J. Phys. Chem. B 2008, 112, 10411–10416
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Structural Dependence of Phase Transition and Dielectric Relaxation in Ferroelectric Poly(vinylidene fluoride-chlorotrifluoroethylene-trifluoroethylene)s Yingying Lu,† Jason Claude,† Luis Enrique Norena-Franco,‡ and Qing Wang*,† Department of Materials Science and Engineering, The PennsylVania State UniVersity, UniVersity Park, PennsylVania 16802, Department of Basic Science, Metropolitan Autonomous UniVersity, Azcapotzalco, Mexico City 02200, Mexico ReceiVed: March 19, 2008; ReVised Manuscript ReceiVed: June 12, 2008
A series of ferroelectric poly(vinylidene fluoride-chlorotrifluoroethylene-trifluoroethylene)s, P(VDF-CTFETrFE), with systematically varied chemical compositions have been synthesized via a two-step approach consisting of copolymerization and dechlorination. The effect of polymer structure on polarization responses and dielectric properties has been investigated over a broad frequency and temperature range. As shown in the X-ray diffraction patterns, multiple phases coexist within the terpolymers as a result of the gauche conformation induced by the CTFE unit. The polarization hysteresis loops reveal the variation of remanent polarization and coercive electric field with the CTFE content due to the changes of crystallinity and crystalline phase. The observed broad dielectric constant peak with Vogel-Fulcher dielectric dispersion behavior suggests a transformation from a normal ferroelectric to a ferroelectric relaxor of the polymers. The relationship between the local relaxation process and relaxor ferroelectric behavior has been examined on the basis of the dielectric and mechanical loss tangents as a function of temperature. Introduction Ferroelectric polymers have raised considerable attention due to their potential applications in a wealth of advanced technologies, ranging from transducers, actuators, sensors, and nonvolatile memories, to capacitors.1–5 Compared to conventional electroactive ceramics, ferroelectric polymers enjoy the inherent advantages such as lightweight, flexibility, easy of processing into large areas, and great opportunities for structural modification. The most well-known and extensively studied family of ferroelectric polymers includes vinylidene fluoride (VDF) based homo- and copolymers.6 The interesting electronic properties of poly(vinylidene fluoride) (PVDF) originate from the presence of highly electronegative fluorine on the polymer chain and the spontaneous alignment of the C-F dipoles in the crystalline phases. PVDF with different chain conformations can be packed into various crystal lattices to form a range of crystalline phases depending on the processing conditions.7 Among them, the R phase is the most common for PVDF crystallized from the melt. The polymer chains with trans-gauche (TG) conformation in the R phase are stacked with their respective polarization in alternating directions, resulting in a nonpolar unit cell and paraelectric behavior. The β phase has all-trans (Tm>4) zigzag chain conformation and parallel dipole moments, and thus is the most polar one with large spontaneous polarization and corresponding ferroelectricity and piezoelectricity. The switch from the nonpolar to polar phases can be induced by electric field poling or mechanical stretching.8 It is also noted that, during the phase transition, large differences in dipole moments and lattice parameters exist among different phases, leading to * To whom correspondence should be addressed. E-mail: wang@ matse.psu.edu. † Department of Materials Science and Engineering, The Pennsylvania State University. ‡ Department of Basic Science, Metropolitan Autonomous University.
significant changes of dielectric constants, piezoelectric and pyroelectric coefficients, and interesting electromechanical responses. The energy barrier of the phase transition, however, is considerably high.9 As a result, PVDF generally displays low electric field sensitivity in terms of its dielectric constant, piezoelectric coefficient, and electromechanical coupling efficiency. It has long been realized that the ferroelectric phase of PVDF can be altered by certain disorders or defects in its molecular structure, which in turn mitigates the activation energy of the phase transition. For example, by incorporating trifluoroethylene (TrFE) into PVDF, the resulting P(VDF-TrFE) copolymers adopts an all-trans chain conformation due to the presence of extra fluorine atoms in TrFE that prevents the polymer chain from accommodating the trans-gauche conformation.10–14 Therefore, the copolymers are able to crystallize directly at room temperature into the ferroelectric β phase that possesses polar unit cells. The ferroelectricity in the copolymer is evidenced by the existence of the ferroelectric-paraelectric (F-P) phase transition at a temperature below the melting point.15 It has also been demonstrated that electron or gamma irradiation can introduce structural defects into the β phase of P(VDF-TrFE), which stabilizes the nonpolar phase at room temperature, and reduces or even eliminates the transition barrier between the polar and nonpolar phases.16–18 Consequently, at room temperature the phase transformation can be induced by an applied electric field with very little hysteresis. Interestingly, the irradiated copolymers display an electrostrictive strain greater than 4% and an elastic energy density higher than 1 J/cm,3 which are orders of magnitude larger than those of untreated P(VDFTrFE).19 Similar effects have been recently achieved by introducing a small amount of bulky ter-monomer units, such as chlorotrifluoroethylene (CTFE), 1,1-chlorofluoroethylene (CFE) and hexafluoropropylene (HFP), as crystalline defects into P(VDF-TrFE).20–22 The resulting terpolymers exhibit high dielectric constant and large electrostrictive response at room
10.1021/jp802413g CCC: $40.75 2008 American Chemical Society Published on Web 08/06/2008
10412 J. Phys. Chem. B, Vol. 112, No. 34, 2008 temperature, making them the best piezoelectric polymers. In this paper, we investigate the phase transition and relaxation behavior of P(VDF-TrFE-CTFE) terpolymers prepared via a newly developed synthetic approach. Unlike the conventional ter-polymerization method, this two-step synthetic route leads to the ferroelectric polymers with exquisitely controlled compositions.23 By varying the polymer structures, the content of different phases can thus be modified, which allows a systematical study of the microstructures of the polymers and their influence on the polarization dynamics and dielectric properties.
Lu et al. SCHEME 1: Synthesis of the Ferroelectric Polymers with Varied Chemical Compositions
Experimental Section Materials. Unless otherwise noted, all solvents and reagents were purchased from VWR and used as received. Ethyl hydroxyethyl cellulose was used as received from Polymer Source Inc. Tetrahydrofuran (THF) was distilled from sodium benzophenone ketyl under nitrogen. Vinylidene fluoride and chlorotrifluoroethylene were purchased from SynQuest Laboratory Inc. and purified by the freeze-thaw process prior to use. Synthesis of P(VDF-CTFE-TrFE). A 300-mL stainless steel autoclave equipped with a mechanical stirrer was charged with hydrogen peroxide (0.27 g, 2.1 mmol), ethyl hydroxyethyl cellulose (0.16 g), diethyl carbonate (0.05 mL, 0.4 mmol) and degassed distilled water (150 mL). The reactor was degassed by repeated freeze-thaw cycles and then cooled by liquid nitrogen. After VDF (33 g) and CTFE (15 g) were condensed and transferred into the reactor, the reaction mixture was heated at 110 °C for 4 h. P(VDF-CTFE) (20 g) was obtained by filtration and purification by washing with methanol and distilled water, and drying in vacuum at 60 °C for 24 h. The hydrogenation of P(VDF-CTFE) to yield P(VDF-CTFE-TrFE) was performed according to literature procedures.23 Preparation of Polymer Films. Polymer films with thickness around 20 µm were prepared by solution casting from DMF onto glass slides. The polymer solutions were filtered through 0.45 µm poly(tetrafluoroethylene) filters before casting to remove any particulates. The films were dried at 60 °C overnight followed by heating at the same temperature under vacuum for 24 h. Free-standing films were created by soaking coated glass slides in distilled water and carefully peeling off the films. Gold electrodes with thicknesses of ∼60 nm were sputtered onto the films for electrical characterization. Characterization. 1H and 19F NMR spectra were recorded on a Bruker AM-300 spectrometer instrument. The molecular weights of the polymers were characterized with a Viscotek gel permeation chromatography (GPC) system equipped with light scattering, refractive index, UV/vis and viscometer detectors in a DMF mobile phase. The thermal transition data were obtained by TA Instruments Q100 differential scanning calorimeter (DSC) at a heating rate of 10 °C/min. Wide-angle X-ray diffraction (WAXD) studies were conducted using a Scintag Cu-KR diffractometer with an X-ray wavelength of 1.54 Å. Polarization hysteresis loops at room temperature were collected using a modified Sawyer-Tower circuit at 10 Hz. Dielectric properties were acquired using an Agilent LCR meter (E4980A) at 1.5 V. Isothermal dielectric spectra *(f, T) were collected using a Novocontrol GmbH Concept 40 broadband spectrometer in the frequency domain (10 MHz to 0.01 Hz). The temperature was controlled by a Novocontrol Quatro Cryosystem with stability better than (0.1 °C. Data collection did not start until the temperature had been stabilized at least 10 min. Dynamic mechanical thermal analysis (DMTA) was performed with a TA DMA 2980 instrument from 20 to 1 Hz.
Results and Discussion Synthesis. As outlined in Scheme 1, the synthetic strategy involves the copolymerization of VDF and CTFE followed by a reductive dechlorination reaction. P(VDF-CTFE)s were prepared via a suspension polymerization of VDF and CTFE with water as the continuous phase, hydrogen peroxide as the initiator, ethyl hydroxyethyl cellulose as the surfactant, and diethyl carbonate as the chain transfer agent at 110 °C.24,25 Since VDF and CTFE possess similar reactivity ratios in free radical polymerization,26 the chemical compositions of the P(VDFCTFE) copolymers are readily controlled by varying the monomer feeding ratios. Reduction of chlorine in P(VDF-CTFE) was achieved upon treatment with AIBN and tri(n-butyl)tin hydride to yield P(VDF-CTFE-TrFE).27 Scheme 1 lists a series of the prepared P(VDF-CTFE-TrFE)s containing 73.6 mol % VDF and a systematic variation of CTFE and TrFE contents. The chemical compositions of the polymers were calculated according to the integrals of the characteristic peaks in 1H and 19F NMR spectra.28 Polymers 1-6 are soluble in common organic solvents such as THF, DMF, and DMSO. The absolute weightaverage molecular weights of the polymers, determined by GPC equipped with light scattering detectors in DMF are approximately 240 kDa with polydispersities of ∼3.40. Crystalline Structure. The evolution of the crystalline structures of the polymers with chemical compositions was examined in wide-angle X-ray diffraction (WAXD) measurements. As shown in Figure 1, polymers 1 and 2 display a diffraction peak at a 2θ angle of ca. 17° arising from the (100) reflection of the nonpolar R phase. As the CTFE content is decreased via increasing dechlorination ratios, two peaks were observed in X-ray diffraction patterns: one is associated with the R phase and the other one corresponds to the (002) or (020)
Figure 1. X-ray diffraction patterns of the polymers.
Structural Dependence of Phase Transition
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Figure 2. Lattice spacing and degree of crystallinity as a function of the CTFE content in the polymers.
diffractions from the γ phase, indicating the coexistence of the R and γ phases in polymers 3, 4, and 5. With a complete reduction of CTFE into TrFE, the crystal phase is finally transformed into the β phase in polymer 6, as evidenced by the appearance of the characteristic (200, 110) diffraction peak at 20°. Concomitant with the change of the crystal phases, the lattice spacing (d) decreases steadily from 5.21 Å (polymer 1) to 4.40 Å (polymer 6) due to the substitution of the bulky chlorine atoms with hydrogen. As also summarized in Figure 2, the degree of crystallinity (χc), assessed from the area of the diffraction peaks from the crystalline phases and amorphous halo, increases gradually from 7.2% for polymer 1 to 35.4% for polymer 6. These results agree with earlier studies, indicating that the CTFE component acts as defects to introduce gauche state into the polymer and thus alters the crystalline phases from the β phase with all-trans chain conformation to the γ form with trans-gauche (T3G) conformation.20,29 Polarization Responses. Typical polarization hysteresis loops of the polymers measured at room temperature under 10 Hz and 150 MV/m are given in Figure 3a, where the polarization hysteresis is reduced markedly from polymers 6 to 4 and further to 3. The coercive field (Ec) of polymer 6 is about 33.3 MV/m, whereas the Ecs of polymers 4 and 3 are 19.3 and 9.4 MV/m, respectively. Compared to the P(VDF-TrFE) synthesized via ter-polymerization, which normally shows a well-defined square polarization hysteresis loop with a Ec of 45 MV/m and a remanent polarization of 60 mC/m2,9 polymer 6 exhibits a much slimmer polarization loop with a reduced Ec. This discrepancy can be explained by the difference in their microstructures, in which TrFE is incorporated as both head-to-tail and tail-to-tail sequences in the P(VDF-TrFE) prepared from direct ter-polymerization.28 On the other hand, the 19F NMR spectrum (Figure 4) of polymer 6 yielded from the complete reduction of P(VDFCTFE)s shows the major chemical shifts at 93 ppm (-CHF-CH2-CF2-CH2-CF2-), 112-114 ppm (-CHFCH2-CF2-CH2-CHF-), 122 ppm (-CH2-CHF-CF2CHF-CF2-), 130 ppm (-CH2-CF2-CF2-CHF-CH2-) and 199 ppm (-CF2-CF2-CHF-CH2-CF2-), suggesting that TrFE is present dominantly as a VDF-TrFE tail-to-tail defect structure. A higher percentage of regio-defects in polymer 6 thereby gives rise to a lower crystallinity and a high mobility of dipoles and thus a reduced Ec. It is also interesting to note that the ferroelectric hysteresis in polymer 6 is a double loop, which is presumably caused by imperfection in crystals and the superposition of two loops with
Figure 3. ( a) Polarization hysteresis loop measured at room temperature under 150 MV/m for the polymers 3, 4 and 6. (b) Remanent polarization (Pr) and maximum polarization (Ps) versus the CTFE content for the ferroelectric polymers.
different coercive fields.14 This is consistent with the observation in X-ray diffraction studies where polymer 6 exhibits a broad diffraction peak. The change of the remanent polarization (Pr) and the maximum polarization (Ps) as a function of the CTFE concentration is summarized in Figure 3b. Both Pr and Ps decrease rapidly as the CTFE content is increased to 17.2 mol % (polymer 3). Apparently, CTFE breaks large polar domains into small ones and makes the dipoles reverse at lower electric field and also reduces the polarization level because of a reduction in crystallinity and ferroelectricity. Interestingly, above 17.2 mol % CTFE content, there are increased Pr and Ps, a trend related to the change of crystalline phases from the γ phase in polymer 3 to the R phase in polymer 2. Compared to the γ phase with T3G chain conformation, a larger energy barrier exists between the R phase with TG chain conformation and all-trans conformation yielded from the orientation under an external electric field. For instance, the crystal cohesive energies of Tm>4, T3G and TG crystalline phases of P(VDF-TrFE) are -0.86, -1.23, and 2.09 kcal/mol per carbon atom,30 which suggests that transformation between the R phase with TG conformation and the β phases with Tm>4 conformation are high-energy consumption and accordingly leads to large polarization hysteresis and high remanent polarization. Dielectric Properties. Figure 5 presents the frequency dependence of dielectric permittivity and loss tangent of the polymers measured at 25 °C. All of the polymers show the same
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Figure 4.
19
Lu et al.
F NMR spectrum of polymer 6.
Figure 5. Dielectric permittivity (′) and loss tangent (tan δ) of the polymers.
general trend of decreasing permittivity with increasing frequency which is related to the reduction of the dipolar contribution at high frequencies. Even with the decreasing trend, the permittivity remains relatively flat through 100 kHz before decreasing significantly. The polymers with higher permittivities have a sharper drop-off at high frequencies because of the larger contribution of dipole polarization to the overall permittivity. At fixed frequencies, polymer 3 shows a maximum in permittivity of ∼24 at 1 kHz. The corresponding loss tangents of the polymers remain flat between 100 Hz and 10 kHz with increases below and above this range. The sharp increase above 10 kHz is related to the dipole relaxation where the slope of the line is greater for the higher permittivity polymers. The increase at low frequencies is most likely due to polymer chain motion. As the crystallinity of the polymers decreases, the chains become more mobile at lower frequencies resulting in increased losses from polymers 6 to 2. The dielectric properties of the polymers have also been investigated at varied temperatures. As shown in Figure 6, the dielectric peaks become broadened and appear at lower temperatures in polymer 4 in comparison to those in polymer 6. Polymer 4 exhibits typical ferroelectric relaxor behavior with the dielectric peak shifting progressively to high temperature as frequency increases,31 while the dielectric peak position of polymer 6 shows much weaker frequency dependence. These results are consistent with earlier studies, indicating that the incorporation of CTFE components leads to the conversion to
Figure 6. Temperature dependence of dielectric constants of polymers 4 and 6 at varied frequency.
a ferroelectric relaxor from a normal ferroelectric.32,33 It should be noted that the broad dielectric peaks in the terpolymers do not correspond to the F-P transition, but rather, they are consequence of the kinetics associated with the dipolar glass freezing transition.34 The frequency dispersion of the dielectric behavior of the polymers was analyzed using the Vogel-Fulcher (V-F) equation: f ) f0 exp(-T0/Tm - Tf), an empirical relation that holds for many spin-glass systems and relaxor ferroelectrics, where T0 is a constant, f0 is associated with the relaxation frequency of local dipole regions, Tm is the peak temperature of dielectric constant measured at frequency f, and Tf is regarded as the freezing temperature.35,36 At temperatures below Tf, a ferroelectric state can be induced by a high electric field, while the remanent polarization increases gradually as the temperature is lowered. As shown in Figure 7a, the experimental results are well fitted by the V-F relation, indicating the presence of local polar-ordered regions and their coupling in the amorphous matrix, analogous to the features observed in relaxor ferroelectric ceramics and in many glass systems.37 The parameters obtained by the fitting the V-F expression to the data are given in Figure 7b, where f0 increases and Tf decreases with the increase of CTFE content up to 17.2 mol % (polymer 3). The results imply a decrease in the micropolar region size and an increase in
Structural Dependence of Phase Transition
Figure 7. (a) Measured (solid symbol) relation between the frequency and dielectric constant peak temperature for the polymers. Solid lines are fits obtained with the Vogel-Fulcher expression. (b) Fitting parameters in the V-F equation as a function of the CTFE content in the polymers.
amorphous area induced by the incorporation of CTFE as defects, consistent with the changes of the crystalline structures from the β to γ phases as revealed in XRD studies. The increase of f0 also indicates that the dielectric behavior depends more strongly on the frequency as the CTFE content increases, a distinct feature for the polymers shifting from normal ferroelectrics to ferroelectric relaxors.38,39 As the crystalline structure changes from mixed R and γ forms to the R phase with the further increase of CTFE content to 22.5 mol % (polymer 2), both the reduction of f0 and the upward shifting of Tf suggest that the coupling of the local polar regions becomes stronger and the average polar ordering in these regions increases, which are coincident with the polarization results presented in Figure 3b. Relaxation Behavior. It is known that, due to their semicrystalline nature, PVDF based ferroelectric polymers exhibit rich relaxation processes depending on the molecular relaxations in amorphous and crystalline domains, and at the crystalamorphous interface.40–42 As shown in Figure 8a, the dielectric loss spectra of polymers 2, 3, 4, and 6 exhibit two strong relaxation processes in the range of -50 to +50 °C, while polymer 5 shows no obvious relaxation peak at 100 Hz. For polymer 6, the low-temperature relaxation peak at -16 °C is often denoted as βa, which is usually attributed to the dynamic glass transition of segments in the amorphous phase.43 As established in more recent studies, the βa relaxation also includes a significant contribution from domain wall and polar-defect
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Figure 8. (a) Dielectric loss spectra of the polymers measured at 100 Hz. (b) DMTA loss spectra of the polymers measured at 1 Hz.
motion in the crystalline phase.44,45 The high-temperature relaxation referred to as βc locates at 27 °C. Although the nature of the βc process is still under discussion, it is generally believed that the primary contribution comes from the relaxation of chains at the amorphous/crystalline interface.13 For polymers 2, 3 and 4 with less ferroelectric phase and lower degrees of crystallinity, a low-temperature relaxation (denoted as βr) emerges at about -33 °C, which is associated with the dynamic manifestation of the glass transition in the amorphous phase of the polymers. The high-temperature relaxation βc process shifts from -9.8 °C for polymer 2 to 22.3 °C for polymer 4 as the crystallinity increases and is rather related to the reorientation dynamics of the crystalline domains.39 To further establish the relationship between the relaxation process and polymer structures, dynamic mechanical thermal analysis (DMTA) has been employed to investigate molecular motions in the polymers. Different from dielectric analysis which reveals the relaxation related to motions of dipoles, DMTA is sensitive to molecular motions associated with mechanical relaxation. As shown in Figure 8b, polymers 5 and 6 exhibit two relaxation peaks in the mechanical loss (tan δ) spectra. Different from the dielectric spectra, the βc process at around 22 °C is more prominent than the βa relaxation at a lower temperature of about -27 °C in DMTA spectra. The βa process is significantly weakened in polymer 5 as the crystallinity decreases. For polymers 2, 3 and 4, as the macroscopic polar domains are replaced by the nanoclusters in the conversion from the normal ferroelectric to the ferroelectric relaxor and the
10416 J. Phys. Chem. B, Vol. 112, No. 34, 2008 fraction of amorphous phase increases greatly with high concentrations of CTFE in the polymer, it is conceivable that the βa and βc relaxations merge into a new broad relaxation peak with the contributions from glass transition of amorphous phase, the segmental motion and reorientation process in the crystalline phase. Conclusion Ferroelectric P(VDF-CTFE-TrFE) containing 73.6 mol % VDF and a systematic variation of the CTFE and TrFE contents has been prepared via a copolymerization of VDF and CTFE and a subsequent dechlorination reaction. Different from the conventional ter-polymerization, this approach leads to the polymers with accurately controlled compositions, and reduced remanent polarizations and coercive fields. It has been found that the changes in the crystalline structures, phase transition properties, polarization hysteresis, and relaxation processes can be effectively modulated by the concentration of polarization defects in the polymers. The defects induced by the CTFE units break up coherent polarization domain (all-trans conformation) and randomize the inter- and intrachain polar coupling via a trans-gauche conformation, which is responsible for the evolution of ferroelectric properties from a normal ferroelectric into a relaxor ferroelectric phase. Correspondingly, the terpolymers exhibit slim polarization hysteresis loops, diffused dielectric peaks, and strong frequency dispersion which can be fitted well by the Vogel-Fulcher relationship. Acknowledgment. This work was support by the National Science Foundation (CAREER DMR-0548146) and the Office of Naval Research. References and Notes (1) Herbert, J. M., Glass, A. M., Wang, T. T., Eds. The Application of Ferroelectric Polymers; Chapman & Hall: New York, 1988. (2) Naber, R.; Tanase, C.; Blom, P.; Gelinck, G. H.; Marsman, A. W.; Touwslager, F. J.; Setayesh, S.; de Leeuw, D. M. Nat. Mater. 2005, 4, 243. (3) Mu¨ller, K.; Paloumpa, I.; Henkel, K.; Schmeisser, D. J. Appl. Phys. 2005, 98, 056104. (4) Stadlober, B.; Zirkl, M.; Beutl, M.; Leising, G. Appl. Phys. Lett. 2005, 86, 242902. (5) Chu, B.; Zhou, X.; Ren, K.; Neese, B.; Lin, M.; Wang, Q.; Bauer, F.; Zhang, Q. M. Science 2006, 313, 334. (6) Nalwa, H. S., Ed. Ferroelectric Polymers; Marcel Dekker: New York, 1995. (7) Lovinger, A. J. Science 1983, 220, 1115. (8) Furukawa, T. Phase Transitions 1983, 18, 14. (9) Furukawa, T. AdV. Colloid Interface Sci. 1997, 71-72, 183.
Lu et al. (10) Davis, G. T.; Furukawa, T.; Lovinger, A. J.; Broadhurst, M. G. Macromolecules 1982, 15, 329. (11) Higashihata, Y.; Sako, J.; Yagi, T. Ferroelectrics 1981, 32, 85. (12) Tajitsu, Y.; Chiba, A.; Furukawa, T.; Date, M.; Fukada, E. Appl. Phys. Lett. 1980, 36, 286. (13) Yagi, T.; Tatamoto, M.; Sako, J. Polym. J. 1980, 12, 209. (14) Yamada, T.; Ueda, T.; Kitayama, T. J. Appl. Phys. 1981, 52, 948. (15) Furukawa, T.; Date, M.; Fukada, E.; Tajitsu, Y.; Chiba, A. Jpn. J. Appl. Phys 1980, 19, L109. (16) Lovinger, A. J. Macromolecules 1985, 18, 910. (17) Daudin, B.; Dubus, M. J. Appl. Phys. 1987, 62, 994. (18) Odajima, A.; Takase, Y.; Ishibashi, T.; Yuasa, K. Jpn. J. Appl. Phys 1985, 24, 881. (19) Zhang, Q. M.; Bharti, V.; Zhao, X. Science 1998, 280, 2102. (20) Xu, H.; Cheng, Z.; Olson, D.; Mai, T.; Zhang, Q. M.; Kavarnos, G. Appl. Phys. Lett. 2001, 78, 2360. (21) Xia, F.; Cheng, Z.; Xu, H.; Li, H.; Zhang, Q. M.; Kavarnos, G. J.; Ting, R. Y.; Abdel-Sadek, G.; Belfield, K. D. AdV. Mater. 2002, 14, 1574. (22) Xu, H.; Shen, D.; Zhang, Q. M. Polymer 2007, 48, 2124. (23) Lu, Y.; Claude, J.; Neese, B.; Zhang, Q. M.; Wang, Q. J. Am. Chem. Soc. 2006, 128, 8120. (24) Claude, J.; Lu, Y.; Li, K.; Wang, Q. Chem. Mater. 2008, 20, 2078. (25) Claude, J.; Lu, Y.; Wang, Q. Appl. Phys. Lett. 2007, 91, 212904. (26) Yagi, T.; Tatemoto, M. Polym. J. 1979, 11, 429. (27) Cais, R. E.; Kometani, J. M. Macromolecules 1985, 18, 1354. (28) Lu, Y.; Claude, J.; Zhang, Q. M.; Wang, Q. Macromolecules 2006, 39, 6962. (29) Klein, R. J.; Runt, J.; Zhang, Q. M. Macromolecules 2003, 36, 7220. (30) Su, H.; Strachan, A.; Goddard, W. A. Phys. ReV. B 2004, 70, 064101. (31) Cross, L. E. Ferroelectrics 1994, 151, 305. (32) Klein, R. J.; Xia, J.; Zhang, Q. M.; Bauer, F. J. Appl. Phys. 2005, 97, 094105. (33) Zhang, S.; Huang, C.; Cheng, Z.; Xia, F.; Klein, R.; Zhang, Q. M.; Bobnar, V.; Levstik, A.; Bauer, F. Ferroelectrics 2006, 339, 37. (34) Pirc, R.; Blinc, R.; Kutnjak, Z. Phys. ReV. B 2002, 65, 214101. (35) Viehland, D.; Jang, S. J.; Cross, L. E. J. Appl. Phys. 1990, 68, 2916. (36) Tagantsev, A. K. Phys. ReV. Lett. 1994, 72, 1100. (37) Cheng, Z.; Katiyar, R. S.; Yao, X.; Bhalla, A. Phys. ReV. B 1998, 57, 8166. (38) Lau, S. T.; Chan, H. L. W.; Choy, C. L. Appl. Phys. A: Mater. Sci. Process. 2005, 80, 289. (39) Bobnar, V.; Vodopivec, B.; Levstik, A.; Kosec, M.; Hilczer, B.; Zhang, Q. M. Macromolecules 2003, 36, 4436. (40) Sasabe, H.; Saito, S.; Asahina, M.; Kakutani, H. J. Polym. Sci. Polym. Phys. 1969, 7, 1405. (41) Kassem, M. E.; Issa, M. R.; Higazy, A. A. J. Polym. Mater. 1993, 10, 239. (42) Ang, C.; Yu, Z. AdV. Mater. 2004, 16, 979. (43) Ishida, Y.; Saito, S.; Asabina, M.; Kakutani, H. J. Polym. Sci. Part A2 1969, 7, 1405. (44) Omote, K.; Ohigashi, H.; Koga, K. J. Appl. Phys. 1997, 81, 2760. (45) Zhang, S.; Klein, R. J.; Ren, K.; Chu, B.; Zhang, X.; Runt, J.; Zhang, Q. M. J. Mater. Sci. 2006, 41, 271.
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