Domain wall motion and its contribution to the dielectric and piezoelectric properties of lead...

JOURNAL OF APPLIED PHYSICS VOLUME 89, NUMBER 2 15 JANUARY 2001

Domain wall motion and its contribution to the dielectric and piezoelectricproperties of lead zirconate titanate films

F. Xu,a) S. Trolier-McKinstry, W. Ren, and Baomin XuMaterials Research Laboratory, The Pennsylvania State University, University Park, Pennsylvania 16802

Z.-L. Xie and K. J. HemkerDepartment of Mechanical Engineering, The Johns Hopkins University, Baltimore, Maryland 21218

~Received 21 October 1999; accepted for publication 25 September 2000!

In this article, domain wall motion and the extrinsic contributions to the dielectric and piezoelectricresponses in sol–gel derived lead zirconate titanate~PZT! films with compositions near themorphotropic phase boundary were investigated. It was found that although the films had differentthicknesses, grain sizes, and preferred orientations, similar intrinsic dielectric constants wereobtained for all films between 0.5 and 3.4mm thick. It was estimated that about 25%–50% of thedielectric response at room temperature was from extrinsic sources. The extrinsic contribution to thedielectric constant of PZT films was mainly attributed to 180° domain wall motion, which increasedwith both film thickness and grain size. In studies on the direct and converse longitudinalpiezoelectric coefficients of PZT films as a function of either stress or electric driving field, it wasfound that the ferroelastic non-180° domain wall motion was limited. Thus extrinsic contributionsto the piezoelectric response were small in fine grain PZT films~especially those under 1.5mm inthickness!. However, as the films became thicker (.5mm!, nonlinear behavior between theconverse piezoelectric coefficient and the electric driving field was observed. This indicated thatthere was significant ferroelectric non-180° domain wall motion under high external excitation inthicker films. The activity of the non-180° domain walls was studied through non-180° domainswitching. For fine grain films with film thicknesses less than 2mm, non-180° switching wasnegligible. Transmission electron microscopy plan-view micrographs evidenced non-180° domainfringes in these films, where the vast majority of grains were 50–100 nm in diameter and showeda single set of domain fringes. Taken together, these measurements suggest that the pinning ofnon-180° domain walls is very strong in films with thickness less than 2mm. In thicker films,non-180° domain switching was evidenced when the poling field exceeded a threshold field. Thethreshold field decreased with an increase in film thickness, suggesting more non-180° domain wallmobility in thicker films. Non-180° domain switching in large grained PZT films was found to bemuch easier and more significant than in the fine grained PZT films. ©2001 American Institute ofPhysics. @DOI: 10.1063/1.1325005#

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I. INTRODUCTION

The demands for miniaturized mechanical devicescomponents that are integrated with microelectronics hled to great interest in the design and fabrication of micelectromechanical systems~MEMSs!. Incorporating smartmaterials such as ferroelectrics into MEMS designs canhance the sensing and actuation functions of such devAmong the ferroelectric materials, lead zirconate titan~PZT! is especially attractive for MEMS applications. Numerous MEMS devices based on PZT thin films depositedsilicon substrates have been demonstrated.1–4

For the construction of MEMS devices using PZT thfilms, one of the primary advantages is the large actuatiand high energy densities PZT can provide due to its lapiezoelectric coefficients. Since PZT is a multiaxial ferrelectric ceramic, there are both intrinsic and extrinsic conbutions to its dielectric and piezoelectric responses at ro

a!Author to whom correspondence should be addressed; [email protected]

1330021-8979/2001/89(2)/1336/13/$18.00

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temperature. The intrinsic contributions originate from tdielectric and piezoelectric responses of single domains;extrinsic contributions can be mainly attributed to domawall motions. For PZT ceramics with compositions near tmorphotropic phase boundary~MPB!, domain wall motionsare substantial at room temperature. The superior dielecand piezoelectric properties of PZT ceramics near the Mcan be at least partly attributed to the large extrinsic conbutions in these materials.

For the best performance of PZT MEMS devices, Pfilms with large piezoelectric coefficients (d33 or d31) aredesirable. Despite the tremendous efforts devoted to thevelopment of piezoelectric PZT thin films in the past 10 ythe electromechanical behavior of PZT films is incompletunderstood. The effective longitudinal and transverse pieelectric coefficients of PZT films have been measured tomuch smaller than those of their bulk counterparts, escially when the film thickness is small.5–7 It has been shownthat this can be partially attributed to the clamping effectthe substrate.7 It was also suggested that limited extrinsil:

6 © 2001 American Institute of Physics

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1337J. Appl. Phys., Vol. 89, No. 2, 15 January 2001 Xu et al.

contributions to the piezoelectric properties, due to theability to activate the non-180° domain wall motion, is aother important source that also reduces the piezoelectrin ferroelectric thin films.8,9

In this study, the domain wall motions and their contbutions to the dielectric and piezoelectric properties of sgel derived PZT films with composition near the MPB weinvestigated. The temperature dependence of the dieleconstant was measured to discriminate the intrinsic andtrinsic contributions to the dielectric response in PZT filmThe intrinsic dielectric properties were obtained from mesurements at 4 K by freezing out domain wall motioThrough separating the intrinsic and extrinsic dielectric prerties of PZT films, an evaluation of the intrinsic and tdomain wall contributions to the dielectric response wastained. Such analysis can provide important informatabout the influence of the grain size, preferred orientatthe film thickness, and the mechanical boundary condition the intrinsic responses, the domain structures, and domwall motion in these films. The extrinsic contribution to thpiezoelectric properties in PZT films was investigatedstudying the stress and electric field dependence of theezoelectric coefficients. In addition, by studying the inflence of uniaxial stress and dc electric field on the dielecproperties, both the ferroelectric and ferroelastic activitiesthe non-180° domain walls in these films were evaluatThe effect of the grain size and the film thickness on n180° domain wall pinning and extrinsic contributions to tdielectric and piezoelectric properties of PZT films is adiscussed. The domain structures in the PZT films werestudied using transmission electron microscopy and n180° domain walls were observed in all the films.

II. EXPERIMENTAL PROCEDURE

The PZT films investigated were all deposited oncoated Si substrates with a Zr/Ti ratio of 52/48~near theMPB!. Three sol–gel processes were used to preparePZT films to allow large variations in film thickness, grasize, and preferred orientation. The first one was2-methoxyethanol solution and rapid thermal anneal~RTA! process, which is described in detail elsewhere.10 PZTfilms with thicknesses from 0.25 to 3.4mm were preparedusing this method. Films thicker than 4mm prepared by thismethod often resulted in cracks due to the stress accumtion. To make thicker films, a second sol–gel method wused where acetylacetone was added into the precursortion of the first method. In addition, a 600–650 °C preanneing step was introduced into this process for densificationthe spin-coated layers to prevent further shrinkage duringfinal annealing at 700° by RTA.11 This adjustment enablepreparation of PZT films up to 10mm thick without cracks.The third sol–gel method was used to prepare PZT fiwith relatively large grain size. It used the same precursas the first two methods but a different solvent~acetic acid!.In addition, a conventional furnace annealing processused for the final crystallization of the PZT films in orderhave large grain size. This method was also able to depPZT films up to 10mm thick. Details of this sol–gel proces


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can be found elsewhere.12 The substrates used were Pt~111!/Ti/SiO2/Si and Pt~100!/SiO2/Si wafers commercially pro-duced by Nova Electronics and Tong Yang Central Labotories, respectively. For the Pt~111!/Ti/SiO2/Si substrates, thethickness of SiO2, Ti, and Pt layers were 1, 0.02, and 0.1mm, respectively. For the Pt~100!/SiO2/Si substrates, thethickness of SiO2 and Pt layers were 0.3 and 0.15mm, re-spectively.

X-ray diffraction ~XRD! patterns of the PZT films weremeasured at room temperature with a Scintag DMC-105fractometer~Scintag, Inc., Sunnyvale, CA! using CuKa ra-diation to determine the crystalline structure and preferorientation of the films. The surface morphology and crosectional microstructure of the films were observed usinJEOL JSM6300F field-emission scanning electron micscope~FE-SEM! and an atomic force microscope~AFM!.TEM bright field, dark field, and selected area electron dfraction studies were used to reveal the growth texturgrain sizes, and domain structures of the thin films. Sampfor TEM plan-view observation were prepared by first mchanical grinding and then ion milling from the substraside. The conditions for ion milling were as follows: HV55kV, current50.5 mA, at 12.5° with liquid nitrogen coolingTo reduce possible contamination, the thin film surface wprotected while ion milling. The cross-section TEM imagwere obtained using a FE Hitachi HF-2000 TEM.13

The film thickness was measured using an Alpha-S500 surface profilometer from Tencor Instruments~MountainView, CA!. Top electrodes of platinum, approximately 60Å in thickness and usually 1.6 mm in diameter, were spudeposited on top of the PZT film surface through a shadmask. The dielectric constant and loss tangent of the Pfilms were measured using either a Hewlett Packard 427multi-frequencyLCR meter or a 4192 A LF impedance anlyzer. All low-field tests were conducted at a frequency okHz and ac electric field of 0.5 kV/cm. The dielectric costant and loss were also measured as a function of dc pofield. For this measurement, the samples were poled unvarious dc electric fields for 1 min at room temperature adielectric measurement was made after 5 min aging. In socases, the dependence of the dielectric constant and lossgent on the amplitude of the applied electric field was ameasured at 1 kHz. The polarization hysteresis loops ofPZT films were measured using a Radiant TechnologRT66A ferroelectric tester, from which the saturation polaization Psat, the remanent polarizationPr , and the coercivefield Ec were determined. Typically, the maximum electrfield was 500 kV/cm and the measurement frequency wasHz for the hysteresis measurements. An external ampl~AVC Instrumentation 790 series power amplifier! was usedto allow high field measurements on thicker films.

The effectived33 of PZT films was measured by botdirect and converse piezoelectric measurements. The etive d33 was measured as a function of stress amplitude usthe pneumatic pressure charge technique, which appliestatic stress perpendicular to the film and measures theduced surface charge.7 The back of the silicon wafer wapolished with 1mm alumina powder to improve the accuracof the measurement. The electric field-induced longitudi


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1338 J. Appl. Phys., Vol. 89, No. 2, 15 January 2001 Xu et al.

strain of PZT films was measured as a function of applfield through a modified double beam laser interferomesystem.14 The effective d33 was calculated based on thstrain measurements. A small piece of Pt-coated silicon sstrate was adhered to the back of the specimen to incrthe reflectance from the back surface. A bipolar ac elecfield of 1 kHz was applied between the two electrodesgenerate an ac longitudinal strain in the specimen. Forpiezoelectric measurements, PZT films were poled underkV/cm at room temperature for 10 min.

The dielectric and ferroelectric properties of PZT filmwere also measured as a function of temperature from rotemperature to 4 K using a liquid helium Dewer cryogenisystem. The cooling rate was about 6 °C/min. The dielecconstant and loss measurements were made using a HePackard 4284 ALCR meter. The oscillation signal used watypically 1 kHz in frequency and 0.5 kV/cm in amplitudThe high field electrical properties, including the saturatpolarizationPs , the remanent polarizationPr , and the coer-cive field Ec were also measured as a function of tempeture down to 4.2 K. The measurements were made usinRT66A ferroelectric tester at discrete temperatures.

The dielectric properties of PZT films were measureda function of applied compressive uniaxial stress. The msurement setup is shown schematically in Fig. 1. A hydrastress rig was used to apply uniaxial stress to PZT films.sample was put between two stainless steel parts, with cties both above and beneath it. By pumping hydraulic oil inthese cavities, high pressure was exerted on both sides osample, thus imposing compressive normal stress onsample. Low field dielectric measurements were made ua Hewlett Packard 4192A LF impedance analyzer at aquency of 10 kHz and an amplitude of 1 kV/cm. High fieferroelectric hysteresis loop measurements were made ua RT66A ferroelectric tester. A 2 min waiting period wasallowed before any measurements were made. The mmum pressure applied was 20 MPa.

III. RESULTS AND DISCUSSION

A. Microstructural observations

The surface morphology and texture of the PZT filmdeposited by the three different sol–gel methods were sied by SEM, AFM, XRD and TEM. The XRD studies indcated that all of the films had the perovskite structure. Pfilms deposited by the first method had a strong prefer

FIG. 1. A schematic drawing of the experimental setup for uniaxial strapplication.


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orientation, which was controlled by the orientation of thebottom electrodes. Films deposited on substrates with^111&oriented Pt bottom electrodes were highly^111& oriented,while films deposited on substrates with^100& oriented Ptbottom electrodes were found to have a strong^100& pre-ferred orientation. The relative intensities of the x-ray peadid not change substantially with film thickness for eith^111& or ^100& oriented PZT films.10 TEM diffraction pat-terns taken at various depths in the^111& oriented films con-firmed that the texture remained strong when the film thiness was on the order of 0.3mm, and indicated that thetexture was weaker in the top~nonsubstrate! portion of thefilms that were 1mm thick.

SEM and TEM observations indicated that these filhad small grain sizes, as shown in Fig. 2. SEM, AFM, aTEM all confirmed that the grain size in the^111& orientedfilms was between 50 and 100 nm, while that of^100& ori-ented films was slightly larger~about 100–200 nm! due to alarger grain size of 100& oriented Pt bottom electrodesThese grain sizes remained relatively constant from theto the bottom of the film and no appreciable variation

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FIG. 2. Microstructures of^111& oriented films deposited using a2-methoxyethanol precursor solution and RTA process:~a! SEM surfacemorphology,~b! TEM cross-sectional image.


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FIG. 3. TEM plan-view micrographsof PZT ~52/48! film ~0.7 mm thick!showing various domain configurations in the main body of the film.Sample was annealed at 600 °C for 1h before picture~d! was taken.

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grain size was observed for films of different thickness.very thin layer was sometimes observed by SEM on topthe film surface, which appeared to be either amorphoupyrochlore.

The domain structures of 0.3, 0.7, and 1mm thick PZT~Zr/Ti 5 52/48! films prepared by the first method oPt~111!/Ti/SiO2/Si substrates were observed by plan-vieTEM. Microstructures were observed at different depthsthe films by ion milling the PZT films from the substrate sifor different times. Domain fringes were observed in all thrfilms. Figures 3~a!–3~d! are plan-view micrographs that wertaken in the middle of a 0.7mm film. The strong fringecontrast evidenced in these photos is consistent with prously published observations of non-180° domain walls15

No 180° domain walls were observed. The fact that somethe grains in these figures do not show fringe contrast canbe taken as an indication of the absence of non-180° domin these grains. The visibility of domain fringes is strongdependent on orientation. Tilting experiments~not shown!led to changes in contrast and indicated that most of grainthese films contained non-180° domains. Variations inthickness of the films did not result in obvious changes indomain configuration of the films that were observed. Toverall width of the domains was found to be on the order5–10 nm in all three films. The majority of the grains thshowed domain fringes contained only one set of fringes,more complex~multivariant! domain structures were observed in larger grains.

The domain wall density was observed to be signcantly lower in the region very near to the free surface inthree films, as shown in Fig. 4. The reason for the appadecrease in the number of domains in the near surface reis not obvious, but the importance of this observation is mgated by the fact that the affected region is but a small frtion of the thin films. The majority of grains in the 0.3, 0.


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and 1mm films were observed to have non-180° domainsthem.

XRD patterns showed that the films prepared by the sond method had a gradual change in the preferred orientawith increasing film thickness. When the film thickness wsmall (,2 mm!, the PZT films had 111& preferred orienta-tion on ^111& oriented platinum coated substrates, indicatinucleation from the bottom electrode. With increasing fithickness, the films became more and more randomlyented as the influence of the Pt/PZT interface was reduce16

The films also had small grain size~around 100 nm! whichdid not change much with film thickness@Fig. 5~a!#. Cross-sectional SEM images showed a layered structure forthick films, in which each layer corresponded to one cryslization step. There was no columnar growth found in thefilms, which was in agreement with the randomly orientXRD patterns.

All the films prepared using the third method westrongly^100& oriented although they were deposited on sustrates with^111& oriented Pt bottom electrodes. XRD paterns showed that the relative intensity of the^100& peakincreased with increasing film thickness.12 Unlike the filmsannealed by RTA, films prepared by this method~crystal-lized using the conventional box furnace! had larger grainsizes. The average grain size increased with the film thness, typically ranging from 300 to 700 nm. Figure 5~b!shows the AFM surface image of a 3mm thick ^100& ori-ented PZT 52/48 film prepared using this method, whgives an average surface grain size of 500 nm. It was afound from the cross-sectional SEM study that dense layestructures with well-defined grains and columns wereserved in these films.17


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B. Temperature dependence of the dielectric andferroelectric properties

Two experimental methods have been used to sepaintrinsic and extrinsic responses in ferroelectric ceramThe first one is measurement of the frequency dispersiothe dielectric response.18,19At high enough frequencies~wellinto the GHz range!, extrinsic contributions are too slow tcontribute to the dielectric response, leaving only the intrsic response. This experiment requires frequencies inGHz range, and can only give information about the clampdielectric properties. The other type of experiment to dcriminate the intrinsic and extrinsic response is a measment of the temperature dependence of the dielectricpiezoelectric constants.20,21Since domain wall motion, phasboundary motion, and defect reorientation are thermallytivated processes, they can be frozen out at temperatures0 K. Therefore, the piezoelectric and dielectric propertmeasured at this temperature are completely from the insic contribution. This method can provide informationboth dielectric and piezoelectric responses. But it should

FIG. 4. Plane-view morphologies observed near the surface of the 0.3mm~a! and 0.7 mm ~b! thick PZT ~52/48! films prepared using a2-methoxyethanol precursor solution and RTA process.


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noted here that the properties measured near 0 K may not bedirectly related to the properties at high temperatures.

In this work, the temperature dependence of the low fielectrical characteristics of the PZT films was investiga

FIG. 5. ~a! SEM surface morphology of a 7mm thick PZT film preparedusing the second sol–gel method.~b! AFM surface image of a 3mm thickPZT film prepared using the third sol–gel method.

FIG. 6. Temperature dependence of the dielectric constant of sol–gel52/48 films at l kHz. Films 1 and 3 are100& oriented, films 2, 4, 5, 6, and7 are^111& oriented.


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from room temperature down to 4.2 K. Figure 6 showsresults of the dielectric constant as a function of temperafor films with different thicknesses, orientations, and grasizes. At room temperature, there are large differencesserved for the dielectric constants of the various PZT filinvestigated. The dielectric constant increases with fithickness for films with the same orientation and grain siand is larger in films with larger grain sizes. As the tempeture decreases, the dielectric constants of all the filmscrease. However, the temperature coefficients of capacitare different. Eventually the dielectric constants of all thefilms converge to approximately the same value as the tperature approaches 0 K~except films,0.5 mm thick!.

The results from Fig. 6 indicate that although the sol–PZT films have much different dielectric constants at rotemperature~due to the difference in thickness and grasize!, they all have similar intrinsic contributions to the delectric constant near 0 K. A close look at the dielectconstant at 4 K showed that there is still a weak dependenof the intrinsic dielectric constant on the film thickness. This probably due to the effect of the top surface layer, andan interfacial layer between the PZT and the bottom etrode. These layers usually have lower dielectric constathan ferroelectric PZT, and thus can result in a smallerparent intrinsic dielectric constant.

If the surface and interfacial layers have a combincapacitance ofCint for all the films, then the apparent capactance of the filmCtot is related toCint as follows:

1/Ctot51/Cfilm11/Cint , ~1!

whereCfilm is the capacitance of the bulk portion of the PZfilms. Also

1/Ctot5t/«0«S11/Cint ~2!

in the above equations, wheret is the film thickness andS isthe area of the top electrode. This assumes that the interflayer is negligibly thin compared to the measured thickneFigure 7 shows the plot of 1/Ctot versus film thicknesst for^111& oriented fine-grained PZT films. The linear relatioship between 1/Ctot and t along with the nonzero intercepfor the plot support the hypothesis that the thickness depdence of the dielectric constant at 4 K is mainly due to the

FIG. 7. Reciprocal of the film capacitance measured at 4 K and 1 kHz as afunction of film thickness for 111& oriented fine-grained PZT films.


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surface and interfacial layers. It is expected that with iproved film processing this effect could be eliminated, orleast substantially lessened.

Since all the films have similar intrinsic dielectric constants at 4 K, it is reasonable to expect that they also hsimilar intrinsic contributions to the dielectric constantroom temperature. Part of the difference in dielectric costants at room temperature can be attributed to the surand interfacial layers. However, deviation in the linear retionship between the reciprocal capacitance and thethickness was found at room temperature. This suggestsvariations in extrinsic contributions to the dielectric responalso contribute to the difference in dielectric constantroom temperature. To quantitatively separate the effect ofsurface and interfacial layer on the dielectric constant offilms at room temperature, the capacitance of the surfaceinterfacial layers must be determined. This was attemphowever, it was not possible to determine the intercept inCtot versust line at 4 K accurately enough, due to the lackdielectric data in the very thin film thickness range. Vesmall variations in this number lead to very large over-undercorrections of the room temperature dielectric datathin films. In addition, the temperature dependence ofdielectric constant of the surface and interfacial layers is aunknown, which adds more complexity in the calculationNevertheless, it is clear that the interfacial layer doesaccount for all of the differences in room temperature dieltric constants for films with different thickness.

The existence of extrinsic contributions to the dielectconstant in PZT films is manifested in the ac driving-fiedependence of the dielectric constant at room temperaand 4 K ~Fig. 8!. At 4 K, the nonlinearity of the dielectricconstant is very small, reflecting the intrinsic dielectric rsponse of the material. However, at room temperature,dielectric constant increases with the amplitude of theplied electric field, showing a large dielectric nonlinearitythese films. This dielectric nonlinearity is believed to be asociated with extrinsic sources and can be attributed tomain wall motion.22,23

Despite the large difference in the texture and grain sbetween the PZT films and bulk ceramics, the dielectric c

FIG. 8. The dependence of the dielectric constant at 1 kHz on the ampliof the ac electric driving field for a 0.6mm thick PZT film.


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stant of PZT films at 4 K was similar to that of PZT buceramics at 4 K.24 There are a number of factors that miglead to a variation in the measured intrinsic dielectric costant between bulk and thin film PZT specimens. First,PZT films have strong preferred orientation; thereforepolarization vectors in these films are concentrated alseveral specific directions relative to the electrodes. Tmay give different averages of the intrinsic dielectric costant for films with different textures, and for thin films anbulk PZT. Second, with a decrease in grain size, there isincrease in the internal stress due to the difficulty in formnon-180° domains. Therefore the intrinsic dielectric constshould increase as the grain size decreases.20 This effectwould result in a larger dielectric constant in thin films thin bulk ceramics. Third, due to the constraints of the sstrates, the measured dielectric constant of PZT films ileast partially clamped~clamping should decrease« r). Fi-nally, since the thickness of the PZT films is much smathan bulk samples, low dielectric constant layers at the Pelectrode interfaces can cause a noticeable decrease imeasured dielectric constant of thin films while they havirtually no effect on bulk samples. The experimental resuindicate that these factors are either very modest or thatcounterbalance each other. Thus, the intrinsic dielectric cstants are similar for bulk ceramics and most of the thin fiPZTs.

From the phenomenological theory developed for PZthe intrinsic contribution to the relative dielectric constantroom temperature for PZT 52/48 bulk ceramics was callated to be about 670.20 Using this value as an approximatiofor the intrinsic contribution to the dielectric constant of PZfilms at room temperature, it was estimated that 25%–5of the total dielectric constant at room temperature in thfilms arises from extrinsic sources. The extrinsic contributto the dielectric constant is likely to increase with film thicness, and is larger in films with larger grain size.

The extrinsic contribution to the dielectric and piezelectric response in PZT ceramics with compositions nthe MPB is believed to be mainly from domain wall motioand phase boundary motion. There are two types of domwalls in ferroelectric materials: 180° domain walls and no180° domain walls. Since there is no strain change associwith 180° domain wall motion, it contributes only to thdielectric properties. Thus, 180° domain walls are purferroelectric walls. In contrast, non-180° domain walls aboth ferroelectric and ferroelastic; they can be excitedboth external electrical and mechanical signals. Such wmotion can cause changes in both the polarization andstrain, so that it can contribute to both the dielectric apiezoelectric properties.

As will be shown in Secs. III C and III D, non-180° domain wall motion in the small grain (,0.2 mm! and thin~thickness smaller than 3mm! PZT films is negligible.Therefore the extrinsic contribution to the dielectric constin these films was most probably from 180° domain wmotions or from phase boundary motion. Due to the coexence of tetragonal and rhombohedral phases, phase boaries may exist in PZT films with MPB composition. Similato the non-180° domain wall motion, phase boundary mot


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changes both the polarization and the strain, and thus ctributes to both the dielectric and piezoelectric propertiSince the extrinsic contribution to the piezoelectric respoin most of the PZT films is very small, it suggests that tphase boundary motion in these films is also limited. Thefore, the major source for the extrinsic contribution to tdielectric properties in PZT films can be attributed to 18domain wall motion.

The mechanism for the increase in extrinsic contributto the dielectric constant with film thickness is not clear. Opossibility is that there may be pinning centers located atinterface between the bottom electrode and the PZT filmthicker films, fewer grains are affected by these pinning cters. Therefore, the 180° domain wall motion becomeasier. There is also evidence from the TEM observatithat the domain wall density is lower near the film surfacThe increased extrinsic contribution to the dielectric constin large grained films is believed to be due to the high dmain wall density and less domain wall pinning.~It has beenshown in bulk PZT ceramics that the domain wall densdecreases with decreasing grain size in the deep submirange.!20 There is also more grain boundary area in the fingrained films where space charge could be trapped. Aresult, domain wall pinning is expected to be strongerfine-grained films. In addition, non-180° domain wall motiomay contribute to the dielectric constant in thick and largrained films, as indicated by the extrinsic contribution to tpiezoelectric coefficient in these films~see Sec. III C!.

The dielectric loss of the PZT films was also measuas a function of temperature from room temperature to 4.~shown in Fig. 9!. Compared to the dielectric constant, thtemperature dependence of the loss factor was more comcated. In general, there was a broad anomaly in thefactor centered at temperatures between 200 and 250 K.films with thicknesses larger than 1mm, the dielectric loss attemperature around 200 K was larger than its room tempture value. As the temperature dropped below 50 K, a radecrease in the loss factor was observed in all films. Wthe temperature approached 0 K, the dielectric loss in thfilms converged to about 0.6%, though there was a livariation.

The behavior of the dielectric loss with temperature

FIG. 9. Temperature dependence of the dielectric loss of PZT 52/48 film1 kHz.


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still not well understood. Similar behavior was also observin PZT bulk ceramics.21 The authors suggested that there aseveral relaxation processes depending on compositiondoping. The relaxation processes could be connectedthe impurity ions and with domain walls or phase-boundmotions. At temperatures near 0 K, these thermally activaprocesses were frozen out, therefore the intrinsic sindomain properties dominated the measured dielectric los21

A similar mechanism may also be responsible for the dietric loss behavior in PZT films.

The P–E hysteresis loop of PZT films was also mesured as a function of temperature from room temperatur4 K @Fig. 10~a!#. Figure 10~b! shows the temperature depedence of the saturation polarization and remanent polartion of a 0.6mm PZT 52/48 film. The saturation polarizatiowas almost independent of temperature between 300 anK. As the temperature continued to decrease, there wslight decrease in the saturation polarization. In contrast,remanent polarization showed a much larger change withtemperature. It monotonically increased with decreasing tperature. As a result, the difference between the satura

FIG. 10. ~a! P–E hysteresis loops at different temperatures.~b! Saturationand remanent polarizations as a function of temperature. Data meafrom a 0.6mm ^111& oriented PZT 52/48 film at maximum field of 50kV/cm.


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polarization and the remanent polarization decreased attemperatures. Figure 11 shows the coercive field of themm PZT film as a function of temperature. The coercive fiewas found to increase with decreasing temperature. Astemperature dropped below 50 K, a rapid increase incoercive field was observed with the decrease of tempture.

The temperature dependence of the coercive field incates that domain reversal in PZT films becomes more dcult as the temperature decreases. Since non-180°domaorientation is very limited in these films, domain reversalmainly achieved by 180° domain wall motion. With decreaing temperature, the domain wall motion becomes moreficult, and thus a larger electric field is needed to accompit. The rapid increase of the coercive field at temperatubelow 50 K indicated a dramatic decrease in the 180°main wall mobility, which is also illustrated by the rapidecrease in the dielectric loss in this temperature range.increase of the remanent polarization with decreasing tperature may also be due to the reduced domain wall moity. It was suggested that the large difference between sration and remanent polarizations in PZT films is most likea result of significant domain reversal after the removalthe electric field. At low temperatures, the back switchingthe domains is partially frozen out. This therefore results ismall difference between the saturation and remanent poizations. The slight decrease in the saturated polarizatiotemperature near 0 K can be attributed to the inability toreverse all the domains at that temperature, using akV/cm excitation field.

C. Stress and field dependence of dielectric andpiezoelectric properties

Due to the ferroelastic nature of non-180° domain wain ferroelectric materials including PZT, it is reasonableexpect that the properties of these materials can be inenced by mechanical stresses. To investigate the ferroelactivity of the non-180° domain walls, the low and high fieelectrical characteristics of^111& PZT films were measuredas a function of applied normal stress~perpendicular to thefilm plane!. Figure 12 shows the effect of the applied normstress on the dielectric constant of a poled 1mm thick ^111&PZT film, which has an average grain size of about 0.05–mm. The result indicates that the dielectric constant has vweak stress dependence, increasing less than 2% for an

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FIG. 11. The coercive field of a 0.6mm ^211& oriented PZT film as afunction of temperature. The maximum field applied was 500 kV/cm.


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plied normal stress up to 20 MPa. The change of the dietric constant with applied stress is reversible, i.e., removathe applied stress results in the full recovery of the dielecconstant. On the other hand, the dielectric loss is almindependent of the applied stress within the stress rangein this investigation. The ferroelectricP–E hysteresis loopwas also measured during the application of normal strThere was no noticeable change found in either the remapolarization or the coercive field with applied stresses up20 MPa. Both high field and low field measurements wperformed on several films with film thickness ranging fro0.5 to 3mm. Similar results were obtained on all the samptested.

The uniaxial stress effect on the dielectric and piezoetric characteristics of PZT ceramics has been the subjecmany investigations.25–29 Krueger found that for an appliecompressive uniaxial stress of 20 MPa, there were 12%5% increases in the dielectric constant for hard and soft Pceramics~poled!, respectively.25 This behavior is believed tobe due to the increased domain wall contributions todielectric and piezoelectric properties caused by the appuniaxial stress. In hard PZT bulk ceramics, the defect dipoare easy to reorient due to the mobile nature of these defThey tend to align with the polarization vectors of the dmains to stabilize the domain structure. A displacementhe domain walls then increases the material’s free energthat the domain walls are inactive against external excitatTherefore, domain wall pinning in hard PZT is strong, resuing in small extrinsic contributions to the dielectric constaBy applying a compressive stress parallel to the polarizavector, the ferroelectric domains tend to reorient to directiomore perpendicular to the stress, which results in a mstable domain configuration. Therefore the domain wall mtion is much enhanced by the compressive stress, so thadielectric constant and tand increase as extrinsic contributions increase.29 Since domain wall pinning is less severesoft PZT, a smaller stress dependence of the dielectric cstant was found in soft PZT over this stress range.

The stress dependence of the dielectric constant ofpoled PZT films was much smaller than that in bulk PZceramics. In addition, there was no increase in the dielecloss with the applied stress. These results suggest that180° domain wall motion in response to the applied stres

FIG. 12. The effect of applied normal stress on the dielectric constant1 mm thick poled^111& PZT 52/48 film which has an average grain sizeabout 50–100 nm.


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negligible in these films for this stress range. Since strodomain wall pinning in these films is expected due to thsmall grain size and to interfacial effects, a large dependeof the dielectric constant and loss on the applied stress wobe expected if the non-180° domain walls were ferroelacally active. The effect of normal stress on the high fiecharacteristics of the PZT films also suggests that the feelastic activity of the non-180° domain walls is negligiblOtherwise, a decrease in the remanent polarization anchange in the coercive field would be expected. Howevthey were not observed.

Both the low and high field electrical measurements afunction of the normal stress are consistent with the hypoesis that ferroelastic motion of the non-180° domain wallslimited. This is similar to previous reports on measuremeof PZT films under biaxial stresses.6 It was proposed that thesubmicron grain structure is one critical reason for thishavior in sol–gel PZT films.6,8 Besides, the reduced domawall density and strong pinning by the film/substrate intface, point defects~such as lead and oxygen vacancies! mayalso play roles in reducing the non-180° domain wall activin PZT films.6

It is believed that the primary source for the extrinspiezoelectric response in ferroelectric materials is the n180° domain wall motion.30,31 The fact that the non-180°domain wall motion is severely limited in PZT films suggesthat little extrinsic contribution to the piezoelectric coefcient is expected in these films. To verify this, the effectid33 of PZT films was measured using the pneumatic presscharge technique as a function of the amplitude of theplied stress. Figure 13 shows the effect of stress ampliton thed33 coefficients of both a 1mm thick ^111& PZT filmwith 50–100 nm average grain size and a PZT-5A bulkramic sample.

The d33 of the bulk PZT ceramic increased with thstress amplitude, showing a significant piezoelectric nonearity. This nonlinearity is believed to be of extrinsic natuand can be attributed to non-180° domain wall motion.22,23 Itwas found in barium titanate ceramics that both thed33 co-efficient and its nonlinearity with stress amplitude decreaas the grain size decreased, due to the reduction of non-

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FIG. 13. The effect of stress amplitude on thed33 coefficients of a 1mmthick ^111& oriented PZT film ~50–100 nm average grain size! and aPZT-5A bulk ceramic.


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domain wall motion in fine grain ceramics.32 Unlike the PZTbulk ceramics, measurements on the fine grained~50–100nm in grain size! PZT thin films ~thickness less than 2mm!showed that there was no increase in the effectived33 withan increase of the applied stress amplitude up to 6 MPa.result is again a strong indication that the ferroelastic motof the non-180° domain walls in these films is limited.

The piezoelectric nonlinearity of the PZT films was almeasured as a function of the applied electric driving fiusing double-beam laser interferometry. The PZT films stied here were prepared using the acetylacetone mod2-methoxyethanol solution and RTA process. The film thicness was between 1.5 and 6.7mm. These films had an average grain size of 100 nm and almost random orientation.effective d33 of the films was found to increase with filmthickness under an ac field of 2 kV/cm and 1 kHz~Fig. 14!.To investigate the piezoelectric nonlinearity of the effectd33, measurements were made as a function of subcoerac electric field~Fig. 15!. For films with small thickness~1.5mm!, the effectived33 remained unchanged as the appliedfield increased to 10 kV/cm. Above that, only a small icrease in the effectived33 was measured. For the 6.7mmthick film, the onset of the piezoelectric nonlinearity o

FIG. 14. The effectived33 of PZT films as a function of film thickness atkHz. The films had an average grain size on the order of 0.1mm and werepoled at room temperature under an electric field of 250 kV/cm.

FIG. 15. Normalized effectived33 of PZT films as a function of subcoercivac electric field at 1 kHz. The films had an average grain size on the ord0.1 mm and were poled at room temperature under an electric field ofkV /cm.


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curred below 4 kV/cm, and the effectived33 increased muchmore rapidly with higher ac electric fields. Similar behaviwas also observed by Kholkin.33

The piezoelectric coefficients were found to increawith the driving field for PZT bulk ceramics.22 Accompaniedby the onset of this nonlinearity, a minor hysteresis wasserved. The increase of nonlinearity was also accompaby an increase in loss. Applying a dc bias field or decreastemperature reduced the piezoelectric nonlinearity. Thesesults indicated that the nonlinear behavior is of extrinsicture and related to the irreversible non-180° domain wmotions.22 The relatively large piezoelectric nonlinearity ithick PZT films suggests that there is significant irreversinon-180° domain wall motion at sufficiently high ac electrfields. The decrease of the threshold field for piezoelecnonlinearity with film thickness suggests that the degreenon-180° domain wall pinning becomes less severe asfilms become thicker. This, in turn, suggests that there isincreasing tendency to have extrinsic contributions topiezoelectric response as the films become thicker. Agthe large threshold field for piezoelectric nonlinearity andsmall degree of piezoelectric nonlinearity in thin PZT filmagrees with the hypothesis that there is strong non-180°main wall pinning and thus little extrinsic contribution to thpiezoelectric coefficient in thin PZT films.

Unlike the effectived33, which showed small nonlinearity with ac field for most of the films, the dielectric constaof all PZT films showed very large amplitude dependen~Fig. 16!. All the films displayed a significant dielectric nonlinearity, which was found to increase with the film thickness. The large difference between the nonlinearity in pieelectric coefficients and dielectric constants suggests180° domain wall motion is much more significant than no180° domain wall motion. As a result, a majority of thextrinsic contribution to the dielectric constant in PZT filmis due to the 180° domain wall motion. The increase indielectric nonlinearity with film thickness suggests that textrinsic contribution to the dielectric constant also increawith film thickness, which is consistent with the results frothe temperature dependence of the dielectric constant msurements.

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FIG. 16. The normalized dielectric constant of PZT films as a function ofelectric field at 1 kHz. Films were randomly oriented and had an avergrain size of 0.1mm.


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Kholkin also observed a large difference in the nonlinebehavior of dielectric constants and piezoelectric coefficiein PZT films.33 However, the author attributed this differencto the geometry of the PZT films. It was proposed thatlarge dielectric nonlinearity was due to extensive non-1domain wall motion, and the small non-linearity in effectivd33 was due to the111& orientation and the tetragonal struture of the films. The author argued that since all the allowpolarization vectors have equal projections onto the subsnormal, 90° domain wall motion should not produce amechanical strain, and therefore did not contribute topiezoelectric response.33 However, this argument is not applicable to the randomly oriented, MPB films used in thinvestigation. In addition, if 90° domain wall motion is important in tetragonal PZT films, then the effectived33 of^100& PZT films is expected to be larger than^111& PZTfilms due to extrinsic contributions in100& films. However,our experimental results showed that this was not the ca34

D. Non-180° domain switching in PZT films

Experimental results in Sec. III C showed that the extrsic contribution to the effectived33 in fine grain PZT films issmall due either to strong pinning of the non-180° domwalls or to a low density of ferroelastic walls in the filmThere was also evidence that with an increase in film thiness, non-180° domain wall motion might begin to contrute to the piezoelectric response. It has been suggestedthe degree of non-180° domain wall pinning is also stroninfluenced by the grain size of the PZT films.8 In this section,the activity of the non-180° domain walls of the PZT filmand their relationship with the film thickness and grain swere investigated.

The activity of non-180° domain walls can be relatedthe ease of non-180° domain switching. In general, domswitching occurs via domain wall motion driven by the eternal field. Therefore the smaller degree of domain wpinning, the more extensive domain wall motion shouldthus the more easily domain switching should occur. Domswitching can be identified by monitoring the change ofrelative intensity of the 001& and ^100& peaks in the XRDpatterns for tetragonal PZT ceramics and thin films.8,35 Thismethod, though, cannot be used for PZT films at the Mcomposition. However, single domain PZT with the MPcomposition has very large anisotropy in the dielectric prerties, i.e.,a domains have much larger dielectric constathan c domains.24 Therefore, a large change in the intrinscontribution to the dielectric constant is expected as a reof non-180° domain switching. This may be used to identnon-180° domain switching in PZT films.

The effect of dc poling on the dielectric constant of tPZT films was investigated at room temperature first. Dieltric measurements were made under small field~0.5 kV/cm!without dc bias. PZT films were poled at different dc fielfor 1 min at room temperature. Then a 5 min aging time wasallowed before measuring the dielectric constant. Figureshows the dielectric constant as a function of dc poling fie

For poling fields below the coercive field, all the filmshowed small increases in both the dielectric constant


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loss immediately after poling. The dielectric constant aloss gradually decreased with aging time, and finallyturned to their original values after about 72 h aging. Tobservation suggests that the initial increases in the dieleconstant and loss result from deaging. 180° domain wmay be depinned by the dc poling field through the formtion of a metastable domain structure or redistribution ofdefect dipoles, thus leading to larger 180° domain wall cotributions to the dielectric response in these films. Simideaging behavior was also observed after heat treatmetemperatures above the Curie temperature, in which themain structure and space charge distributions werechanged.

As the dc poling field continued to increase, the dieletric constant remained flat for small grained PZT films whthe film thickness was less than 2mm. The dielectric con-stant changed little through dc poling in either^100& or^111& oriented PZT thin films up to the breakdown fieldThis behavior suggests that the non-180° domain wall pning is so strong in these films that the ferroelectric switcing of the non-180°domain is not achievable at room teperature. This result is consistent with the x-ray diffractistudy on tetragonal fine grain PZT thin films, which alshowed that electrical switching of 90° domains in the sumicron films is severely limited.8 However, in thicker films adecrease in the dielectric constant was observed whenpoling field exceeded a certain level, while little change wobserved in the dielectric loss. When heated aboveTc thedielectric constant recovered to its original value, sochange was not due to cracking or electrode delaminatio

The decrease in the dielectric constant in thick PZT filcan be explained by non-180° domain switching. Strong ping of PZT films may switch some of thea-oriented domainsinto c-oriented domains, which may result in a decreasethe measured room temperature dielectric constant sinceintrinsic contribution usually dominates the dielectric behaior of PZT films. Similar results on permittivity changehave been reported in XRD studies on hot poling of tetronal PZT films, which showed that the room temperatudielectric constant decreased as more domains switched

FIG. 17. Weak field dielectric constant of PZT films as a function ofpoling field.


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Since the non-180° domain walls are strongly pinnedPZT films, some minimum electric field is required to ovecome the potential barrier to realize extensive domain wmotion. Therefore, the threshold field for the non-18switching ~the onset of the decrease of the dielectric costant! can be used as an indication of the degree of non-1domain wall pinning in PZT films. It was found that for filmwith similar average grain sizes on the order of 0.1mm, thethreshold field decreased with increasing film thickne~from 600 kV/cm for a 3mm thick film to 160 kV/cm in a6.5 mm thick film!, indicating a reduction in the degree othe non-180° domain wall pinning. The percentage changthe dielectric constant through dc poling increased slightlythicker films and was less than 20% in all the fine grain Pfilms, indicating that only a small fraction ofa domainscould be switched toc domains by the electric field.

Compared to the film thickness, the grain size of tfilms had a much stronger effect on ferroelectric non-18switching in PZT films. The large grain PZT films preparusing conventional furnace annealing showed much msignificant decreases in the dielectric constant~40%–50%! ata much smaller threshold field. The large decrease in dietric constant suggests that there is a significant volume ftion of a domains which are switched toc domains by dcpoling when the poling field exceeds the threshold field. Tsmall threshold field for non-180° domain switching indcates that the pinning of non-180° domain walls in thelarge grain films is much less severe than that in the smgrained films. It was also found that with an increase in grsize, there was a decrease in the threshold field~from 100kV/cm in films with an average grain size of 0.47mm to 60kV/cm in films with an average grain size of 0.68mm!.

The grain size effect on non-180° domain switchingPZT films agrees very well with the results from PZT buceramics. In fine grain PZT bulk ceramics, it was found twith the decrease of grain size, there was a large decreaboth the remanent polarization and piezoelectric coefficiebut an increase in the coercive field. This behavior waslieved to be due to the clamping of domain walls exhibitby neighboring grains and grain boundaries in fine-grainPZT.20,36,37

To verify that there is non-180° domain switching acompanying the decrease of the dielectric constant aftepoling, the dielectric constant of the PZT films was measuas a function of temperature before and after poling. Sithe dielectric constant at room temperature is composeboth intrinsic and extrinsic contributions, a variation in textrinsic contribution to the dielectric constant may alsosult in a change in the room temperature dielectric constThe domain configuration of the PZT films can be changby dc poling, and this could also lead to a change inextrinsic contribution to the dielectric constant. On the othhand, non-180° domain switching changes the observedelectric constant of the PZT films via the intrinsic contribtion. Figure 18 shows the temperature dependence ofdielectric constant of a 3mm thick conventional furnace annealed PZT film which had an average grain size of 0.5mm.


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The large decrease of the dielectric constant at 4 K aftepoling clearly showed that the intrinsic contribution to thdielectric constant was reduced. This result indicated tnon-180° domain switching did accompany the decreasethe room temperature dielectric constant in coarse graiPZT films. In contrast, there was little change in the dieletric constant either at room temperature or 4 K for finegrained thin films, again suggesting little non-180° domaswitching in these films.

IV. CONCLUSIONS

Domain wall motions and extrinsic contributions to thdielectric and piezoelectric response in sol–gel derived Pfilms were investigated. The temperature dependence ofdielectric properties showed that although PZT films hmuch different dielectric constants at room temperaturefilms with different film thickness, grain size, and preferrorientation, they all had similar intrinsic contributions to thdielectric constant. It was estimated that at room tempera25%–50% of the dielectric response of PZT films arisfrom the extrinsic contribution. The extrinsic contributionthe dielectric constant in PZT films was mostly attributed180° domain wall motion, and was likely to increase wiboth film thickness and grain size.

TEM observations of PZT films prepared from2-methoxyethonal solution and RTA showed the presencnon-180° domain fringes in the vast majority of grains in tfilms with thickness ranging from 0.3 to 1mm. Multivariantdomain fringes were observed in a few large grains, butmajority of grains were of comparable size and containonly stripe domains. No significant changes in domain cfiguration or density were observed in these films.

By contrast, ferroelastic non-180° domain wall motiowas found to be limited in fine grain PZT films, as waindicated by the small effect of normal stress on the low ahigh field electrical characteristics of the films and the linerelationship between the piezoelectric response and the sexcitation amplitude. This, in turn, suggested that the extsic contribution to the piezoelectric coefficient in the fingrain PZT films was small. However, as the films becathicker (.5 mm!, significant ferroelectric non-180° domai

FIG. 18. Temperature dependence of the dielectric constant before anddc poling. The film was 3mm thick annealed by a conventional furnace.poling was performed at room temperature under 150 kV/cm.


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wall motion was observed under high external excitatiwhich suggested that the degree of non-180° wall pinndecreased with increasing film thickness.

The activity of the non-180° domain walls was studithrough the characteristics of non-180° domain switchiFor fine grained films with film thickness less than 2mm,non-180° switching was negligible, suggesting that the pning of non-180° domain walls was very strong. As the filmbecame thicker, a decrease in the dielectric constant wasserved when the poling field exceeded a threshold fiwhich was attributed to non-180° domain switching. Tthreshold field for switching was found to decrease withincrease in film thickness, suggesting larger non-180°main wall mobility in thicker films. It was also found that thnon-180° domain switching in large grained PZT films wmuch easier and more significant than that in fine graiPZT films.

ACKNOWLEDGMENTS

The authors would like to thank the National ScienFoundation for providing financial support for this work.

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Domain wall motion and its contribution to the dielectric and piezoelectric properties of lead...

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