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Relaxor FerroelectricPolymers–Fundamentals and ApplicationsQin Chen a , Kailiang Ren a , Baojin Chu a , Yiming Liu a , Q. M. Zhanga , Vid. Bobnar b & A. Levstik ba Department of Electrical Engineering and Materials ResearchInstitute , The Pennsylvania State University , University Park , PA ,16802b Jozef Stefan Institute , P.O. Box 3000, Ljubljana , SloveniaPublished online: 10 Oct 2011.

To cite this article: Qin Chen , Kailiang Ren , Baojin Chu , Yiming Liu , Q. M. Zhang , Vid. Bobnar & A.Levstik (2007) Relaxor Ferroelectric Polymers–Fundamentals and Applications, Ferroelectrics, 354:1,178-191, DOI: 10.1080/00150190701454891

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Ferroelectrics, 354:178–191, 2007Copyright © Taylor & Francis Group, LLCISSN: 0015-0193 print / 1521-0464 onlineDOI: 10.1080/00150190701454891

Relaxor Ferroelectric Polymers–Fundamentalsand Applications

QIN CHEN,1 KAILIANG REN,1 BAOJIN CHU,1 YIMING LIU,1

Q. M. ZHANG,1,∗ VID. BOBNAR,2 AND A. LEVSTIK2

1Department of Electrical Engineering and Materials Research Institute,The Pennsylvania State University, University Park, PA, 168022Jozef Stefan Institute, P.O. Box 3000, Ljubljana, Slovenia

We present recent studies in the fundamentals and applications of the relaxor ferro-electric polymers, i.e., Poly(vinylidene fluoride-trifluroethylene-chlorofluoroethylene)(P(VDF-TrFE-CFE) terpolymers and high energy electron irradiated P(VDF-TrFE)copolymers. We show that the dynamic processes in these relaxor ferroelectric polymersare very similar to these observed in various polar-glasses. We further show that thelarge and reversible polarization change in these polymers leads to giant electrostric-tion and large electro-optic effect. By employing active electric boundary conditions,we demonstrate that the large electromechanical responses in these electroactive poly-mers can be made use of effectively for energy harvesting with high harvested electricenergy density and efficiency. Terpolymer composites with ZnS nanoparticles providean advanced E-O polymer which refractive index can be varied widely by small amountof ZnS in the composites while maintaining the large E-O effect of the matrix.

I. Introduction

One area in which Dr. Vitaly L. Ginzburg has made fundamental impact in the ferroelectricis the phase transitions, which has lead to many decades of studies and great advances in theferroelectrics. Besides their fundamental scientific impacts, phase transitions in ferroelectricmaterials have also found broad range of applications. In this paper, we present our recentworks in the relaxor ferroelectric polymers, in which the normal ferroelectric transition ismodified by random defect fields and a reversible phase transformation between the polarand non-polar molecular forms is realized. As a result of the random defect modification,the normal ferroelectric polymer P(VDF-TrFE) is converted into a relaxor ferroelectric. Weexamine the dynamic behavior of the ferroelectric relaxor polymers. We will also present theexperimental results which show giant electrostrictive response in the polymer, due tothe reversible transformation in the molecular conformations and nanostructures betweenthe polar- and non-polar forms. Moreover, the large polarization change associated with themolecular conformation change also leads to very large electro-optic effect in the relaxorferroelectric polymers, which will also be presented in this paper.

Received in final form March 25, 2007.∗Corresponding author. E-mail: qxy1@psu.edu

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 179

Figure 1. The small signal dielectric constant and dielectric loss of the terpolymer P(VDF-TrFE-CFE) 63/37/7.5 mol% as a function of temperature measured at different frequencies (for dielectricconstant, curves from top to bottom and for the dielectric loss, curves from bottom to top: 100 Hz,1 kHz, 10 kHz, 100 kHz, and 1 MHz).

II. Relaxor Behavior of the Electroactive Polymers

Presented in Fig. 1 is the temperature dependence of dielectric properties of P(VDF-TrFE-CFE) at 63/37/7.5 mol% measured at several frequencies. A distinctive feature of the dielec-tric data is the broad dielectric constant maximum which shifts to higher temperature withfrequency, typical for relaxor ferroelectrics [1, 2]. Furthermore, structural data indicate thatthere is no macroscopic phase transition associated with the broad dielectric constant max-imum [3, 4]. Detailed dielectric data analysis has shown that the broad dispersive dielectricmaximum, as in dipolar glass, is a result of the dipolar freezing transition in which differentfrequency component of the polarization response will freeze at different temperature [5,6].

The dynamic processes in the relaxor P(VDF-TrFE-CFE) terpolymers as well as thehigh energy electron irradiated P(VDF-TrFE) copolymers were investigated and it wasfound that the dynamic processes in these relaxor ferroelectric polymers are very similar tothese in the classical inorganic relaxors such as lead magnesium niobate [5, 6]. The resultsshow that the broad dispersive dielectric maximum is a result that, as in dipolar glasses,the dielectric constant ε1 (the linear dielectric constant) at a certain temperature whichdepends on the experimental time scale, i.e., frequency, starts to deviate from its static value(Fig. 2). The temperaure dependence of the static dielectric constants, measured via chargeaccumulation method, is also shown in Fig. 2. Due to the conduction in the terpolymer,the static dielectric constant cannot be measured at temperatures above 240 K. The rise ofthe dielectric constant at high temperature (for the static dielectric constant as well as theone at 20 Hz and 100 Hz) is caused by the conduction in the sample. The inset to Fig. 2shows that the characteristic relaxor frequency, determined from peaks in the imaginarydielectric constant ε1,” follows the Vogel-Fulcher (V-F) law f = f0 exp [−U/k(T−T0)]with the Vogel-Fulcher temperature T0 = 254 K for the P(VDF-TrFE-CFE) 68/32/9 mol%terpolymer.

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Figure 2. Temperature dependence of the real, ε′1, and imaginary, ε′′

1 , parts of the complex lineardielectric constant, measured at several different frequencies in the P(VDF-TrFE-CFE) terpolymer.The temperature dependence of the static dielectric constant εs , measured via charge accumulationtechnique, is also shown. The inset shows that the characteristic relaxation time follows the Vogel-Fulcher law.

The temperatrure dependence of the third-order nonlinear dielectric constant ε3 and di-electric nonlinearity a3 (= ε3/ε3

0ε41, where ε0 is the vacuum permittivity) indicates a crossover

from decreasing paraelectric-like to increasing glass-like temperature behavior (Fig. 3).Such a behavior is in accordance with the prediction of the spherical random-bond-random-field model of relaxor ferroelectrics. Although one standard way to analyze the dielectricdata is to fit to the Cole-Cole plot, this procedure can’t provide direct and independentinformation about the actual relaxation spectrum. Furthermore, this method is not suitablewhen the relaxation spectrum becomes extremely polydispersive in the relaxor with de-creasing temperature and when the freezing temperatrure is approached. The informationon the relaxation spectrum and thus on the dynamic processes can be directly extracteedusing the so-called temperature-frequency plot. By varying the reduced dielectric constant

δ = ε′1 − ε∞

εs − ε∞=

∫ z2

z1

g(z)dz1 + (ω/ωa)2 exp(2z)

(1)

between the values 1 and 0 different segments of the relaxation spectrum g(z) are beingprobed. The temperature-frequency plot for the terpolymer P(VDF-TrFE-CFE) 68/32/9

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 181

Figure 3. Temperature dependences of the third-order nonlinear dielectric constant ε3 and dielectricnonlinearity a3, measured at three different frequencies in the P(VDF-TrFE-CFE) terpolymer. Theinset shows paraelectric-to-glass crossover in the temperature dependence of a3.

mol% is shown in Fig. 4(a). The solid lines through the δ = 0.95, . . . , 0.7 data are fits tothe V-F law, which demonstrate the diverging behavior of the relaxation time. The freez-ing temperature of the system is determined by the divergence of the longest relaxationtime, i.e., T f = T0(δ → 1) = 269 K, which is higher than the V-F temperature determinedfrom the characteristic relaxation time in Fig. 2 because the bulk of the relaxation timesremains finite below T f . In fact, the high frequency component of the relaxation spec-trum remains active down to the lowest temperatures. That is, the solid line through δ =0.01 data in Fig. 4(a) is a fit to the Arrhenius law f = f0 exp(−E/kBT). The temperature-frequency plot for the high energy electron irradiated P(VDF-TrFE) 68/32 mol% relaxorpolymer is also shown in Fig. 4(b). The high frequency components (curves with smallδ) remain active down to temperatures much below T f and follow the Arrhenius lawf = f0exp(−E/kBT).

II. Electromechanical Responses of the Electroactive Polymers

One major objective for the development of the relaxor ferroelectric polymers in PVDFbased ferroelectric polymers is to exploit the large electromechanical responses, orig-inating from the large strain associated with the molecular conformation changes be-tween the polar- and non-polar forms. As an illustration, we show in Fig. 5 the latticedimensions for the β-phase, in which the polymer chains adapt an all-trans-conformation,and α-phase, in which the molecular conformation is in trans-gauche form (TGTG′).

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Figure 4. Temperature-frequency plot for several fixed values of the reduced dielectric constant δ

in (a) the P(VDF-TrFE-CFE) terpolymer and (b) the high energy electron-irradiated P(VDF-TrFE)copolymer. Solid lines are fits obtained with Vogel-Fulcher (δ = 0.95, 0.8, . . . , 0.05) or Arrhenius(δ = 0.03, 0.01) expression. Because of high polydispersivity of the relaxation spectrum and largeelectrical conductivity, the temperature-frequency plot for terpolymer could not be determined in therange of 0.01 < δ < 0.70. Insets show the dependence of the Vogel-Fulcher temperature T0 on δ inboth relaxor-like systems.

Apparently, a very large strain (>10% along the polymer chain direction) can be obtainedfor a conformation change between the two forms. However, in the traditional ferroelectricPVDF homopolymer and P(VDF-TrFE) copolymers, no reversible conformation changebetween the polar- and non-polar forms can be realized near room temperature. By con-verting the normal ferroelectric polymer into a relaxor, the defects introduced modify theenergy balance between the polar- phase and non-polar phase and render the non-polarphase to be thermodynamically stable at room temperature. The diffused transformationprocess between the non-polar and polar-molecular bonds results in very little hysteresisin the polarization and strain responses at near the broad dielectric maximum (near roomtemperature). Consequently, large electromechanical responses can be achieved. Presentedin Fig. 6 is the polarization hysteresis loop measured at room temperature for the terpoly-mer and its comparison with the polarization loop measured for the copolymer with thesimilar VDF/TrFE ratio. The relaxor terpolymer shows very little polarization hysteresis.Furthermore, the strain responses of the terpolymer are presented in Fig. 7. For uniax-ially stretched terpolymer, a transverse strain of 5% along the stretching direction wasobtained.

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 183

Figure 5. Schematic views of form I (i.e. β phase) having lattice dimensions of a = 8.58A, b =4.91 A, and a chain direction (or fiber axis) of c = 2.58 A, and form II (i.e. α phase) having latticedimensions of a = 4.96A, b = 9.64 A, and a chain direction (or fiber axis) of c = 4.62 A.

Using the relationship

Sv = S3 + S1 + S2 (2)

where S1 and S2 are the transverse strains, the volume strain can be determined and it wasfound that for the terpolymer of P(VDF-TrFE-CFE), the volume strain is very small (onthe order of 10% of the thickness strain S3). This is quite different from the irradiated

Figure 6. The polarization hysteresis loop of the terpolymer P(VDF-TrFE-CFE) 63/37/7.5 mol%measured at 10 Hz. For comparison, the polarization loop of P(VDF-TrFE) 65/35 copolymer is alsopresented.

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Figure 7. (a) The field-induced longitudinal strain (S3), (b) transverse strain for the uniaxiallystretched terpolymer measured along the stretching direction, and (c) transverse strain for the un-stretched P(VDF-TrFE-CFE) 68/32/9 mol% as a function of the applied field. The measurementfrequency is in the 1 to 10 Hz range.

copolymers, which exhibit high volume strain and suggests different electrostrictive straingeneration mechanisms in these two classes of relaxor ferroelectric polymers [7, 8].

III. Energy Harvesting with the Electroactive Polymers

During the past several decades, there have been many attempts in utilizing the electroactivematerials, especially piezoelectric materials, to harvest electric energy from a wide rangeof acoustic and mechanical energy sources. There are many factors which will influencethe efficiency and output electric energy and power density of an energy harvesting system.In general, an electroactive material with a high electromechanical conversion efficiencyis highly desirable. However, from the device as well as system point of view, the overallsystem efficiency in the energy conversion process can be significantly larger than that ofthe electroactive material by properly designing the electric controlling circuit (referredto as the electric boundary condition in this paper). This point has been largely ignoredin the past investigations on energy harvesting using the electroactive materials. In ad-dition, compared with ceramic materials, polymers in general offer a better mechanicalor acoustic impedance matching with the external mechanical environment, ensuring aneffective energy transfer from the mechanical energy sources to the electroactive materi-als. Moreover, the high elastic energy density (∼1 J/cm3) of the electrostrictive polymersmakes it much easier to achieve high electric energy density for the energy harvestingsystem.

Presented in Fig. 8 is a classical energy harvesting cycle used in the literature to illustratethe energy conversion process in a piezoelectric material [9]. In this energy harvestingcycle, the total input mechanical energy density from the external mechanical source tothe piezoelectric material is W1+W2 and the harvested electric energy density is W1. Thecoupling factor or the energy conversion efficiency, therefore, is

W1

W1 + W2= k2 (3)

Equation (3) illustrates two important factors for an energy harvesting system. In order toachieve high electric energy output from a mechanical source, a high coupling factor of theelectromechanical material is required and a high mechanical energy density available inthe material is also highly desirable (large W1+ W2). In this section, we will illustrate how

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 185

Figure 8. An illustration of a typical energy harvesting cycle for an electroactive material such aspiezoelectric material.

external electrical and mechanical conditions can influence the energy harvesting processwith these electrostrictive polymers, especially the output electric energy density and energyconversion efficiency.

3.1. Analysis of the Performance of the Two Energy Harvesting Cycles

Presented in Fig. 9 is an energy harvesting cycle which can be realized for the electrostrictivepolymers (as well as in general electroactive materials). In this energy harvesting cycle, aconstant electric field EL is applied as the stress is increased to Tmax from the state 1 to thestate 2, followed by an increase in the electric field from EL to EH (from the state 2 to thestate 3), and then a constant field EH as the stress is reduced from Tmax to 0 (from the state3 to the state 4). Finally, at zero stress the electric field is reduced to EL , returning to thestate 1. The harvested energy density can be found as [10]:

W1 = Tmax M(E2

H − E2L

)(4)

Figure 9. Energy harvesting cycle corresponding to the constant electrical field conditions as theelectroactive material is stressed and unstressed.

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Figure 10. Energy harvesting cycle corresponding to the sinusoidal excitations of stress and electricfield.

where M is the electrostrictive coefficient of the material (S=ME2). Since W2 = 12 sT 2

max,where s is the elastic compliance of the material, the coupling factor is given by

k =√

M(E2H − E2

L )12 sTmax + M(E2

H − E2L )

(5)

The maximum harvested energy density and coupling factor occurs when EL is zero andEH is set to be Emax, which is the maximum applied electric field allowed by the materialand system.

W1 max = TmaxME2max (6)

and

kmax =√

ME2max

12 sTmax + ME2

max

(7)

Using a M = 2.2 × 10−18 m2/V2, Emax = 100 MV/m, s = 3.3 × 10−9/Pa, and Tmax = 7.5MPa, a kmax = 0.8 and W1max = 0.16 J/cm3 are derived.

Another energy harvesting cycle which can also be easily realized is that presented inFig. 10 where the mechanical and electrical excitations both consist of constant componentsand sinusoidal-varying components of the same frequency,

E = Eac sin(ωt) + Eb

T = Tac sin(ωt + φ) + Tbias

where Eb is a DC bias electric field, Tbias is the DC stress bias field. Assuming a phasedifference φ between the sinusoidal component of the electric field and mechanical field,the harvested electric energy density in one period is:

W1 =∫ T

0JEdt = 2πMEb EacTac sin φ (8)

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 187

Assuming the electric field and stress are bound between zero and their maximum allowedvalues, the maximum energy density W1 is achieved when Eb = 1

2 Emax, Tac = 12 Tmax and

φ = −π/2, which is

W1 max = −π

4ME2

maxTmax (9)

Using the values as listed in the previous example, W1max = 0.12 J/cm3.

3.2. Experimental Results

A set up was developed to perform the energy harvesting experiment of the electrostrictivepolymer (the irradiated P(VDF-TrFE) copolymer) for the energy harvesting cycle with si-nusoidal stress signal as the input mechanical energy source and a sinusoidal electric fieldwith different phase delays. In the experiment, to provide reliable electrodes for the elec-trostrictive polymer films during the energy harvesting cycle in which the films experiencelarge strain change, a thin polypyrrole conducting polymer electrode was deposited on bothsides of the polymer films. After that, a 20 nm Al layer was evaporated on both surfaces ofthe polymer films to provide high conductivity of the electrodes. A DC bias electric field isapplied to the electrostrictive PVDF films to establish an effective piezoelectric state.

From Eq. (8), it is obvious that when the phase difference φ between the mechanicalstress and electric field is negative, the specific energy is negative, which means that anelectric energy has been harvested during this process. Figure 11 is the electric energydensity vs. the phase difference φ for the electrostrictive polymer under different electricfield. For the experimental condition of Eb = 61 MV/m and Eac = 15 MV/m, a harvestedelectric energy density of 39.4 mJ/cm3 with an conversion efficiency of 10% (correspondingto an coupling coefficient of 0.31) were obtained. Due to the experimental constraints, theapplied electric field and stress field were smaller than the maximum allowed values of thematerial. However, the results obtained here are much better compared with earlier resultsreported. For instance, in reference [11] it was reported that using PVDF piezoelectricpolymers the harvested energy density is 0.044 J/cm3 with an conversion efficiency of0.5%, while using pizeoceramic PZT the energy density was improved to 2.1 J/cm3 withan efficiency in the range of 1.5 to 5% [11].

Figure 11. (a) The harvested power (V DC = 800 V, VAC = 200 V), sample dimension: 23.04 mm× 7.72 mm × 13 μm. (b) Optimum energy harvesting occurs when phase angle between AC voltageand stress is 90◦. (See Color Plate IX)

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Figure 12. (a) FTIR transmission spectra and (b) thickness strain of the terpolymer film underdifferent electric fields. (See Color Plate X)

IV. Large Electro-Optic Response in the Relaxor Terpolymers

The large polarization change in the relaxor ferroelectric polymers also suggests a largeelectrooptic effect (E-O) through the Kerr effect. Although the refractive index of relaxorferroelectric polymers originates solely from electronic dipoles, the large change in themolecular conformation (and hence the electric dipoles) induced by the applied field hasthe potential of leading to large refractive index change.

In order to measure the E-O response, polymer films were coated with conductingpolymer polypyrrole on both sides, resulting in a sandwiched structure. The coated con-ducting layers functioned as both electrodes and reflectors. The optical transmissions of thesandwiched structure under different electric fields were measured by a Fourier TransformInfrared Spectrometer (FTIR) (Fig. 12(a)). The polymer under study was terpolymer ofP(VDF-TrFE-CFE) with a composition of 68/32/9 mol%. The optical path length changeis related to the spectrum shift by [12]

(nd)

nd= λ

λ1(10)

where n, d are the refractive index and thickness of the polymer, λ1 is the peak wavelengthof the transmission, and λ is the shift of peak wavelength. The optical path length changewas a result of both pure refractive index change (n/n) and longitudinal strain (d/d):

(nd)

nd= n

n+ d

d(11)

The thickness strain of the film was measured separately (Fig. 12(b)). Using Eqs. (10) and(11), we obtain that the pure refractive index change under electric field of 80MV/m isn/n = −2.6%. Comparing to inorganic crystals such as LiNbO3where n/n < 0.1%, therefractive index change in the terpolymer is more than one order larger. As the terpolymerhas a central symmetric structure, the E-O response is independent of the polarity of appliedfield. The observed E-O effect in the terpolymer possibly originates from both the pure E-O effect and a secondary effect due to photoelasticity. For example, under a condition ofpartially mechanical clamping (the polymer film was coated on a substrate) the refractive

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 189

Figure 13. Properties of nanocomposites (a) Refractive index of nanocomposite as a function of theconcentration of ZnS (b) UV-Vis absorption of nanocomposite (10.6 vol% ZnS) (c) FTIR spectrumof nanocomposite film (10.6 vol% ZnS) under different electric fields. (See Color Plate XI)

index change was 0.52% at 100 MV/m [13]. The difference in E-O response may be due tothe much smaller contribution of photoelasticity in clamped films.

The large refractive index change in the terpolymer opens up new possibilities to tun-able optical devices. For example, tunable long-period fiber gratings (LPGs) incorporatingthe terpolymer as active layer achieved a resonance wavelength shift of 18 nm under electricfield of 50 MV/m, which is much larger than any other reported result of E-O tuning inLPGs. In tunable optical devices the tuning range usually strongly depends on the indexmatching between the active E-O layer and the passive layers [14]. In tunable LPGs, tuningrange under a matched condition can be more than twice as large as in unmatched conditionwith same amount of refractive index change in E-O material [15]. Therefore, it is highlydesired that the native refractive index of E-O material can be varied in a wide range so asto match various passive materials in different tunable devices. At the same time large E-Oresponse and optical transparency must be maintained. In order to meet these requirements,we adopt a nanocomposite approach where zinc sulfide (ZnS) nanoparticles were embed-ded in terpolymer matrix [16]. ZnS has high refractive index (n = 2.27@1.55 μm) and istransparent over wide wavelength range (0.4–20 μm). Due to their large refractive indexdifference (terpolymer n ≈ 1.4 @1.55 μm), only small amount of ZnS is required to change

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the refractive index of the polymer matrix. Therefore, most volume of the nanocompositeis occupied by the terpolymer and large E-O response is likely to be retained. Furthermore,the size of nanoparticles is much smaller than light wavelength and hence light scatteringshould be negligible. One of the difficulties in nanocomposite preparation is that due totheir large surface energy, the nanoparticles tend to agglomerate to form large particlesand induce serious scattering. We introduced thiophenol and methyl 3-mercaptopropionateas surfactants in the synthesis of nanoparticles to passivate the surface of the nanopar-ticles and possibly improve the compatibility between the nanoparticles and polymermatrix.

The refractive indexes of the nanocomposites with different concentration of nanopar-ticles are measured by ellipsometer (Fig. 13(a)). Refractive index can be readily increasedby 0.1 with about 10 vol% of ZnS. The optical loss of the nanocomposite is similar to neatterpolymer in visible and near IR wavelengths (Fig. 13(b)). The large UV loss is possiblydue to intrinsic loss of ZnS. Under electric field of 45 MV/m, nanocomposite film (10 vol%ZnS) free from mechanical clamping exhibited an optical path length change of −1.7% (Fig.13(c)), which has a 37% decrease comparing to neat terpolymers. If we assume that the ratiobetween contributions of pure index change and strain remains the same, the pure indexchange is estimated to be −1.3%. Tunable LPG was fabricated using the nanocomposite asactive layer. Resonance wavelength shift of more than 50nm was achieved under electricfield of 50 MV/m, which is more than enough to cover an entire fiber optic communicationband. Comparing to the spectrum of LPG immersed in different index-matching fluids, therefractive index change under such a partially clamped condition (the length of fiber isfixed) is about 1%.

V. Conclusion and Acknowledgement

The dynamics of the relaxor ferroelectric polymers, i.e., the P(VDF-TrFE-CFE) terpolymersand high energy electron irradiated copolymers has been studied. Utilizing the temperature-frequency plot analysis of the linear dielectric constant, it was shown that the relaxor fer-roelectric polymers display typical glassy asymmetric behavior of the relaxation spectrum:while the ergodicity of the system is effectively broken due to divergency of the longestrelaxation time at the freezing temperature T f (∼270 K), the bulk of the relaxation timesremain active below T f , the high frequency part, obeying the Arrhenius law, even to thelowest temperatures. Also the dielectric nonlinearity undergoes a crossover from decreas-ing paraelectric-like to increasing glass-like temperature behavior when approaching thefreezing transition from above.

Owing to the reversible transformation between the polar- and non-polar nanoregions, agiant electrostrictive response was observed in the relaxor ferroelectric polymers. The largeelectrostrictive strain and elastic energy density, as well as improved electromechanicalcoupling factor make them attractive candidates to be used for the energy harvesting. Weexamined how to further improve the harvested electric energy density and system efficiencyby exploiting an active electronic circuit to vary the electric boundary conditions imposedon the electroactive polymers. We showed that the electromechanical coupling factor of thesystem can be significantly higher than that of the materials by employing an appropriateelectric boundary conditions. It was demonstrated that a relatively high harvested electricenergy density (∼40 mJ/cm3) and efficiency (∼10% or 31% of coupling factor) can beobtained from the electrostrictive P(VDF-TrFE) polymers under a sinusoidal electric fieldapplied during the energy harvesting cycle on the polymer.

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Relaxor Ferroelectric Polymers–Fundamentals and Applications 191

The large reversible polarization change, reflecting the large changes in the molecularconformations, results in a large E-O effect (Kerr effect) in which n/n ∼ −2.6% can beobtained in the relaxor ferroelectric terpolymer. Furthermore, making use of the nanoparti-cles of ZnS which has a reflective index 2.27, much higher than that of P(VDF-TrFE-CFE),we demonstrated that the reflective index of the polymer composite can be varied over alarge range while maintaining the high E-O effect of the polymer matrix, which is attractivefor many optical devices.

This work was supported by ONR under Grant No. N00014-05-1-0455 and No.N00014-05-1-0541. We would like to thank Dr. D. Y. Jeong and Dr. Heath Hofmann fortheir assistance in these investigations.

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