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Materials Chemistry and Physics 130 (2011) 381– 386

Contents lists available at ScienceDirect

Materials Chemistry and Physics

j ourna l ho me pag e: www.elsev ier .com/ locate /matchemphys

lectrical behavior of Pb(Zr0.52Ti0.48)0.5(Fe0.5Nb0.5)0.5O3 ceramics

awnit Kumar ∗, Avijit Ghosh, R.N.P. Choudharyepartment of Physics and Meteorology, Indian Institute of Technology Kharagpur, 721 302, India

r t i c l e i n f o

rticle history:eceived 31 July 2010eceived in revised form 22 April 2011ccepted 28 June 2011

eywords:iffuse phase transition

a b s t r a c t

Polycrystalline sample of Pb(Zr0.52Ti0.48)0.5(Fe0.5Nb0.5)0.5O3 (PZTFN) ceramics was prepared by a high-temperature solid-state reaction technique. Preliminary X-ray diffraction analysis confirms the formationof single phase compound with tetragonal crystal structure. The dielectric behavior of the materialwas studied in a wide frequency (100 Hz–1 MHz) and temperature (30–500 ◦C) range. The relation-ship between dielectric constant εr and temperature showed a single Curie transition temperature(Tc = 255 ◦C), with as large a εr value as 4200 at Tc. A low value of dissipation factor (tan ı), as small as 0.04

xygen vacancyonductivityysteresis

up to temperature above than Tc (frequency f = 1 kHz), with good insulating properties made the samplefor prospective applications in microwave dielectric components and high-power-density capacitors. Thehigher diffusivity (� ∼ 1.6) value of the dielectric peak in the material indicates the greater disorderingin the system. The signature of ac conductivity depending upon frequency and temperature exhibits thepresence of different type of charge species. It is observed that oxygen vacancy plays an important rolefor high temperature conduction mechanism and dynamics of polarization of the system.

. Introduction

Recently multiferroics in bulk or thin film form are explored forighly-sensitive ac magnetic field sensors and electrically tunableicrowave devices such as filters, oscillators and phase shifters inhich the ferri-, ferro- or antiferro-magnetic resonance is tuned

lectrically instead of magnetically [1]. As compared to singlehase compounds, composites yield larger magnetoelectric cou-ling coefficients. But recently, Kumar et al. [2,3,4] has made seriesf ambient temperature complex single phase multiferroics withood magnetoelectric coupling. It is well known that lead zirconateitanate, Pb(Zr0.52Ti0.48)O3 (PZT), near the morphotropic phaseoundary (MPB) exhibits an excellent piezoelectric property dueo a high electromechanical coupling factor resulting from the co-xistence of the rhombohedral and tetragonal phases. PZT ceramicsormally have (a) high remnant polarization value at room tem-erature and (b) very high Curie temperature (Tc) in the vicinity of00 ◦C [5]. In addition, lead iron niobate Pb(Fe0.5Nb0.5)O3 (PFN) isn attractive material for use in multilayer ceramic capacitors andther electronic devices due to its high dielectric constant, diffusehase transition behavior and low sintering temperature [6]. PFNlso exhibits the magnetoelectric behavior undergoing a ferroelec-

ric to paraelectric and a paramagnetic to antiferromagnetic phaseransition at 112 ◦C [7] and −130 ◦C [8,9], respectively. Recently aeak ferromagnetism at room-temperature was observed in bulk

∗ Corresponding author. Tel.: +91 9933363461.E-mail address: nawnitphy@gmail.com (N. Kumar).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.06.059

© 2011 Elsevier B.V. All rights reserved.

as well as thin film form of PFN due to its spin canting [10,11]. Itis observed that PZT does not have magnetoelectric (ME) proper-ties like wise PFN, and again PFN itself does not have high value ofelectromechanical coupling factor and remanat polarization valuelike wise PZT. In addition, PFN exhibits large dielectric loss aroundroom temperature, and the dielectric loss increases rapidly with theincrease of temperature, which makes it unsuitable for the appli-cation in the multilayer ceramics capacitor industry [12]. Keepingall those behavior in our mind we have prepared PZTFN ceramics,which will combine several properties in the same phase with opti-mum value at ambient transition temperature. Many applicationsrequire Tc to be closed to ambient temperature. Therefore, there isa general interest to reduce the Tc of PZT ceramics to optimize theiruses. In this paper we have reported only structural and electricalproperties of the solid solution of PZT-PFN ceramics. It is observedthat the presence of defect and vacancies are detrimental factor tothe performance of ferroelectric behavior.

2. Experimental details

Polycrystalline sample of Pb(Zr0.52Ti0.48)0.5(Fe0.5Nb0.5)0.5O3 was synthesizedusing high-purity (≥99.9%) oxides: PbO, Nb2O5, TiO2, ZrO2 and Fe2O3 (M S−1 LobaChemicals) by a high-temperature solid-state reaction method. The powders weretaken in a stoichiometric amount (with 3% excess PbO to compensate the lead loss atelevated temperatures) and mixed them thoroughly in an agate-mortar in a wet con-dition (methanol) for 3 h and calcined at 1000 ◦C in a high-purity alumina crucible inair for 8 h. The crystalline structure of the calcined powder was examined by an X-

ray diffraction (XRD) technique using an X-ray powder diffractometer (using PHILIPS(PW3373 XPERT-PRO diffractometer)). Room temperature high resolution XRD pat-tern of the sample was obtained in a wide range of Bragg angles 2� (20◦ ≤ 2� ≤ 60◦)with Cu K� (� = 1.54 A) radiation. Finally, the calcined powder was used to makecylindrical pellets of diameter 10 mm and thickness 1–2 mm under a hydro-static

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mal Curie Weiss type dielectric) and 2 (for typical diffuse typedielectric). After performing linear fit, the average value of thediffuseness exponent (�) was observed as 1.6 (Fig. 4). The valueof � suggests that the compound does not follow a Curie Weiss

ig. 1. Room temperature XRD pattern of the sample, SEM micrograph is shown innset.

ressure of about 6 × 106 N m−2 using a hydraulic press. Polyvinyl alcohol (PVA) wassed as binder to reduce the brittleness of the pellets. The binder was burnt out dur-

ng the sintering process. The pellets were sintered in a covered alumina cruciblet 1050 ◦C in air atmosphere. The grain morphology and size were directly imagedsing scanning electron microscopy (ZEISS–SEM) and the average grain size wasetermined by using a mean linear intercept method.

For dielectric and electrical characterization, high purity silver-paste was thenainted on the flat polished surfaces of the sintered pellets to act as an electrodend then dried at 150 ◦C in an oven for 4 h. The electrical parameters were obtaineds a function of frequency (100 Hz–1 MHz) using an impedance analyzer (N4L PSM,735) in conjunction with a laboratory made sample holder in a wide tempera-ure range from 30 to 500 ◦C. The hysteresis loop of the poled sample was obtainedsing workstation of loop tracer (M/S Radiant Technology Inc., USA). Piezoelectricoefficient d33 is measured using Piezo-Meter (PM35 Take-control, UK) at 110 Hzpressure 1 N m−2) at room temperature of the poled sample. For the piezoelectricnd P–E hysteresis loop measurements, the sample was poled at 110 ◦C temperaturen a silicon oil bath under the dc field of 20 kV cm−1 for 1 h using APLAB high voltagec power supply (model 7342P).

. Results and discussion

.1. Structural property and microstructure

Fig. 1 exhibits the room temperature X-ray diffraction (XRD)attern of the sample. All the XRD peaks of the system were indexedsing a standard computer program POWDMULT [13]. On theasis of best agreement between observed and calculated valuesf d (inter-planar spacing) and (i.e. ��d = dobs − dcal = minimum),t is found that the structure is tetragonal (i.e. responsible forerroelectric property) with least-squares refined lattice parame-ers, a = 4.005(12) A, c = 4.0711(11) A. The experimental density ofhe sample is found as 97% of the theoretical density. The aver-ge crystallite size of the sample was calculated from the X-rayine broadening using Scherer’s equation [14]; D = K�/ˇ1/2 cos�hkl,

here K (constant) is 0.89, ˇ1/2 is the full width at half maximaFWHM) (in radian) and �hkl is the Bragg angle. The value of averagerystallite size is found to be 25 nm (the effects of strain, instrumen-al factor and other defects have been ignored in the calculations).

The prepared perovskite oxides are normally satisfying the fol-owing conditions: (a) charge neutrality and (b) tolerance factort). A simple description of the geometric packing within the per-

vskite structure can be characterized by tolerance factor (t). Its defined as t = (rA + rO)/

√2 (rB + rO) where rA is the average ionic

adius of the A site atoms; rB is the average ionic radius of B-sitetoms and rO is the ionic radius of O2− = 1.40 A. For ideal perovskite,

nd Physics 130 (2011) 381– 386

the value of t should be 1 and for distorted perovskites, this valueshould be less than 1 [15]. In the Pb(Zr0.52Ti0.48)0.5(Fe0.5Nb0.5)0.5O3system, the A-site is occupied by Pb2+ (1.19 A) ions, and the Fe3+,Nb5+, Ti4+ and Zr4+ ions occupy the B-site of the ABO3 perovskitecrystal structure. The average ionic radius of the B-site ions inthe system can be calculated from the following equation rb =0.26rZr4+ + 0.25rFe3+ + 0.25rNb5+ + 0.24rTi4+ where the ionic radiiof Fe3+, Nb5+, Ti4+ and Zr4+ are 0.55, 0.64, 0.61 and 0.84 A, respec-tively. The calculated value of the tolerance factor t = 0.89 suggeststhat the system has distorted peroveskite structure [16]. Fig. 1(inset) exhibits SEM micrograph of plane surface with good densi-fication and homogenous grain size distribution of polycrystallinesample. The grain size varies from 1.5 to 3.0 �m and some degreeof porosity is clearly seen.

3.2. Dielectric properties

Fig. 2 represents the dispersion curves of the dielectric constant(εr) and the loss tangent (tan ı), respectively. It is found that εr andtan ı decrease with increase in frequency. This is a typical charac-teristic of a dielectric or ferroelectric material. The high value ofεr and tan ı at lower frequency indicates the presence of all typeof polarization (i.e. electronic, ionic, orientation, and space chargepolarization). From the nature of variation of εr with frequency, itis possible to find out the different contributions of polarizationprecisely in a particular frequency range of the system.

Fig. 3 shows the temperature dependent dielectric constantof the sample. The ferroelectric to paraelectric phase-transitiontemperature (Tc) is observed at about 255 ◦C with a high dielec-tric constant value as 4200 at Tc for 1 kHz. It is observed thatthere is no shift of maximum dielectric constant temperature Tc

(εmax) as function of frequency, which leads to the behavior of thesystem as nonrelaxor. This behavior indicates that the present sys-tem is nonrelaxor in contrast to the complex perovskites such asPb(Mg1/3Nb2/3)O3 and Pb(Fe2/3W1/3)O3 ceramics [17,18]. The dif-fuseness of peak was analyzed using an empirical formula [19]:1/εr − 1/εr max = C(T − Tc)� or, ln(1/εr − 1/εr max) = lnC + �ln(T − Tc)where εr is the maximum value of dielectric constant at temper-ature T. The value of the exponent � gives the value of degreeof diffuseness. The value of � is measured from the slope ofln(1/εr − 1/εr max) versus ln(T − Tc) which is in between 1 (for nor-

Fig. 2. Variation of relative permittivity and the loss with frequency at room tem-perature.

N. Kumar et al. / Materials Chemistry and Physics 130 (2011) 381– 386 383

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Fig. 5. Variation of the loss with temperature at selective frequencies and the lossvariation at 1 MHz with temperature are shown in the inset.

ig. 3. Variation of relative permittivity with temperature at selective frequencies.

ype of phase transition, and hence system has diffuse type phaseransition. The diffuse phase transition indicates the existence ofelaxation process in the material, and this broad dielectric relax-tion cannot be described by a single relaxation. This behavioran be associated with compositional fluctuation which leads tohe formation of polar micro or nano-regions with different Curieemperature [20]. In present system the B site cation disorderingre most favorable. The B-site cation disordering affects the peakroadening [21] because A B−1 site cation disordering promoteshe short range ordering, this is responsible for the formation of

icro or nano-polar regions [22,23]. The oxygen vacancy also playsn important role to break the long range ordering which is ulti-ately responsible for diffuse phase transition. It is also observed

hat oxygen vacancy plays a dominant role in peak broadening (forhe development of ferroic domains) rather than B site cation dis-rdering [24]. The observed diffuse phase transition may be theesults of B-site cation disordering or oxygen vacancy present inhe system or due to both. One interesting feature (peak around

62 ◦C for 1 kHz) is also observed in the high temperature region.his peak is shifted towards higher temperatures with increasingrequency and fully disappears at higher frequencies. This typef diffused dielectric anomaly in the high-temperature region is

Fig. 4. Representative plot of log(1/�r − 1/�r max) vs. log(T − Tc) at 10 kHz.

normally investigated in perovskite ferroelectric ceramics or sin-gle crystals. Several experimental evidences have proved that thephase transition (first peak) is not related to the diffused dielectricanomaly (second peak) present at high temperatures. Therefore,this is not an intrinsic property of the system, but it is related tothe oxygen vacancy conductivity relaxation process. The disappear-ance of dielectric anomaly at higher frequencies provides that thisphenomenon is predominant only for low-frequency relaxationprocess. Thus, this anomaly can be correlated to a low-frequencyrelaxation process due to the oxygen vacancies [25].

Fig. 5 exhibits that the relationship between temperature andtan ı. The value of tan ı as small as 0.04 from room tempera-ture to above than Tc (frequency f = 1 kHz), with good insulatingproperties and then increases on increasing temperature due tothe enhancement of conductivity in the material. A low value oftan ı and high dielectric constant value at Tc made the samplefor prospective applications in microwave dielectric componentsand high-power-density capacitors. A peak or anomaly near thetransition temperature is observed in the dielectric loss graph forfrequency 1 MHz (Fig. 5(inset)). A peak near transition temperatureat high frequency (1 MHz) from the dielectric loss is representa-tive of strong absorption. This strong absorption is attributed torelaxation processes associated with domain reorientation, domainwall motion, and the dipolar behavior of ferroelectric materials.This feature is observed in many Pb-based, Fe-containing complexperovskite compounds [20]. At lower frequency no dielectric losspeak is observed due to the intrinsic conduction by space chargeand other charge carriers which mask the ferroelectric contribu-tion, which is discussed in conductivity analysis. The value of tan ıis almost constant up to 350 ◦C, and then shows a sharp rise upto500 ◦C. The increase of tan ı with the decrease of frequency andthe increase of temperature have been reported in bulk [26] andthin film [27] ceramics. The dielectric loss has been interpretedin terms of the polarization of mobile charge carriers, most prob-ably oxygen vacancies which are activated at high temperature.As the polarization is induced by oxygen vacancies increased withthe increase of temperature and the decrease of frequency due tothe inertia of oxygen vacancies [26]. Therefore, significant dielec-tric loss at the reduced frequency and at high temperature maybe attributed to the enhanced space charge relaxation due to the

increase of oxygen-vacancy concentration in the system.

3 istry and Physics 130 (2011) 381– 386

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.3. Conductivity analysis

To understand the effect of bulk conductivity on the dielectric,erroelectric properties and fatigue resistance, it is imperative totudy the mechanism of formation of different conducting speciesn the sample. There are three conduction mechanisms associated

ith the ferroelectric characteristics of the material.If system isure and stoichiometric state then it should be of highly resistiveature and with low dielectric losses. However, during the calcina-ion and sintering processes, oxygen vacancies and charge carrierselectrons) are generated. In perovskite ferroelectric materials oxy-en vacancies are considered as one of the mobile charge carriers28]. There is a probability of formation of singly ionized and/doublyonized oxygen vacancy with the release of single or two electronsespectively from the oxygen in the system. Using Kröeger–Vinkotation, the occurrence of oxygen vacancies can be representedy the following thermally activated reactions:

0 = 1/2O2 + V′0 + e−

0 = 1/2O2 + V′′0 + 2e−

here the liberated electrons can be captured by the Fe3+, Ti4+,nd Nb5+ ions [20,29]. During these creation processes, some ofhe released electrons are captured by Ti4+ to form Ti3+, Nb5+ toorm Nb4+ and Fe3+ to form Fe2+. The electronic hopping betweenTi3+–Ti4+), (Fe2+–Fe3+) and (Nb4+–Nb5+) state leads to over all elec-ronic conductivity, which is dependent on the temperature andrequency. There is a probability of Ti4+ ion may combine withingle ionized oxygen vacancy (Ti3+–V′

0), Nb5+ ion may combineith single ionized oxygen vacancy (Nb4+–V′

0) and Fe3+ may com-ine with single ionized oxygen vacancy (Fe2+–V′

0) to form dipoles.hese dipoles affect the dielectric relaxation in the system.Latticetrains originated during the increment of the temperature belowhe ferro-paraelectric transition temperature crystal attains polar

ode. The electrons interact with the polar modes of the crystalso form small polaron which is responsible for conduction pro-ess in ferro- to paraelectric transition region [30]. The carriers areransported by hopping process.The motion of oxygen vacancies isot localized in one unit cell but can extend to the whole system

eading to ionic conductivity. The ionic conductivity is an intrin-ic and typical conduction mechanism in ferroelectric materialst high temperatures [31]. Therefore, each one of the above men-ioned mechanisms and the electrical conductivity are thermallyctivated.

.3.1. ac conductivityThe study of conduction mechanism shows that bulk conduc-

ivity is modified by dielectric properties. It is well known thatonductivity in ferroelectrics appreciably affects the properties inhe sense that there will be a competition between the growth ofhe ferroelectric phase and the free charge carriers. Conductivitytudies provide the information about the effects of conductivityn the domain structure and its motion, as well as the nature ofharge carriers present in the system.

The ac conductivity �ac of the system is calculated using theelation �ac = ωε0εrtan ı (ω = angular frequency, and ε0 = dielectriconstant in free space) from the measured dielectric parame-ers. The variation of ac conductivity (�ac) with 103/T at differentrequencies is shown in Fig. 6. The �ac increases on increasingrequency and temperature. The nature of variation of �ac over aide temperature range supports the thermally activated transport

roperties of the material. Different slope changes appear in differ-nt regime of temperature, indicating multiple activation processith different energies. In the lower temperature region, there is

inear variation in conductivity with temperature but strong depen-

Fig. 6. Variation of ac conductivity (�ac) with inverse of absolute temperature(1000/T) at selective frequencies.

dence on frequency. As the temperature increases, the conductivityshows an increasing behavior and the observed peak at Tc (255 ◦C)corresponds to phase transitions in dielectric study. Although thepeak is relatively small compared with the peak observed in dielec-tric constant curve (Fig. 3). In the higher temperature region,conductivity varies exponentially with the temperature but veryweak with frequency dependence and ultimately all the curvesmerge.

The apparent increase in the conductivity aroundferroelectric–paraelectric transition temperature is attributedto the increased polarizability of the material [32]. As the fre-quency decreases, oxygen vacancies and other defects becomeactive and mask the growth of the ferroelectric phase and noanomalies appear in temperature dependent conductivity (onlya slope change is observed). At higher temperatures (above Tc

the conductivity data tend to fall onto a straight line), the spacecharges are released and recombined, and hence conductivities atvarious frequency merge [18]. This is a typical behavior of the dccomponent of conductivity [19].

The ac conductivity is very much related to activationenergy and frequency. The activation energy can be calcu-lated from slope of ln �ac vs 1000/T using empirical formula�ac = �oexp(−Ea/KBT), (Ea = activation energy, KB = Boltzmann con-stant, �o is pre-exponential factor). The value of activation energyat high temperature is higher (0.81 eV for T ≥ Tc). The lower valueof activation energy (0.21 eV) obtained at low temperatures in cor-respondence with the hopping charge mechanisms is associatedmainly with the creation of large number of space charge carrierin oxidation–reduction process as explained above. The value ofactivation energy (0.46 eV) in phase transition region indicates theconduction due to polaron creation in distorted lattice [12]. Theactivation energy for the oxygen vacancy to move in the octahe-dra of peroveskite structure is 1 eV [33]. The calculated value ofactivation energy (Ea = 0.81 eV) from conductivity plots suggestsconduction due to hopping of oxygen vacancy in high tempera-ture region. The increase in conductivity is solely responsible for anincreased loss (tan ı vs temperature plot) at higher temperaturesas shown in Fig. 5.

3.4. Hysteresis

Fig. 7 displays the polarization–electric field (P–E) hysteresisloop of the poled sample. The nature of hysteresis loop behav-ior indicates ferroelectric character and dynamic polarizability of

N. Kumar et al. / Materials Chemistry and Physics 130 (2011) 381– 386 385

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ig. 7. Room temperature P–E loops for the sample corresponding to full hysteresisor selected Emax values: (a) 5 kV cm−1, (b) 9 kV cm−1, (c) 13 kV cm−1, (d) 18 kV cm−1,e) 22 kV cm−1, (f) 27 kV cm−1, (g) 31 kV cm−1, and (h) 36 kV cm−1.

he sample. In ferroelectric materials, domain wall movement canesult in a complex nonlinear macroscopic response, displayingignificant hysteresis and slow aging [34]. Fig. 7 shows the electric-eld amplitude dependence of hysteresis loops. The hysteresis loopecomes saturated as the field amplitude is increased. If an exter-al electric field is applied then the domains are favorably orientedith respect to the field, and the domain wall motion grows at

he expense of other less favorably oriented domains. With theias electric field increasing, the P–E hysteresis loops grow largernd wider. The variation of maximum polarization (Pmax), rem-ant polarization (Pr), and coercive field (Ec) at room temperature

rom the material with the applied electric field are observed fromig. 8. The obtained values of Pmax, Pr and Ec from the system aren between the reported value of these parameters for PZT and PFN5,35].

The ferroelectric characteristic of the ceramics can be assessedith the hysteresis loop squareness (Rsq), Rsq = (Pr/Ps) + (P1.1Ec/Pr),

here P1.1Ec is the polarization at the field equal to 1.1Ec. Pr is the

emnant polarization at zero electric field and Ps is the saturatedolarization obtained at some finite field strength below the dielec-

ig. 8. Variation of Pmax, Pr, and Ec as functions of the maximum applied electriceld.

Fig. 9. ln–ln plot of the hysteresis loop area versus field amplitude.

tric breakdown. For the ideal square loop, Rsq is equal to 2.00. Thehysteresis loop is of a typical “square” form as a result of domainswitching in an applied field. This is a typical characteristic of aphase that contains long-range interaction between dipoles in theferroelectric micro-domain state [36]. The value of Rsq is found to be(Rsq = 1.05). This confirms that PZTFN is a ferroelectric which lacksa long range interaction between dipoles. The long range orderingis broken by B-site disordering or the presence of oxygen vacancydefects [22–24].

The value of Ec is very large compared to that of other ordinaryperovskite-type ferroelectric materials such as BaTiO3 [37,38]. TheXRD pattern shows that the sample has good crystallinity withoutany impurity phase. Therefore, it is speculated that the large Ec ismainly caused by two reasons: (a) one is that the Ec of ferroelectricmaterial including Pb is large and (b) other is that the increasingof the number of ferroelectric domain boundaries which restrictsreversing of polarization [39]. From the hysteresis loop three dif-ferent regions are observed: (i) a major saturated hysteresis loop atlarge field well above Ec, (ii) a minor hysteresis loop broadly satu-rated at relatively small fields above Ec, and (iii) a minor hysteresisloop with area close to zero below Ec. Fig. 9 shows the ln–ln plotof the hysteresis loop area versus field amplitude, where the areaof hysteresis loops increases with the increment of field amplitude.The value of loop area change discontinuously at Ec = 9.12 kV cm−1

and then increases gradually to 36 kV cm−1. This value saturateswith further applied voltage. Above Ec (10 kV cm−1), the scaling fitfor the loop areas, A ∝ Ea, with respect to the field strength alsoshows a power law behavior with the scaling exponent a = 1.5.

The values of Pmax, Pr, and Ec obtained from each full hysteresisloop, are plotted as functions of maximum electric field for sample(Fig. 8), respectively. In system, Pmax as well as Pr steadily increaseswith applied electric field illustrating a nonlinear behavior of thematerial. Significant differences are observed in the values of Pmax

and Pr for the sample. Ferroelectric ceramics can exhibit the fol-lowing type of polarization at room temperature [40]:

(1) Polarization associated with highly polarized system con-tributed by polaron.

(2) Polarization associated with movement of space charge, cation

and anion vacancy.

(3) Polarization associated with orientation of the switchable fer-roelectric domains.

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The two former components, corresponding to the nonrem-ant component, are the most important contributions to theotal polarization of the hysteresis loop; whereas, the switch-ble ferroelectric domains or remnant component contributesn a lesser degree to the loop. For higher electric fields theb(Zr0.52Ti0.48)0.5(Fe0.5Nb0.5)0.5O3 sample exhibits an increase inhe Ec values. This can be due to higher domain pinning at the grainoundary for the PZTFN sample.

The coercive fields (2Ec) for the sample with Ec(+) = 11.6 kV cm−1

nd Ec(−) = 9.12 kV cm−1. An offset (2.4 kV cm−1) is observed inhe coercive voltage. The shift of hysteresis loop in the samples due to the generation of an “internal bias” or dipole moment.he observation can be attributed to the presence of the com-lex defects, which provide the internal field-bias. Defect-induced

attice deformations can modify local fields and polarization pat-erns in ferroelectrics [41]. Generally, oxygen vacancy is trapped inerroelectric domain boundaries because these energetically favor-ble positions act as potential well [42,43]. Such vacancy-inducedntiphase polarizations can play an important role in domain pin-ing. The defect dipole model can easily explain the phenomenonf the asymmetric hysteresis loops in poled sample. Defect dipolesormed by the oxygen-vacancy–lattice site associates might beandomly orientated in the system. When poling the sample, theefect dipoles orient to the poling direction forming an internalias. The ferroelectric interaction between these defect dipoleshich are trapped near domain walls and the other dipoles in theomain makes the polarization switching and domain reorienta-ion processes, as well as the domain wall motions, more difficultn one field direction than the other direction [44]. This results insymmetric loops, and the differences between Ec(−) and Ec(+)ere observed in Fig. 7. The poled sample has been taken for

he measurement of piezoelectric constants. The obtained valuef piezoelectric constant d33 = 115 pC N−1 from the system whichs in between the reported value for PZT and PFN [5,45] is foundo be suggested that the material can be used for some peizoevices.

. Conclusions

The single phase system with tetragonal crystal structure isbtained by a high temperature solid state reaction method. Theystem exhibits a diffuse type phase transition at 255 ◦C. Theifferent conduction mechanisms at (a) lower temperature dueo (i) trapped electron, and (ii) polaron, and (b) high temper-ture because of oxygen vacancy are observed. The presencef defect (oxygen vacancy) in the system modifies the localeld in the crystal system which affects the dynamic polar-

zation of the system and produce asymmetric P–E hysteresisoop.

cknowledgement

One of the authors (Nawnit Kumar) would like to thank CSIR,ew Delhi, India for research fellowship.

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nd Physics 130 (2011) 381– 386

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