies.

It is well accepted that the defect complexes play animportant role in (deliberately or non-) acceptor-doped per-ovskites, such as PZT and BaTiO3.115 – 117,119 – 123,127,128 Eventhough the macroscopic behavior of BiFeO3 (e.g., pinchedP – E loop) is consistent with the behavior of these perovsk-ites, the defect structure may be different in BiFeO3, andother domain-wall pinning mechanisms must be considered,

for example, the pinning defects may be located within thedomain walls. In addition, one should probably also considerother pinning defects, such as dislocations.129,130

In the case of BiFeO3, the defects located within thedomain walls might play a role in the domain-wall pinning.In fact, several studies on BiFeO3 thin films131 – 134 show evi-dence of conductive 180°, 109°, and 71° domain walls, indi-cating the presence of mobile charge carriers at the walls, asassumed originally by Carl and H€ardtl113 and Postnikovet al.135,136 in their microscopic models. In addition, “head-

to-head” domain configurations were directly identified bypiezoresponse-force microscopy (PFM) in fatigued (111)pc

oriented BiFeO3 thin film.7 The associated compensatingcharges at these domain walls may cause pinning, as wasproposed for the BiFeO3 thin films,7 and/or, in principle,may reduce the wall motion under an applied field.113,135,136

On the other hand, the model of Li et al.137 suggests that theconcentration of the neutralizing charged defects within thedomain walls may lead to either an enhancement or reduc-tion of the piezoelectric effect, depending on whether thecharges in the walls are under or overcompensated. Theseresults and the observations for thin films are consistent withthe observed macroscopic behavior of the BiFeO3 ceramics:the domain-switching strain [see Section IV(2)] and the irre-versible domain-wall contribution to d 33 (see Section V) are

considerably enhanced once the domain walls are depinned,which occurs by driving the BiFeO3 ceramics with low driv-ing-field frequencies, that is, <1 Hz.86,138

(2) High-Field Strain Hysteresis LoopsIn contrast to the numerous polarization-switching studies,data on the electric-field-induced strain in BiFeO3 ceramicsunder switching conditions are limited.68,86,112,139 We recentlymeasured a large bipolar strain of up to ~ 0.4% peak-to-peakat a low frequency (0.1 Hz) and high amplitude (140 kV/cm)of the driving field.86 The strain – electric-field (S – E ) loops of BiFeO3, measured with increasing field amplitude, are shownin Fig. 12(a). The loop at 160 kV/cm shows a clear “butter-fly-like” shape [see the inset of Fig. 12(a)], suggesting switch-

ing and movement, particularly of the non-180° domainwalls, which, depending on the switchable lattice strain,involves a significant change in the dimensions of the grainsin ferroelectric ceramics.140 – 142 We note that the pinched-likestrain loops with small remanence, measured at low-fieldamplitudes [see the loop measured at 110 kV/cm denotedwith an arrow in Fig. 12(a)], are consistent with the pinningof domain walls by defects, in agreement with the P – E data(see previous section). The loops evolve progressively with

(a) (b)

Fig. 11. Influence of Bi2O3 loss upon annealing on the domain-switching behavior of BiFeO3. Polarization – electric-field (P – E )hysteresis loops of BiFeO3 sintered at 820°C for 10 h with anadditional 10 h of postannealing at 820°C or 840°C with the ceramicsample (a) immersed in a BiFeO3 packing powder or (b) treated inair (without the packing powder). Note that the loops shown in

panels (a) and (b) were measured at different electric-field amplitudes(140 and 90 kV/cm, respectively). To allow a comparison under thesame electric-field conditions, the P – E loop of the sample annealedin the packing powder and measured at 90 kV/cm is added in panel(a) (see dashed line).

(a) (b)

Fig. 12. (a) Strain – electric-field (S – E ) hysteresis loops of BiFeO3 ceramics measured with increasing field amplitudes at 0.1 Hz (the inset showsthe loop at 160 kV/cm). (b) Peak-to-peak strain as a function of driving-field frequency for 110, 120, and 130 kV/cm. The data shown in panel

(b) were compiled from measurements reported in Ref. [86]. As indicated with arrows in panel (a), the S –

E loops evolve from a pinched-likestate with little remanent strain at low driving-field amplitudes (110 kV/cm) to a de-pinched state at high amplitudes (160 kV/cm).




increasing field, that is, from a pinched state at low fields toan opened or de-pinched state at high fields [see the loopevolution with field in Fig. 12(a)]. In agreement with thepolarization-field studies,68 this loop evolution suggestsprogressive wall depinning possibly by the reorientation (dis-ordering) of defects.

A particular aspect of the high-field strain response of Bi-FeO3 ceramics is its strong dependence on the frequency of the driving field [Fig. 12(b)]. Especially at lower field ampli-tudes [Fig. 12(b), see 110 and 120 kV/cm], the strain exhibits

an abrupt increase below 1 Hz. Two origins were proposedfor the observed release of the strain at low driving frequen-cies: (i) multistep non-180° domain switching, which was pre-dicted by Kubel and Schmidt30 and confirmed in BiFeO3 thinfilms;6,7 and (ii) rearrangement (disordering) of the pinningcenters.86 The latter scenario is consistent with the domain-wall depinning effect confirmed by quenching (see the previ-ous section) and field cycling [see Fig. 12(a) and Ref. (68)].Interestingly, the same strong frequency dependence is alsoseen in the weak-to-moderate field piezoelectric d 33 coefficient(see next section), suggesting that both the high- andweak-to-moderate field responses of BiFeO3 at low drivingfrequencies are strongly affected by the dynamic interactionbetween the domain walls and the pinning centers.

Some authors have reported that the large electric-field-

induced strain in BiFeO3-based thin films is related to theswitching between tetragonal-like and rhombohedral-likephases.143 – 145 This does not seem to be the dominant mecha-nism in polycrystalline bulk BiFeO3. The role of non-180°domain switching in the large strain response of the BiFeO3

ceramics was confirmed by synchrotron X-ray diffractionanalysis. Details of the measurement and analysis methods asapplied to other compositions have been reported previously,respectively, in Refs. [146,147] and [148,149].

We next discuss the results of our own observations of domain reorientation and its potential contribution to theswitching strain of BiFeO3. The studies were performed usingX-ray diffraction analysis and were carried out ex situ onpoled and unpoled samples. The samples were prepared bythe authors using the mechanochemical method [details of

the synthesis procedure are reported in Ref. (68)].In rhombohedral ceramics, such as BiFeO3, the extent of

non-180° domain switching is determined by measuring therelative intensities of (111)pc and (111)pc diffraction peaksalong a particular specimen direction. Figure 13(a) illustratesthe intensity changes of {111}pc reflections as a function of the orientation of the poled BiFeO3 sample relative to theelectric-field direction. Parallel to the electric-field direction,that is, at the angle of 0°, the intensity of the (111)pc peak,

relative to the (111)pc, is stronger than what would beexpected for randomly oriented rhombohedral ceramics (inwhich case the (111) pc:(111)pc intensity ratio would be closeto 1:3). The difference in the present measurement is due todomain switching that occurred during the poling process,which preferentially oriented the rhombohedral polar [111]pc

direction more closely to the electric-field direction. In addi-tion, a decrease in the intensity of the (111)pc peak is evidentwith increasing angle relative to the electric-field direction,that is, toward the angle of 90°, which shows the distribution

of the preferred orientation, or domain texture, present inthe poled ceramics.Using the diffracted intensities, the extent of domain orien-

tation induced by poling can be quantified. For this purpose,the {111}pc peaks were fitted with two Gaussian profiles andthe integrated intensities were extracted. A degree of domainorientation was then calculated using the ratios of the inte-grated intensities for {111}pc peaks in the poled (oriented)state relative to those in the unpoled (nonoriented) state inall sample directions, defined with respect to the electric-field direction. The degree of orientation can be representedas either a multiple of a random distribution (MRD or f 111)or as a fraction of domains reoriented (g111). These valuesquantify the same preferred orientations or orientation distri-butions, but use different units.148

Figure 13(b) shows the MRD ( f 111) or, equivalently, thefraction of [111]pc-oriented domains that were reoriented(g111) as a function of the angle to the electric field in thepoled ceramic BiFeO3. For nonoriented ceramics, we wouldexpect f 111 = 1 and g111 = 0 [as indicated by the dashed linein Fig. 13(b)], whereas lower or higher values, respectively,correspond to ceramics with a decreased or increased volumefraction of [111]pc-oriented domains in a particular direc-tion.149 Thus, with respect to the nonoriented (unpoled)state, the electric field applied to the BiFeO3 during polinglowered the fraction of the [111]pc domains (g111 < 0) atangles >45° away from the field direction, whereas itincreased the fraction (g111 > 0) along directions more closelyoriented to the field direction (<45°). This can be interpretedas a field-induced [111]pc preferred orientation in the ceramic,

which occurs through non-180° domain switching, that is,through 71° and 109° in the case of rhombohedral BiFeO3.The angle where g111 = 0 [Fig. 13(b)] corresponds to the ori-entation along which the average of all [111]pc domains lie inan energetically equivalent state with respect to the fielddirection.

Figure 14 illustrates a simplified schematic of a hypotheti-cal non-180° domain-wall movement or switching processunder an applied field with the corresponding increase in the

(a) (b)

Fig. 13. (a) Portion of X-ray diffraction (XRD) pattern and (b) calculated degree of domain alignment in poled BiFeO3 as a function of theangle to the applied electric field. Before XRD analysis, the BiFeO 3 ceramic sample was poled with electric fields of increasing amplitude from 10to 160 kV/cm in steps of 10 kV/cm. The corresponding S – E loops of this sample are shown in Fig. 12(a). The degree of domain alignment canbe represented as either multiples of random distribution (MRD) for the polar [111]pc direction ( f 111) or the domain switching fraction (g111);

both of these values are shown as separate y-axes in panel b. An initially unpoled state would be represented as f 111 = 1 MRD and g111=0 in alldirections of the polycrystalline ceramic; such a state is shown on panel b as a dashed line.




fraction of [111]pc-oriented domains along the field axis andthe decrease of this fraction along the direction perpendicularto the applied field. Note the accompanying deformation of the grain as switching occurs (denoted as DL on the schemat-ics).

Being coupled to strain, the non-180° domain-wall motioncontributes to the macroscopic strain realized in ferroelectricmaterials by application of the electric field. In Joneset al.,150 a relationship is given through which the non-180°domain-wall orientation distribution, such as that shown inFig. 13(b), and the spontaneous ferroelastic strain can beused to determine the macroscopic strain resulting from suchdomain reorientation. Using this approach, we found that

the degree of the domain orientation for the poled BiFeO3

ceramics [Fig. 13(b)] yields a longitudinal strain contributionof 0.21% relative to the unpoled state. This large strain issimilar to the magnitude of the measured remanent strain,shown in Fig. 12(a) (see strain at zero field), reinforcing theconclusion that the significant contribution to the macro-scopic electric-field-induced strain of BiFeO3 comes from thenon-180° domain-wall motion and ferroelectric/ferroelasticdomain-switching processes (see also schematics in Fig. 14)and not from phase transformation, as reported for epitaxialthin films.

Finally, we note that the magnitude of the switchingstrain of BiFeO3 ceramics is comparable to that measuredin lead-based ferroelectric ceramics with morphotropic phaseboundary composition, such as PZT and PMN-PT.86 This

is exceptional for a simple oxide such as BiFeO3 and opensup opportunities in the search for potentially efficient lead-free systems based on the ferrite. A good, recent example,which confirms this, is the BiFeO3 – BaTiO3 MPB systemwith an excellent piezoelectric response and a high T C(~ 600°C).15,17,18

V. Direct Piezoelectric Response

Despite the intensive research on BiFeO3, an understandingof the macroscopic piezoelectric response of BiFeO3 ceramicshas been progressing slowly. Some of our recent studies,however, reveal a peculiar piezoelectric response of theferrite, which is discussed in detail in this section.138 Cross-checking the values reported for the piezoelectric d 33 coeffi-

cient,12,13,78,85,151 – 153 we found a very broad range of reported d 33 values, that is, from 2 to 60 pm/V. The problem

is primarily the high electric conductivity of the BiFeO3,which prevents the application of sufficiently high electricfields to the material for an efficient and reproducible poling.

A characterization of the direct piezoelectric response of BiFeO3 ceramics, including the field and frequency depen-dence of the d 33 coefficient and the piezoelectric phase angle,was recently carried out by the authors. The details arereported in Rojac et al.138 Here, we focus on the behavior atlow driving stress frequencies, where the piezoelectricresponse of the BiFeO3 exhibits a significant enhancement.

As shown in Fig. 15(a), the ferrite exhibits a strong fre-quency dispersion of d 33 and tand below ~ 1 Hz. In the low-frequency range (0.02 – 1 Hz), the d 33 increases from 26 pC/N(1 Hz) to 35 pC/N (0.02 Hz), that is, by 35%; likewise, thepiezoelectric tand exhibits an abrupt increase from 0.03(1 Hz) to 0.19 (0.02 Hz). The lossy low-frequency piezoelec-tric response of the BiFeO3 is reflected in the stronglyhysteretic charge density (D) – stress (r) loops measured at0.1 and 0.01 Hz [Fig. 15(c)]. In contrast to these two loops,a small hysteresis with tand<0.05 was measured at frequen-cies ≥1 Hz [see the loops measured at 1 Hz and 10 Hz inFig. 15(c)].

The low-frequency d 33 dispersion [Fig. 15(a), <1 Hz] wasfound to be linked to the piezoelectric nonlinearity, that is, itcorrelates with the amplitude of the driving field.138 This is

seen in Fig. 15(b), which shows the plots of d 33 versus stressamplitude for different stress frequencies. An increase of d 33

with increasing amplitude is observed for all the measuredfrequencies; however, there are significant quantitative differ-ences. While the maximum relative increase in d 33 with stressis similar at 10 and 1 Hz [between 12% and 16%; see the rel-ative d 33 at 3.2 MPa in Fig. 15(b)], d 33 increases with field upto 40% and 66% when the material is cycled with 0.1 and0.01 Hz, respectively. The piezoelectric nonlinear response istherefore strongly enhanced at frequencies below 1 Hz, inagreement with the d 33 frequency dispersion [Fig. 15(a)]. Inrelative terms, the maximum increase of d 33 with the drivingstress at the lowest measured frequency (66% at 0.01 Hz) iscomparable to that measured in Nb-doped “soft” PZT andcoarse-grained BaTiO3.154

Macroscopic piezoelectric nonlinearity and hysteresis inferroelectric materials are usually attributed to extrinsic ori-gins, most commonly to the irreversible movement of ferro-electric-ferroelastic non-180° domain walls under an appliedfield.149,154 – 157 This implies a significant contribution of theirreversible domain-wall motion in BiFeO3 at low driving fre-quencies [Fig. 15(b)].138 The same process also gives rise toenergy dissipation in the material and a hysteretic response.Figure 15(d) shows the D – r loops measured at 0.01 Hz withincreasing stress amplitude. As the amplitude is increased, inaddition to an increase in the slope of the loop, which is pro-portional to d 33, the piezoelectric response of the BiFeO3

becomes increasingly hysteretic. The piezoelectric tand wasalso found to increase with increasing stress amplitude (notshown, see Rojac et al.138), which is expected for the irrevers-

ible contribution from non-180° domain-wall displacements.We note, however, that an additional linear (stress-indepen-dent) and frequency-dispersive mechanism operates in thebackground and contributes to the overall d 33 and piezoelec-tric dispersion.138

What distinguishes BiFeO3 from other piezoelectricallynonlinear ceramics, such as the morphotropic PZT, is its“low-frequency” nonlinearity [Fig. 15(b)] and the associateddispersion of d 33 at low driving frequencies [Fig. 15(a)]. Thisbehavior is very different from the rather general logarithmicfrequency dependence of the piezoelectric coefficient and thenonlinearity observed in “soft” morphotropic PZT and someother materials, in which d 33 increases linearly with decreas-ing log(x) where x is the frequency of the driving field.156,158

Considering that the piezoelectric nonlinearity and hysteresis

in BiFeO3 originates from the irreversible non-180° domain-wall motion, the low-frequency dispersion of d 33 suggests

Fig. 14. Schematics of non-180° domain-wall movement orswitching process under a field (E) and the resulting increase in thefraction of [111]pc-oriented domains (g111) along the direction of theapplied electric field where [111]pc is the rhombohedral polar axis.The deformation (DL) of a hypothetical grain that accompanies theswitching process is also illustrated. For simplicity, the angles

between the polarization vectors are drawn as 90°.




that the domain-wall motion in BiFeO3 is different than inthese other ferroelectric materials.

A possible explanation for the particular behavior of BiFe-

O3 is the coupling between the domain-wall motion and theelectric conductivity, which is reflected in the lossy loopsmeasured at high amplitudes and low frequencies of the driv-ing field [Figs. 15(c) and (d)]. Assuming that mobile chargecarriers are accumulated at the domain-wall area [see discus-sion in Section IV(1)], the domain walls will move under anapplied external field, providing there is a sufficiently longfield exposure time (equivalent to low frequencies). This isbecause domain-wall motion requires charge migration, thatis, electrical conduction in the material. The hypothesis isconsistent with the measured nonlinearity and hysteresis atlow driving frequencies [Figs. 15(b),(c),(d)]. In addition, pos-sible local conductivity at domain walls, as found in BiFeO3

thin films,131,133 could affect the macroscopic piezoelectricresponse, its frequency dispersion and hysteresis through the

Maxwell-Wagner piezoelectric effect.159 A strong indicationof this mechanism is the negative piezoelectric phase angle,which was actually measured in BiFeO3 ceramics.138 All thesepossibilities have recently been considered to explain themacroscopic piezoelectric behavior of the ferrite;138 however,more focused studies are needed to confirm them.

VI. Summary

In ceramic form, BiFeO3 exhibits a number of interestingfunctional properties, however, it presents difficulties when itcomes to processing, both of which are systematicallyreviewed in this contribution. An analysis of the thermody-namics and kinetics of this system unveiled the multiple ori-gins of the frequently observed formation and stabilization

of the Bi-rich and Fe-rich secondary phases. These nonper-ovskite phases can be stabilized by:

1. spontaneous decomposition of the BiFeO3 in the tem-perature range 447°C – 767°C according to the reaction[see Eq. (1)]:.

BiFeO3 ! 1=49Bi25FeO39 þ 12=49Bi2Fe4O9 (4)

2. Contamination with impurities that are likely toincorporate into the Bi25FeO39 and/or Bi2Fe4O9

phases, which leads to the stabilization of these phasesaccording to the Bi2O3 – Fe2O3 – AOx ternary phase rela-tions (AOx is an impurity oxide and also includesrefractory oxides, such Al2O3 (AOx = AlO1.5) andSiO2, which are in often contact with the BiFeO3 dur-ing thermal treatments),

3. kinetic reasons, whereby, due to solid-state diffusion-limited processes, the reaction between Bi2O3 andFe2O3 remains incomplete with the two secondaryphases in the ceramics,

4. melting of the easily formed Bi25FeO39 phase, leadingto segregation and de-mixing of the Bi-rich phase andFe-rich reaction counterparts,

5. sublimation and/or evaporation of Bi2O3 at elevatedtemperatures, which leads to the formation of typicallylarge (tenths of lm) Bi2Fe4O9 crystals in the ceramics.

All these second-phase sources may operate simulta-neously, rendering the synthesis of BiFeO3 and the identifica-tion of the origin of the secondary phases rather difficult. An“ideal” processing method for BiFeO3 should take intoconsideration all of these issues.

In contrast to the general belief that Fe2+ and/or oxygenvacancies are responsible for the high electric conductivity

of BiFeO3, recent theoretical and experimental studies sug-gest that it is p-type conductivity for ceramics processed

(a) (b)

(c) (d)

Fig. 15. Direct piezoelectric response of BiFeO3 ceramics. (a) Piezoelectric d 33 coefficient and piezoelectric losses (tand) as a function of thefrequency of the driving stress (measured at 1.1 MPa of peak-to-peak stress amplitude); (b) relative piezoelectric coefficient ( d 33 relative) versusamplitude of the driving stress, measured at 10, 1, 0.1, and 0.01 Hz; 138 (c, d) charge density – stress loops measured with increasing (c) frequencyand (d) amplitude of driving stress. Note the considerable increase of the slope and the hysteresis of the loops as the frequency of the drivingfield is reduced (panel c) or the amplitude is increased (panel d). The static stress was set to either 2.7 MPa (panel a) or 3.7 MPa (panels b, c,and d).




under ordinary oxygen-rich conditions (air atmosphere).Thus, it appears that the conductivity of the ferrite could becontrolled and lowered by a heat treatment in an oxygen-deficient environment, such as in N2, as shown recently forCa-doped BiFeO3. Systematic experiments, such as conduc-tivity versus the partial pressure of oxygen in a wide pressurerange (e.g., p(O2) = 1020

– 1), would be an important steptoward an understanding of the defect chemistry of the fer-rite and its electrical conductivity.

Macroscopic evidences exist to suggest that BiFeO3

behaves as a “hard” ferroelectric material with typicallypinched or constricted P – E and S – E loops. We showed thatthe overall behavior of BiFeO3, including de-aging or depin-ching by quenching and electric-field cycling, is consistentwith the accepted view of hardening in other perovskites,such as the acceptor-doped PZT and BaTiO3. In these casesthe domain walls are pinned by charged defects, involvingoxygen vacancies, and the wall movement and switching in afield is thus reduced. It has to be emphasized, however, thatthe nature of the pinning centers in BiFeO3 may be differentfrom those in acceptor-doped PZT or BaTiO3, despite manysimilarities in the macroscopic behavior.

Once the domain walls are depinned from the defects,which occurs by application of fields of high amplitude andlow frequency, the BiFeO3 produces large field-induced

strains, comparable to those measured in morphotropic lead-based ferroelectric ceramics, such as PZT and PMN-PT. Syn-chrotron XRD analysis confirmed that the large strains inBiFeO3 ceramics are due to the switching and movement of non-180° domain walls, and not due to the motion of phaseboundaries, as reported for thin films.

The piezoelectric response of BiFeO3 is characterized by astrong nonlinearity and hysteresis, which we attribute to irre-versible domain-wall displacements. At high driving stressamplitudes and low stress frequencies, this domain-wall con-tribution reaches levels comparable to those measured indonor-doped PZT and BaTiO3. This is rather unexpected,considering the “hard” nature of the BiFeO3 and the strongpinning of domain walls as seen through the domain-switch-ing behavior. This inconsistency appears to be, in part, rec-

onciled by the particular dependence of the domain-wallcontribution on the frequency of the applied stress. Indeed,the irreversible non-180° domain-wall motion in BiFeO3 isstrongly restricted to low driving frequencies (<1 Hz). Thislow-frequency nonlinear piezoelectric behavior suggests cou-pling between the domain-wall motion and the conductivityand has not been reported so far for other ferroelectrics. Thebehavior could have multiple origins, but might be, in part,caused by the conductive domain walls, which have beenextensively studied in BiFeO3 thin film, but not yet shown toexist in polycrystalline bulk BiFeO3.

Acknowledgments

This work was supported by the Slovenian Research Agency (programme P2-0105 and project J2-5483). TR would like to thank Prof. Dr. Nava Setter forher financial and technical support related to parts of this work. DD acknowl-edges the financial support of FNS-PNR62. JJ and GT acknowledge the U.S.Department of the Army for support under contract number W911NF-09-1-0435. JD acknowledges financial support from an AINSE research fellowshipand ARC DP120103968. Use of the Advanced Photon Source was supportedby the US Department of Energy, Office of Science, Office of Basic EnergySciences, under Contract No. DE-AC02-06CH11357.

References

1J. Wang, J. B. Neaton, H. Zheng, V. Nagarajan, S. B. Ogale, B. Liu, D.Viehland, V. Vaithyanathan, D. G. Schlom, U. V. Waghmare, N. A. Spaldin,K. M. Rabe, M. Wuttig, and R. Ramesh, “Epitaxial BiFeO3 Multiferroic ThinFilm Heterostructures,” Science, 299 [5613] 1719 – 22 (2003).

2J. Yu and J. Chu, “Progress and Prospect for High Temperature Single-

Phased Magnetic Ferroelectrics,” Chi. Sci. Bull., 53 [14] 2097 – 112 (2008).3

G. Catalan and J. F. Scott, “Physics and Applications of Bismuth Ferrite,”Adv. Mater., 21 [24] 2463 – 85 (2009).

4D. Lebeugle, D. Colson, A. Forget, and M. Viret, “Very Large Spontane-ous Electric Polarization in BiFeO3 Single Crystals at Room Temperature and

Its Evolution Under Cycling Fields,” Appl. Phys. Lett., 91 [2] 022907, 3pp(2007).

5N. Balke, S. Choudhury, S. Jesse, M. Huijben, Y. H. Chu, A. P. Baddorf,L. Q. Chen, R. Ramesh, and S. V. Kalinin, “Deterministic Control of Ferro-elastic Switching in Multiferroic Materials,” Nat. Nanotechnol., 4 [12] 868 – 75(2009).

6S. H. Baek, H. W. Jang, C. M. Folkman, Y. L. Li, B. Winchester, J. X.Zhang, Q. He, Y. H. Chu, C. T. Nelson, M. S. Rzchowski, X. Q. Pan, R. Ra-mesh, L. Q. Chen, and C. B. Eom, “Ferroelastic Switching for NanoscaleNon-Volatile Magnetoelectric Devices,” Nat. Mater., 9 [4] 309 – 14 (2010).

7S. H. Baek, C. M. Folkman, J. W. Park, S. Lee, C. W. Bark, T. Tybell,

and C. B. Eom, “The Nature of Polarization Fatigue in BiFeO 3,” Adv. Mater.,23 [14] 1621 – 5 (2011).

8J. R€odel, W. Jo, K. T. P. Seifert, E. M. Anton, T. Granzow, and D. Dam- janovic, “Perspective on the Development of Lead-Free Piezoceramics,” J.Am. Ceram. Soc., 92 [6] 1153 – 77 (2009).

9R. Mohan, “Green Bismuth,” Nat. Chem, 2 [4] 336 (2010).10

J. M. Moreau, C. Michel, R. Gerson, and W. J. James, “Ferroelectric Bi-FeO3 X-Ray and Neutron Diffraction Study,” J. Phys. Chem. Solids, 32 [6]1315 – 20 (1971).

11G. L. Yuan and S. W. Or, “Enhanced Piezoelectric and PyroelectricEffects in Single-Phase Multiferroic Bi1xNdxFeO3 (x = 0 – 0.15) Ceramics,”Appl. Phys. Lett., 88 [6] 062905, 3pp (2006).

12C. Sun, X. Chen, J. Wang, G. Yuan, J. Yin, and Z. Liu, “Structure and

Piezoelectric Properties of BiFeO3 and Bi0.92Dy0.08FeO3 Multiferroics at HighTemperature,” Solid State Commun., 152 [14] 1194 – 8 (2012).

13X. Chen, Y. Wang, Y. Yang, G. Yuan, J. Yin, and Z. Liu, “Structure,Ferroelectricity and Piezoelectricity Evolutions of Bi1xSmxFeO3 at VariousTemperatures,” Solid State Commun., 152 [6] 497 – 500 (2012).

14T. P. Comyn, S. P. McBride, and A. J. Bell, “Processing and Electrical

Properties of BiFeO3 – PbTiO3 Ceramics,” Mater. Lett., 58 [30] 3844 – 6 (2004).15H. Yang, C. Zhou, X. Liu, Q. Zhou, G. Chen, H. Wang, and W. Li,

“Structural, Microstructural and Electrical Properties of BiFeO3 – BaTiO3

Ceramics With High Thermal Stability,” Mater. Res. Bull., 47 [12] 4233 – 9(2012).

16T. P. Comyn, T. Stevenson, and A. J. Bell, “Piezoelectric Properties of Bi-FeO3 – PbTiO3 Ceramics,” J. Phys. IV France, 128 [1] 3 – 7 (2005).

17S. O. Leontsev and R. E. Eitel, “Dielectric and Piezoelectric Properties in

Mn-Modified (1 – X)BiFeO3 – XBaTiO3 Ceramics,” J. Am. Ceram. Soc., 92 [12]2957 – 61 (2009).

18S. O. Leontsev and R. E. Eitel, “Origin and Magnitude of the Large Pie-zoelectric Response in the Lead-Free (1 – X)BiFeO3 – XBaTiO3 Solid Solution,”J. Mater. Res., 26 [1] 9 – 17 (2011).

19S. Karimi, I. M. Reaney, Y. Han, J. Pokorny, and I. Sterianou, “CrystalChemistry and Domain Structure of Rare-Earth Doped BiFeO3 Ceramics,” J.Mater. Sci., 44 [19] 5102 – 12 (2009).

20S. Fujino M. Murakami, V. Anbusathaiah, S.-H. Lim, V. Nagarajan, C. J.Fennie, M. Wuttig, L. Salamanca-Riba, and I. Takeuchi, “Combinatorial Dis-covery of a Lead-Free Morphotropic Phase Boundary in a Thin-Film Piezo-electric Perovskite,” Appl. Phys. Lett., 92 [20] 202904, 3pp (2008).

21

D. Kan, L. Pa ´ lova ´ , V. Anbusathaiah, C. J. Cheng, S. Fujino, V. Nagara- jan, K. M. Rabe, and I. Takeuchi, “Universal Behavior and Electric-Field-Induced Structural Transition in Rare-Earth-Substituted BiFeO3,” Adv. Funct.Mater., 20 [7] 1108 – 15 (2010).

22T. L. Ivanova and V. V. Gagulin, “Dielectric Properties in the MicrowaveRange of Solid Solutions in the BiFeO3 – SrTiO3 System,” Ferroelectrics, 265[1] 241 – 6 (2002).

23N. Itoh, T. Shimura, W. Sakamoto, and T. Yogo, “Effects of SrTiO 3 Con-tent and Mn Doping on Dielectric and Magnetic Properties of BiFeO3 – SrTiO3

Ceramics,” J. Ceram. Soc. Jpn., 117 [9] 939 – 43 (2009).24Z. Z. Ma, Z. M. Tian, J. Q. Li, C. H. Wang, S. X. Huo, H. N. Duan, and

S. L. Yuan, “Enhanced Polarization and Magnetization in Multiferroic (1 – X)BiFeO3 – XSrTiO3 Solid Solution,” Solid State Sci., 13 [12] 2196 – 200 (2011).

25J. Wu and J. Wang, “Multiferroic Behavior of BiFeO3 – RTiO3 (Mg, Sr,Ca, Ba, and Pb) Thin Films,” J. Appl. Phys., 108 [2] 026101, 3pp (2010).

26Q. Q. Wang, Z. Wang, X. Q. Liu, X. M. Chen, and D. W. Johnson,“Improved Structure Stability and Multiferroic Characteristics in CaTiO3-Modified BiFeO3 Ceramics,” J. Am. Ceram. Soc., 95 [2] 670 – 5 (2012).

27M. I. Morozov, M. A. Einarsrud, and T. Grande, “Polarization and Strain

Response in Bi0.5K0.5TiO3 –

BiFeO3 Ceramics,” Appl. Phys. Lett., 101 [25]252904, 4pp (2012).

28K. Yazawa, S. Yasui, H. Morioka, T. Yamada, H. Uchida, A. Gruver-

man, and H. Funakubo, “Composition Dependence of Crystal Structure andElectrical Properties for Epitaxial Films of Bi(Zn1/2Ti1/2)O3 – BiFeO3 SolidSolution System,” J. Ceram. Soc. Jpn., 118 [8] 659 – 63 (2010).

29C. Michel, J. M. Moreau, G. Achenbach, R. Gerson, and W. J. James,“The Atomic Structure of BiFeO3,” Solid State Commun., 7 [9] 701 – 4 (1969).

30F. Kubel and H. Schmid, “Structure of a Ferroelectric and Ferroelastic

Monodomain Crystal of the Perovskite BiFeO3,” Acta Crystallogr. B, 46 [6]698 – 702 (1990).

31A.Maıtre, M. Francois, and J. C. Gachon,“Experimental Studyof theBi2O3 –

Fe2O3 Pseudo-BinarySystem,” J. Phase Equilib. Diffus., 25 [1] 59 – 67 (2004).32R. Palai, R. S. Katiyar, H. Schmid, P. Tissot, S. J. Clark, J. Robertson, S.

A. T. Redfern, G. Catalan, and J. F. Scott, “Β Phase and c-b Metal-InsulatorTransition in Multiferroic BiFeO3,” Phys. Rev. B, 77 [1] 014110, 11pp (2008).

33E. I. Speranskaya, V. M. Skorikov, E. Y. Rode, and V. A. Terekhova,

“The Phase Diagram of the System Bismuth Oxide – Ferric Oxide,” Bull.Acad. Sci. U.S.S.R, 5 [87] 3 – 4 (1965).

34S. M. Skinner, “Magnetically Ordered Ferroelectric Materials,” IEEE Trans. PMP, 6 [2] 68 – 90 (1970).




35V. S. Filip’ev, N. P. Smolyaninov, E. G. Fesenko, and I. N. Belyaev,

“Synthesis of BiFeO3 and Determination of the Unit Cell,” Sov. Phys. – Crys-tallogr., 5, 913 – 4 (1960).

36S. A. Fedulov, Y. N. Venevtsev, G. S. Zhdanov, and E. G. Smazhevskaya,

“High-Temperature X-Ray and Thermal-Analysis Results for BismuthFerrite,” Sov. Phys. – Crystallogr., 6 , 640 – 1 (1961).

37G. D. Achenbach, W. J. James, and R. Gerson, “Preparation of Single-Phase Polycrystalline BiFeO3,” J. Am. Ceram. Soc., 50 [8] 437 – 7 (1967).

38M. I. Morozov, N. A. Lomanova, and V. V. Gusarov, “Specific Features

of BiFeO3 Formation in a Mixture of Bismuth(III) and Iron(III) Oxides,”Russ. J. Gen. Chem., 73 [11] 1676 – 80 (2003).

39M. W. Lufaso, T. A. Vanderah, I. M. Pazos, I. Levin, R. S. Roth, J. C.Nino, V. Provenzano, and P. K. Schenck, “Phase Formation, Crystal Chemis-

try, and Properties in the System Bi2O3 –

Fe2O3 –

Nb2O5,” J. Solid State Chem.,179 [12] 3900 – 10 (2006).

40T. T. Carvalho and P. B. Tavares, “Synthesis and Thermodynamic Stabil-

ity of Multiferroic BiFeO3,” Mater. Lett., 62 [24] 3984 – 6 (2008).41S. M. Selbach, M. A. Einarsrud, and T. Grande, “On the Thermodynamic

Stability of BiFeO3,” Chem. Mater., 21 [1] 169 – 73 (2009).42S. Phapale, R. Mishra, and D. Das, “Standard Enthalpy of Formation

and Heat Capacity of Compounds in the Pseudo-Binary Bi2O3 – Fe2O3 Sys-tem,” J. Nucl. Mater., 373 [1 – 3] 137 – 41 (2008).

43M. Valant, A. K. Axelsson, and N. Alford, “Peculiarities of a Solid-StateSynthesis of Multiferroic Polycrystalline BiFeO3,” Chem. Mater., 19 [22] 5431 – 6 (2007).

44H. A. Harwig, “On the Structure of Bismuthsesquioxide: The a , b , c , and

d-Phase,” Z. Anorg. Allg. Chem., 444 [1] 151 – 66 (1978).45L. G. Sillen, “X-Ray Studies on Bismuth Trioxide,” Ark. Kemi. Mineral.

Geol., 12A [18] 1 – 15 (1937).46

E. M. Levin and R. S. Roth, “Polymorphism of Bismuth Sesquioxide. IIEffect of Oxide Additions on the Polymorphism of Bi2O3,” J. Res. Nat. Bur.Stand. A, 68A [2] 197 – 206 (1964).

47D. C. Craig and N. C. Stephenson, “Structural Studies of Some Body-Centered Cubic Phases of Mixed Oxides Involving Bi2O3: The Structures of Bi25FeO40 and Bi38ZnO60,” J. Solid State Chem., 15 [1] 1 – 8 (1975).

48T. Rojac, M. Kosec, B. Malic, and J. Holc, “Mechanochemical Synthesisof NaNbO3,” Mater. Res. Bull., 40 [2] 341 – 5 (2005).

49B. Yu, M. Li, J. Wang, L. Pei, D. Guo, and X. Zhao, “Enhanced Electri-

cal Properties in Multiferroic BiFeO3 Ceramics Co-Doped by La3+and V5+,”J. Phys. D, 41 [18] 185401, 5pp (2008).

50W. S. Kim, Y. K. Jun, K. H. Kim, and S. H. Hong, “Enhanced Magneti-zation in Co and Ta-Substituted BiFeO3 Ceramics,” J. Magn. Magn. Mater.,321 [19] 3262 – 5 (2009).

51F. Azough, R. Freer, M. Thrall, R. Cernik, F. Tuna, and D. Collison,“Microstructure and Properties of Co-, Ni-, Zn-, Nb- and W-Modified Multif-erroic BiFeO3 Ceramics,” J. Eur. Ceram. Soc., 30 [3] 727 – 36 (2010).

52Z. Chaodan, Y. Jun, Z. Duanming, Y. Bin, W. Yunyi, W. Longhai, W.Yunbo, and Z. Wenli, “Processing and Ferroelectric Properties of Ti-DopedBiFeO3 Ceramics,” Integr. Ferroelectr., 94 [1] 31 – 6 (2007).

53S. J. Kim, S. H. Han, H. G. Kim, A. Y. Kim, J. S. Kim, and C. l. Cheon,

“Multiferroic Properties of Ti-Doped BiFeO3 Ceramics,” J. Korean Phys. Soc.,56 [12] 439 – 42 (2010).54

I. O. Troyanchuk, A. N. Chobot, O. S. Mantytskaya, and N. V. Tereshko,“Magnetic Properties of Bi(Fe1 xMx)O3 (M = Mn, Ti),” Inorg. Mater., 46

[4] 424 – 8 (2010).55M. S. Bernardo, T. Jardiel, M. Peiteado, A. C. Caballero, and M. Villegas,

“Sintering and Microstuctural Characterization of W6+, Nb5+ and Ti4+ Iron-Substituted BiFeO3,” J. Alloy. Compd., 509 [26] 7290 – 6 (2011).

56M. S. Bernardo, T. Jardiel, M. Peiteado, F. J. Mompean, M. Garcia-Her-

nandez, M. A. Garcia, M. Villegas, and A. C. Cabalero, “Intrinsic Composi-tional Inhomogeneities in Bulk Ti-Doped BiFeO3: MicrostructureDevelopment and Multiferroic Properties,” Chem. Mater., 25 [9] 1533 – 41(2013).

57S. M. Selbach, T. Tybell, M. A. Einarsrud, and T. Grande, “Phase Transi-

tions, Electrical Conductivity and Chemical Stability of BiFeO 3 at High Tem-peratures,” J. Solid State Chem., 183 [5] 1205 – 8 (2010).

58R. Haumont, I. Kornev, S. Lisenkov, L. Bellaiche, J. Kreisel, and B.Dkhil, “Phase Stability and Structural Temperature Dependence in PowderedMultiferroic BiFeO3,” Phys. Rev. B, 78 [13] 134108, 8pp (2008).

59J. D. Bucci, B. K. Robertson, and W. J. James, “The Precision Determina-tion of the Lattice Parameters and the Coefficients of Thermal Expansion of BiFeO3,” J. Appl. Crystallogr., 5 [3] 187 – 91 (1972).

60M. S. Bernardo, T. Jardiel, M. Peiteado, A. C. Caballero, and M. Villegas,“Reaction Pathways in the Solid State Synthesis of Multiferroic BiFeO3,” J.Eur. Ceram. Soc., 31 [16] 3047 – 53 (2011).

61M. V. Slinkina and G. I. Doncov, “Difuzija Sovstvenih Komponentov V

Keramike Cirkonata-Titanata Svinca.,” Neorgani ceskie Materiali., 28 [3] 567 – 70 (1992).

62B. Malic, D. Jenko, J. Holc, M. Hrovat, and M. Kosec, “Synthesis of

Sodium Potassium Niobate: A Diffusion Couples Study,” J. Am. Ceram. Soc.,91 [6] 1916 – 22 (2008).

63P. J. Harrop, “Self-Diffusion in Simple Oxides (a Bibliography),” J.

Mater. Sci., 3 [2] 206 – 22 (1968).64

J. L. Mukherjee and F. F. Y. Wang, “Kinetics of Solid-State Reaction of Bi2O3 and Fe2O3,” J. Am. Ceram. Soc., 54 [1] 31 – 4 (1971).

65A. Beauger, J. C. Mutin, and J. C. Niepce, “Synthesis Reaction of Metati-

tanate BaTiO3,” J. Mater. Sci., 18 [12] 3543 – 50 (1983).66M. R€ossel, H. R. H€oche, H. S. Leipner, D. V€oltzke, H. P. Abicht, O.

Hollricher, J. M€uller, and S. Gablenz, “Raman Microscopic Investigations of

BaTiO3 Precursors With Core-Shell Structure,” Anal. Bioanal. Chem., 380 [1]157 – 62 (2004).

67M. Thrall, R. Freer, C. Martin, F. Azough, B. Patterson, and R. J. Cer-nik, “An in Situ Study of the Formation of Multiferroic Bismuth FerriteUsing High Resolution Synchrotron X-Ray Powder Diffraction,” J. Eur.Ceram. Soc., 28 [13] 2567 – 2 (2008).

68T. Rojac, M. Kosec, B. Budic, N. Setter, and D. Damjanovic, “StrongFerroelectric Domain-Wall Pinning in BiFeO3 Ceramics,” J. Appl. Phys., 108[7] 074107, 8pp (2010).

69M. Li and J. L. MacManus-Driscoll, “Phase Stability, Oxygen Nonstoichi-ometry and Magnetic Properties of BiFeO3d,” Appl. Phys. Lett., 87 [25]252510, 3pp (2005).

70L. B. Kong, J. Ma, T. S. Zhang, W. Zhu, and O. K. Tan, “Pb(Zr xTi1 x)

O3 Ceramics Via Reactive Sintering of Partially Reacted Mixture Produced bya High-Energy Ball Milling Process,” J. Mater. Res., 16 [6] 1636 – 43 (2001).

71S. Kamba, D. Nuzhnyy, M. Savinov, J. Sebek, J. Petzelt, J. Proleska, R.

Haumont, and J. Kreisel, “Infrared and Terahertz Studies of Polar Phononsand Magnetodielectric Effect in Multiferroic BiFeO3 Ceramics,” Phys. Rev. B,75 [2] 024403, 7pp (2007).

72J. Lu, A. G€unther, F. Schrettle, F. Mayr, S. Krohns, P. Lunkenheimer, A.Pimenov, V. D. Travkin, A. A. Mukhin, and A. Loidl, “On the Room Tem-perature Multiferroic BiFeO3: Magnetic, Dielectric and Thermal Properties,”Eur. Phys. J. B, 75 [4] 451 – 60 (2010).

73A. K. Jonscher, Dielectric Relaxation in Solids. Chelsea Dielectrics Press,

U.K., 1983.74

M. I. Morozov and D. Damjanovic, “Charge Migration in Pb(Zr,Ti)O3

Ceramics and Its Relation to Ageing, Hardening, and Softening,” J. Appl.Phys., 107 [3] 034106, 10pp (2010).

75R. Mazumder, S. Ghosh, P. Mondal, D. Bhattacharya, S. Dasgupta, N.

Das, A. Sen, A.K. Tyagi, M. Sivalumar, T Takami, and H. Ikuta, “ParticleSize Dependence of Magnetization and Phase Transition Near Tn in Multifer-roic BiFeO3,” J. Appl. Phys., 100 [3] 033908, 9pp (2006).

76S. Hunpratub, P. Thongbai, T. Yamwong, R. Yimnirun, and S. Maensiri,“Dielectric Relaxations and Dielectric Response in Multiferroic BiFeO3

Ceramics,” Appl. Phys. Lett., 94 [6] 062904, 3pp (2009).77P. Lunkenheimer, R. Fichtl, S. G. Ebbinghaus, and A. Loidl, “Nonintrin-

sic Origin of the Colossal Dielectric Constants in CaCu3Ti4O12,” Phys. Rev. B,70 [17] 172102, 4pp (2004).

78Z. Dai and Y. Akishige, “Electrical Properties of Multiferroic BiFeO3

Ceramics Synthesized by Spark Plasma Sintering,” J. Phys. D, 43 [44] 445403,5pp (2010).

79J. Chen, X. Xing, A. Watson, W. Wang, R. Yu, J. Deng, L. Yan, C. Sun,

and X. Chen “Rapid Synthesis of Multiferroic BiFeO3 Single-CrystallineNanostructures,” Chem. Mater., 19 [15] 3598 – 600 (2007).

80M. Kumar, K. L. Yadav, and G. D. Varma, “Large Magnetization and

Weak Polarization in Sol-Gel Derived BiFeO3 Ceramics,” Mater. Lett., 62 [8 – 9] 1159 – 61 (2008).

81R. Mazumder, D. Chakravarty, D. Bhattacharya, and A. Sen, “Spark

Plasma Sintering of BiFeO3,” Mater. Res. Bull., 44 [3] 555 – 9 (2009).82

R. Mazumder and A. Sen, “Effect of Pb-Doping on Dielectric Properties

of BiFeO3 Ceramics,” J. Alloy Compd., 475 [1 –

2] 577 –

80 (2009).83H. O. Rodrigues, G. F. M. Pires Junior, A. J. M. Sales, P. M. O. Silva, B.F. O. Costa, P. Alcantara Jr, S. G. C. Moreira, and A. S. B. Sombra, “BiFeO3

Ceramic Matrix With Bi2O3 or PbO Added: M€ossbauer, Raman and DielectricSpectroscopy Studies,” Phys. B, 406 [13] 2532 – 9 (2011).

84Z. H. Dai and Y. Akishige, “BiFeO3 Ceramics Synthesized by SparkPlasma Sintering,” Ceram. Int., 38 , S403 – 6 (2012).

85G. L. Yuan, S. W. Or, Y. P. Wang, Z. G. Liu, and J. M. Liu, “Prepara-

tion and Multi-Properties of Insulated Single-Phase BiFeO3 Ceramics,” Solid State Commun., 138 [2] 76 – 81 (2006).

86T. Rojac, M. Kosec, and D. Damjanovic, “Large Electric-Field Induced

Strain in BiFeO3 Ceramics,” J. Am. Ceram. Soc., 94 [12] 4108 – 11 (2011).87E. Markiewicz, B. Hilczer, M. Baszyk, A. Pietraszko, and E. Talik,

“Dielectric Properties of BiFeO3 Ceramics Obtained From MechanochemicallySynthesized Nanopowders,” J. Electroceram., 27 [3 – 4] 154 – 61 (2011).

88Y. P. Wang, L. Zhou, M. F. Zhang, X. Y. Chen, J. M. Liu, and Z. G.

Liu, “Room-Temperature Saturated Ferroelectric Polarization in BiFeO3

Ceramics Synthesized by Rapid Liquid Phase Sintering,” Appl. Phys. Lett., 84[10] 1731 – 3 (2004).

89A. K. Pradhan, K. Zhang, D. Hunter, J. B. Dadson, G. B. Loutts, P.Bhattacharya, R. Katiyar, J. Zhang, D. J. Sellmyer, U. N. Roy, Y. Cui, andA. Burger, “Magnetic and Electrical Properties of Single-Phase MultiferroicBiFeO3,” J. Appl. Phys., 97 [9] 093903, 4pp (2005).

90F. Chen, Q. F. Zhang, J. H. Li, Y. J. Qi, C. J. Lu, X. B. Chen, X. M.

Ren, and Y. Zhao, “Sol-Gel Derived Multiferroic BiFeO3 Ceramics WithLarge Polarization and Weak Ferromagnetism,” Appl. Phys. Lett., 89 [9]092910, 3pp (2006).

91G. L. Yuan, K. Z. Baba-Kishi, J. M. Liu, S. W. Or, Y. P. Wang, and Z.G. Liu, “Multiferroic Properties of Single Phase Bi0.85La0.15FeO3 Lead-FreeCeramics,” J. Am. Ceram. Soc., 89 [10] 3136 – 9 (2006).

92H. Ke, W. Wang, Y. Wang, H. Zhang, D. Jia, Y. Zhou, X. Lu, and P.

Withers, “Dependence of Dielectric Behavior in BiFeO3 Ceramics on IntrinsicDefects,” J. Alloy. Compd., 541 , 94 – 8 (2012).

93B. Yu, M. Li, J. Liu, D. Guo, L. Pei, and X. Zhao, “Effects of Ion Dop-

ing at Different Sites on Electrical Properties of Multiferroic BiFeO3 Ceram-ics,” J. Phys. D, 41 [6] 065003, 4pp (2008).

94Z. Zhang, P. Wu, L. Chen, and J. Wang, “Density Functional Theory Plus

U Study of Vacancy Formations in Bismuth Ferrite,” Appl. Phys. Lett., 96

[23] 232906, 3pp (2010).




95T. R. Paudel, S. S. Jaswal, and E. Y. Tsymbal, “Intrinsic Defects in

Multiferroic BiFeO3 and Their Effect on Magnetism,” Phys. Rev. B, 85 [10]104409, 8pp (2012).

96S. J. Clark and J. Robertson, “Energy Levels of Oxygen Vacancies in

BiFeO3 by Screened Exchange,” Appl. Phys. Lett., 94 [2] 022902, 3pp (2009).97

M. V. Raymond and D. M. Smyth, “Defects and Charge Transport inPerovskite Ferroelectrics,” J. Phys. Chem. Solids, 57 [10] 1507 – 11 (1996).

98I. Wærnhus, T. Grande, and K. Wiik, “Surface Exchange of Oxygen inLa1xSrxFeO3d (x = 0, 0.1),” Top. Catal., 54 [13 – 15] 1009 – 15 (2011).

99A. S. Poghossian, H. V. Abovian, P. B. Avakian, S. H. Mkrtchian, and V.M. Haroutunian, “Bismuth Ferrites: New Materials for Semiconductor GasSensors,” Sens. Actuators B, 4 [3 – 4] 545 – 9 (1991).

100C. Wang, M. Takahashi, H. Fujino, X. Zhao, E. Kume, T. Horiuchi, and

S. Sakai, “Leakage Current of Multiferroic (Bi0.6 Tb0.3La0.1)FeO3 Thin FilmsGrown at Various Oxygen Pressures by Pulsed Laser Deposition and Anneal-ing Effect,” J. Appl. Phys., 99 [5] 054104, 5pp (2006).

101B. Vengalis, J. Devenson, A.K. Oginskis, R. Butkute, A. Maneikis, A.

Steikuniene, L. Dapkus, J. Banys, and M. Kinka, “Growth and Investigationof Heterostructures Based on Multiferroic BiFeO3,” Acta Phys. Pol. A, 113 [3]1095 – 8 (2008).

102T. H. Kim, B. C. Jeon, T. Min, S. M. Yang, D. Lee, Y. S. Kim, S.-H. Baek,

W. Saenrang, C.-B. Eom, T. K. Song, J.-G. Yoon, and T. W. Noh, “ContinuousControl of Charge Transport in Bi-Deficient BiFeO3 Films Through Local Fer-roelectric Switching,” Adv. Funct. Mater., 22 [23] 4962 – 8 (2012).

103N. Maso and A. R. West, “Electrical Properties of Ca-Doped BiFeO3

Ceramics: From p-Type Semiconduction to Oxide-Ion Conduction,” Chem.Mater., 24 [11] 2127 – 32 (2012).

104M. C. Li, J. Driscoll, L. H. Liu, and L. C. Zhao, “The Phase Transitionand Phase Stability of Magnetoelectric BiFeO3,” Mater. Sci. Eng. A, 4 38 – 440,346 – 9 (2006).

105J. F. Scott, “Ferroelectrics Go Bananas,” J. Phys. Condens. Mat., 20 [ 2]021001, 2pp (2008).

106D. Damjanovic, “Hysteresis in Piezoelectric and Ferroelectric Materials”;pp. 337 – 465 in The Science of Hysteresis, Vol. 3 . Edited by G. Bertotti and I.Mayergoyz. Academic Press, Oxford, U.K., 2006.

107S.T. Zhang, M.H. Lu, D.Wu, Y.F. Chen,andN. B.Ming, “LargerPolariza-tion and Weak Ferromagnetism in Quenched BiFeO3 Ceramics With a DistortedRhombohedral CrystalStructure,” Appl.Phys.Lett., 87 [26]262907,3pp(2005).

108S. T. Zhang, L. H. Pang, Y. Zhang, M. H. Lu, and Y. F. Chen, “Prepa-ration, Structures, and Multiferroic Properties of Single Phase Bi1xLaxFeO3

(x = 0 – 0.40) Ceramics,” J. Appl. Phys., 100 [11] 114108, 6pp (2006).109

W. N. Su, D. H. Wang, Q. Q. Cao, Z. D. Han, J. Yin, J. R. Zhang, andY. W. Du, “Large Polarization and Enhanced Magnetic Properties in BiFeO 3

Ceramic Prepared by High-Pressure Synthesis,” Appl. Phys. Lett., 91 [9]092905, 3pp (2007).

110G. L. Yuan, Y. Yang, and S. W. Or, “Aging-Induced Double Ferroelec-tric Hysteresis Loops in BiFeO3 Multiferroic Ceramic,” Appl. Phys. Lett., 91[12] 122907, 3pp (2007).

111S. Zhang, L. Wang, Y. Chen, D. Wang, Y. Yao, and Y. Ma, “Observa-tion of Room Temperature Saturated Ferroelectric Polarization in Dy Substi-

tuted BiFeO3 Ceramics,” J. Appl. Phys., 111 [7] 074105, 5pp (2012).112A. Y. Kim, Y. J. Lee, J. S. Kim, S. H. Han, H. W. Kang, H. G. Lee,and C. I. Cheon, “Ferroelectric Properties of BiFeO 3 Ceramics Sintered UnderLow Oxygen Partial Pressure,” J. Korean Phys. Soc., 60 [1] 83 – 7 (2012).

113K. C ar l a nd K. H. H€ardtl, “Electrical After-Effects in Pb(Ti,Zr)O3

Ceramics,” Ferroelectrics, 17 [1] 473 – 86 (1977).114M. I. Morozov and D. Damjanovic, “Hardening-Softening Transition in

Fe-Doped Pb(Zr,Ti)O3 Ceramics and Evolution of the Third Harmonic of thePolarization Response,” J. Appl. Phys., 104 [3] 034107, 8pp (2008).

115P. V. Lambeck and G. H. Jonker, “The Nature of Domain Stabilization

in Ferroelectric Perovskites,” J. Phys. Chem. Solids, 47 [5] 453 – 61 (1986).116L. Zhang and X. Ren, “Aging Behavior in Single-Domain Mn-Doped Ba-

TiO3 Crystals: Implication for a Unified Microscopic Explanation of Ferro-electric Aging,” Phys. Rev. B, 73 [9] 094121, 6pp (2006).

117X. Ren, “Large Electric-Field-Induced Strain in Ferroelectric Crystals by

Point-Defect-Mediated Reversible Domain Switching,” Nat. Mater., 3 [2] 91 – 4(2004).

118H. Birol, D. Damjanovic, and N. Setter, “Preparation and Characteriza-tion of (K0.5Na0.5)NbO3 Ceramics,” J. Eur. Ceram. Soc., 26 [6] 861 – 6 (2006).

119R. A. Eichel, “Structural and Dynamic Properties of Oxygen Vacancies inPerovskite Oxides – Analysis of Defect Chemistry by Modern Multi-Frequencyand Pulsed EPR Techniques,” Phys. Chem. Chem. Phys., 13 [2] 368 – 84 (2011).

120U. Robels and G. Arlt, “Domain Wall Clamping in Ferroelectrics by Ori-entation of Defects,” J. Appl. Phys., 73 [7] 3454 – 60 (1993).

121W. L. Warren, D. Dimos, G. E. Pike, K. Vanheusden, and R. Ramesh,“Alignment of Defect Dipoles in Polycrystalline Ferroelectrics,” Appl. Phys.Lett., 67 [12] 1689 – 91 (1995).

122W. L. Warren, G. E. Pike, K. Vanheusden, D. Dimos, B. A. Tuttle, andJ. Robertson, “Defect-Dipole Alignment and Tetragonal Strain in Ferroelec-trics,” J. Appl. Phys., 79 [12] 9250 – 7 (1996).

123W. L. Warren, K. Vanheusden, D. Dimos, G. E. Pike, and B. A. Tuttle,

“Oxygen Vacancy Motion in Perovskite Oxides,” J. Am. Ceram. Soc., 79 [2]536 – 8 (1996).

124P. Erhart, P. Tr€askelin, and K. Albe, “Formation and Switching of

Defect Dipoles in Acceptor-Doped Lead Titanate: A Kinetic Model Based onFirst-Principles Calculations,” Phys. Rev. B, 88 [2] 024107, 10pp (2013).

125A. Chandrasekaran, D. Damjanovic, N. Setter, and N. Marzari, “Defect

Ordering and Defect – Domain-Wall Interactions in PbTiO3

: A First-PrinciplesStudy,” Phys. Rev. B, 88 [21] 214116, 7pp (2013).

126Z. Wen, G. Hu, C. Yang, and W. Wu, “Effect of Annealing Process on

Asymmetric Coercitivities of Mn-Doped BiFeO3 Thin Films,” Appl. Phys. A,97, 937 – 41 (2009).

127Z. Feng and X. Ren, “Striking Similarity of Ferroelectric Aging Effect in

Tetragonal, Orthorhombic and Rhombohedral Crystal Structures,” Phys. Rev.B, 77 [13] 134115, 6pp (2008).

128L. Zhang, E. Erdem, X. Ren, and R. A. Eichel, “Reorientation of (MnTi’’V o

●●)9 Defect Dipoles in Acceptor-Modified BaTiO3 Single Crystals:An Electron Paramagnetic Resonance Study,” Appl. Phys. Lett., 93 [20]202901, 3pp (2008).

129A. Kontsos and C. M. Landis, “Computational Modeling of Domain

Wall Interactions With Dislocations in Ferroelectric Crystals,” Int. J. SolidsStruct., 46 [6] 1491 – 8 (2009).

130A. Lubk, M. D. Rossell, J. Seidel, Y. H. Chu, R. Ramesh, M. J. H€ytch,and E. Snoeck, “Electromechanical Coupling Among Edge Dislocations,Domain Walls, and Nanodomains in BiFeO3 Revealed by Unit-Cell-WiseStrain and Polarization Maps,” Nano Lett., 13 [4] 1410 – 5 (2013).

131J. Seidel, L. W. Martin, Q. He, Q. Zhan, Y.-H. Chu, A. Rother, M. E.Hawkridge, P. Maksymovych, P. Yu, M. Gajek, N. Balke, S. V. Kalinin, S.Gemming, F. Wang, G. Catalan, J. F. Scott, N. A. Spaldin, J. Orenstein, andR. Ramesh, “Conduction at Domain Walls in Oxide Multiferroics,” Nat.Mater., 8 [3] 229 – 34 (2009).

132J. Seidel, P. Maksymovych, Y. Batra, A. Katan, S.-Y. Yang, Q. He, A.P. Baddorf, S.V. Kalinin, C.-H. Yang, J.-C. Yang, Y.-H. Chu, E. K. H. Salje,H. Wormeester, M. Salmeron, and R. Ramesh, “Domain Wall Conductivityin La-Doped BiFeO3,” Phys. Rev. Lett., 105 [19] 197603 (2010).

133S. Farokhipoor and B. Noheda, “Conduction Through 71° Domain Walls

in BiFeO3 Thin Films,” Phys. Rev. Lett., 107 [12] 127601 (2011).134

P. Maksymovych, J. Seidel, Y. H. Chu, P. Wu, A. P. Baddorf, L. Q.Chen, S. V. Kalinin, and R. Ramesh, “Dynamic Conductivity of FerroelectricDomain Walls in BiFeO3,” Nano Lett., 11 [5] 1906 – 12 (2011).

135V. S. Postnikov, V. S. Pavlov, and S. K. Turkov, “Internal Friction in

Ferroelectrics Due to Interaction of Domain Boundaries and Point Defects,”J. Phys. Chem. Solids, 31 [8] 1785 – 91 (1970).

136V. S. Postnikov, V. S. Pavlov, S. A. Gridnev, and S. K. Turkov, “Interac-

tion Between 90o Domain Walls and Point Defects of the Crystal Lattice inFerroelectric Ceramics,” Sov. Phys. – Solid State, 10 [6] 1267 – 70 (1968).

137Z. Li, H. Wu, and W. Cao, “Piezoelectric Response of Charged Non-

180° Domain Walls in Ferroelectric Ceramics,” J. Appl. Phys., 11 1 [2] 024106,5pp (2012).

138T. Rojac, A. Bencan, G. Drazic, M. Kosec, and D. Damjanovic, “Piezo-electric Nonlinearity and Frequency Dispersion of the Direct PiezoelectricResponse of BiFeO3 Ceramics,” J. Appl. Phys., 112 [6] 064114, 12pp (2012).

139Y. F. Gong, P. Wu, W. F. Liu, S. Y. Wang, G. Y. Liu, and G. H. Rao,“Switchable Ferroelectric Diode Effect and Piezoelectric Properties of Bi0.9La0.1FeO3 Ceramics,” Chin. Phys. Lett., 29 [4] 047701, 4pp (2012).

140D. Damjanovic, “Ferroelectric, Dielectric and Piezoelectric Properties of Ferroelectric Thin Films and Ceramics,” Rep. Prog. Phys., 61 [9] 1267 – 324(1998).

141K. Uchino, Ferroelectric Devices. Marcel Deker, New York, USA, 2000.

142

T. Tsurumi, Y. Kumano, N. Ohashi, T. Takenaka, and O. Fukunaga,“90o Domain Reorientation and Electric-Field-Induced Strain of TetragonalLead Zirconate Titanate Ceramics,” Jpn. J. Appl. Phys., 36 [1] 5970 – 5(1997).

143R. J. Zeches, M. D. Rossell, J. X. Zhang, A. J. Hatt, Q. He, C.-H. Yang,

A. Kumar, C. H. Wang, A. Melville, C. Adamo, G. Sheng, Y.-H. Chu, J. F.Ihlefeld, R. Erni, C. Ederer, V. Gopalan, L. Q. Chen, D. G. Schlom, N. A.Spaldin, L. W. Martin, and R. Ramesh, “A Strain-Driven MorphotropicPhase Boundary in BiFeO3,” Science, 326 [5955] 977 – 80 (2009).

144J. X. Zhang, B. Xiang, Q. He, J. Seidel, R. J. Zeches, P. Yu, S. Y. Yang,

C.H. Wang, Y. H. Chu, L. W. Martin, A. M. Minor, and R. Ramesh, “LargeField-Induced Strains in a Lead-Free Piezoelectric Material,” Nat. Nanotech-nol., 6 [2] 98 – 102 (2011).

145R. K. Vasudevan, M. B. Okatan, Y. Y. Liu, S. Jesse, J.-C. Yang, W.-I.

Liang, Y.-H. Chu, J. Y. Li, S. V. Kalinin, and V. Nagarajan, “Unraveling theOrigins of Electromechanical Response in Mixed-Phase Bismuth Ferrite,”Phys. Rev. B, 88 [2] 020402, 7pp (2013).

146J. E. Daniels, A. Pramanick, and J. L. Jones, “Time-Resolved Character-ization of Ferroelectrics Using High-Energy X-Ray Diffraction,” IEEE Trans.

UFFC , 56 [8] 1539 –

45 (2009).147

A. Pramanick, J. E. Daniels, and J. L. Jones, “Subcoercive Cyclic Electri-cal Loading of Lead Zirconate Titanate Ceramics II: Time-Resolved X-RayDiffraction,” J. Am. Ceram. Soc., 92 [10] 2300 – 10 (2009).

148J. L. Jones, E. B. Slamovich, and K. J. Bowman, “Domain Texture

Distributions in Tetragonal Lead Zirconate Titanate by X-Ray and NeutronDiffraction,” J. Appl. Phys., 97 [3] 034113, 6pp (2005).

149A. Pramanick, D. Damjanovic, J. E. Daniels, J. C. Nino, and J. L. Jones,

“Origins Of Electro-Mechanical Coupling in Polycrystalline FerroelectricsDuring Subcoercive Electrical Loading,” J. Am. Ceram. Soc., 94 [2] 293 – 309(2011).

150J. L. Jones, M. Hoffman, and K. J. Bowman, “Saturated Domain Switch-

ing Textures and Strains in Ferroelastic Ceramics,” J. Appl. Phys., 98 [2]024115, 6pp (2005).

151Q. H. Jiang, C. W. Nan, and Z. J. Shen, “Synthesis and Properties of

Multiferroic La-Modified BiFeO3 Ceramics,” J. Am. Ceram. Soc., 89 [7] 2123 – 7 (2006).

152V. V. Shvartsman, W. Kleemann, R. Haumont, and J. Kreisel, “Large

Bulk Polarization and Regular Domain Structure in Ceramic BiFeO3

,” Appl.Phys. Lett., 90 [17] 172115, 3pp (2007).




153Y. Yao, B. Ploss, C. L. Mak, and K. H. Wong, “Pyroelectric Properties

of BiFeO3 Ceramics Prepared By a Modified Solid-State-Reaction Method,”Appl. Phys. A, 99 [1] 211 – 6 (2010).

154D. Damjanovic and M. Demartin, “Contribution of the Irreversible Dis-

placement of Domain Walls to the Piezoelectric Effect in Barium Titanate andLead Zirconate Titanate Ceramics,” J. Phys. Condens. Mat., 9 [23] 4943 – 53(1997).

155S. Li, W. Cao, and L. E. Cross, “The Extrinsic Nature of NonlinearBehavior Observed in Lead Zirconate Titanate Ferroelectric Ceramic,” J.Appl. Phys., 69 [10] 7219 – 24 (1991).

156D. Damjanovic, “Stress and Frequency Dependence of the Direct Piezo-

electric Effect in Ferroelectric Ceramics,” J. Appl. Phys., 82 [4] 1788 – 97 (1997).

157R. E. Eitel, T. R. Shrout, and C. A. Randall, “Nonlinear Contributions

to the Dielectric Permittivity and Converse Piezoelectric Coefficient in Piezo-electric Ceramics,” J. Appl. Phys., 99 [12] 124110, 7pp (2006).

158J. E. Garcıa, R. Perez, D. A. Ochoa, A. Albareda, M. H. Lente, and J.

A. Eiras, “Evaluation of Domain Wall Motion in Lead Zirconate TitanateCeramics by Nonlinear Response Measurements,” J. Appl. Phys., 103 [5]054108, 8pp (2008).

159D. Damjanovic, M. Demartin Maeder, P. Duran Martin, C. Voisard, andN. Setter, “Maxwell – Wagner Piezoelectric Relaxation in Ferroelectric Hetero-structures,” J. Appl. Phys., 90 [11] 5708 – 12 (2001). h

Tadej Rojac received his B.Sc. in2003 in Chemical Engineering anda Ph.D. degree in 2007 in thesame field from the University of Ljubljana, Slovenia. In 2009 heperformed a one-year postdoctoralstudy at the Swiss Federal Insti-tute of Technology in Lausanne,Switzerland. Since 2000 he hasbeen working as a researcher atthe Electronic Ceramics Depart-ment of the Jozef Stefan Institutein Ljubljana, Slovenia. His main

research interests cover mechanochemical reaction mecha-nisms, application of the mechanochemical processing in thesynthesis of complex ceramic oxides, and processing-struc-ture-properties relationship in lead-based and lead-free piezo-electric ceramics and thick films. From 2013 he is an Asst.Prof. at the Jozef Stefan International Postgraduate School.He is author and co-author of 25 scientific papers, 3 reviewpapers, 3 chapters in monographs and 1 patent.

Andreja Bencan obtained her B.Sc.in 1998 in Chemical technologyand in 2002 a Ph.D. in the field of Materials, from the Faculty of Chemistry and Chemical Technol-ogy, University of Ljubljana,Slovenia. She is a seniorresearcher at the ElectronicCeramics Department of the Jozef Stefan Institute, Ljubljana. Hermain research interest lies in thedevelopment of lead-based andlead-free piezoceramics, especially

in the microstructure investigations by different scanning andtransmission electron microscopy methods. From 2009 she isalso habilitated as an Asst. Prof. at the Jozef Stefan Interna-tional Postgraduate School. She is co-author of about 60 sci-entific papers and 4 chapters in monographs.

Barbara Malic received her Ph.D.degree in 1995 in the field of chemistry at the University of Ljubljana, Slovenia. She is cur-rently the head of the ElectronicCeramics Department at the Jozef Stefan Institute, Ljubljana, Slove-nia, and an associate professor of chemistry of materials at the Jozef Stefan International PostgraduateSchool. Her research topics coverlead-based and lead-free ferroelec-tric and piezoelectric ceramics and

thin films, tunable ferroelectric thin films and solution-derived materials for transparent electronics. She is alsoinvolved in studies related to chemical solution deposition of thin films, inkjet printing of solutions and particle disper-sions. She is author or co-author of more than 130 papers,10 book-chapters, more than 110 technical reports and 4 Slo-venian patents.

Goknur Tutuncu is currently aPostdoctoral Research Associateat Brookhaven National Labora-

tory (BNL) and before joiningBNL she spent 2 years as a Post-doctoral researcher at the Univer-sity of Florida. She received herM.S. degree in Chemical Engineer-ing in 2002 from the Middle EastTechnical University and herPh.D. degree in Materials Scienceand Engineering in 2010 fromIowa State University. Her pri-

mary research focuses and expertises are structure/propertycorrelations in piezoelectric materials, processing and charac-terizations of lead-free ferroelectric materials utilizing sophis-ticated X-ray and neutron diffraction techniques.

Jacob Jones is an Associate Pro-fessor of Materials Science andEngineering at North CarolinaState University and Director of the Analytical InstrumentationFacility (AIF). Jones received hisPhD from Purdue University inMaterials Engineering in 2004.From 2006-2013, Jones was at theUniversity of Florida where hewas an Assistant and then Associ-ate Professor of Materials Scienceand Engineering. Jones is an

experimental materials scientist with research interests in X-

ray and neutron scattering, crystallography, ceramic materi-als, and mechanical behavior of materials. He has publishedover 100 publications on these topics since 2004. Jones hasreceived numerous research awards including the NationalScience Foundation CAREER award, a Presidential EarlyCareer Award for Scientists and Engineers, the IEEE Ferro-electrics Young Investigator Award, a National NuclearSecurity Administration Defense Program Award of Excellence, the HHMI Distinguished Mentor (DM) Awardfor undergraduate research mentoring, an InternationalEducator award, and twice received the Edward C. HenryAward by the Electronics Division of the American CeramicSociety.




John Daniels is currently in theSchool of Materials Science andEngineering at UNSW as a SeniorLecturer and Australian Instituteof Nuclear Science and Engineer-ing research fellow. He was awar-ded his PhD in 2007 from theSchool of Physics at Monash Uni-versity, Melbourne, for work inthe field of time-resolved neutron

scattering in ferroelectric materi-als. After his PhD, he spent threeyears as a postdoctoral researcher

within the Structure of Materials group at the EuropeanSynchrotron Radiation Facility, Grenoble, France. His cur-rent research interests are in the application of advanced x-ray and neutron scattering techniques to the study of func-tional and mechanical properties of electro-ceramic materials.

Dragan Damjanovic received BScdiploma in Physics from the Uni-versity of Sarajevo in 1980, andPhD in Ceramics Science from thePennsylvania State University(PSU) in 1987. From 1988 to 1991he was a research associate in theMaterials Research Laboratory atthe PSU. He joined the CeramicsLaboratory, Institute of Materials,at the Swiss Federal Institute of Technology in Lausanne (EPFL)in 1991. He is currently a profes-

seur titulaire and teaches undergraduate and graduate cour-ses on electrical properties of materials. He investigatesexperimentally physical processes taking place at differentdriving fields and time scales and how they affect macro-

scopic behavior of ceramics, single crystals and thin layers.His interests include interaction of defects with domain walls,symmetry breaking and its effect on electro-mechanical andelectro-thermal coupling, interface dynamics, dispersion,creep, nonlinearity and hysteresis in dielectric, mechanicaland piezoelectric responses, phase transition-related instabili-ties, structure/microstructure – property relations, and appli-cations of dielectric, piezoelectric and ferroelectric crystals,films, and ceramics. He is an IEEE Fellow, was awarded2007 Outstanding Achievement Award by the International

Symposium on Integrated Ferroelectrics and 2009 Ferroelec-trics Recognition Award by the IEEE Ultrasonics, Ferroelec-trics and Frequency Control (UFFC) Society, wasDistinguished Lecturer for the IEEE UFFC Society for2010/11, and won with his colleagues the 2012 Edward C.Henry Best Paper Award by the Electronics Division of theAmerican Ceramic Society. He authored and co-authoredmore than 190 papers.


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