Article pubs.acs.org/Macromolecules
Crystal Orientation and Temperature Effects on Double Hysteresis Loop Behavior in a Poly(vinylidene fluoride-co-trifluoroethylene-cochlorotrifluoroethylene)-graf t-Polystyrene Graft Copolymer Lianyun Yang,† Elshad Allahyarov,†,‡,§ Fangxiao Guan,∥ and Lei Zhu*,† †
Department of Macromolecular Science and Engineering, Case Western Reserve University, Cleveland, Ohio 44106-7202, United States ‡ Institut für Theoretische Physik, Heinrich-Heine-Universität Düsseldorf, D-40225 Düsseldorf, Germany § Theoretical Department, Joint Institute for High Temperatures, Russian Academy of Sciences, Izhorskaya 13/19, 117419 Moscow, Russia ∥ Polymer Program, Institute of Materials Science and Department of Chemical, Materials and Biomolecular Engineering, University of Connecticut, Storrs, Connecticut 06269-3136, United States S Supporting Information *
ABSTRACT: Recently, double hysteresis loop (DHL) behavior, which is advantageous for the high energy density and low loss dielectric application, was achieved in a poly(vinylidene fluoride-co-trifluoroethylene-co-chlorotrifluoroethylene)-graf tpolystyrene [P(VDF-TrFE-CTFE)-g-PS(14%)] graft copolymer due to the nanoconfinement effect. In this work, we continued to investigate the crystal orientation and temperature effects on the DHL behavior of this graft copolymer. Based on the electric displacement−electric field (D−E) study, crystal orientation had a profound effect on its electrical behavior. For the nonoriented sample, dielectric instead of ferroelectric behavior was observed. After uniaxial stretching, DHLs gradually developed in the oriented films upon increasing the extension ratio. For a fully stretched film, the DHL behavior was stable below 75 °C but gradually disappeared above 100 °C due to enhanced dc conduction and impurity ion migrational loss at elevated temperatures. After subtracting the dc conduction, D−E hysteresis loops from the ion loss were determined for the poling cycles below 100 MV/m. The hysteresis loss from ion migration under an applied field was closely related to ion concentration and diffusion coefficient, which were determined by broadband dielectric spectroscopy. Both parameters were used in a theoretical calculation to obtain hysteresis loops from ion migrational loss. By fitting the theoretical loops with those after dc conduction subtraction, ion mobility was found to be dependent upon both poling field and temperature. This study provides a quantitative understanding of the effects of impurity ions and dc conduction on dielectric and ferroelectric properties of polymers at elevated temperatures.
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INTRODUCTION Polymer dielectrics with high dielectric constant but low loss are excellent candidates for electric energy storage (e.g., plastics)1−3 and electroactive actuation (e.g., elastomers)4,5 applications. Instead of focusing on enhancement of dielectric constant, which has no clear physical meaning at the atomic and molecular scales, our recent research has been focused on polarization of polymers.1 Generally speaking, a high polarization will result in a high dielectric constant for insulating materials. Nonetheless, a high polarization is often accompanied by high dielectric and hysteresis losses.1 Therefore, the challenge lies in the minimization of dielectric and hysteresis © 2013 American Chemical Society
losses when increasing polarization in polymers. This requires fundamental understanding of different types of polarization mechanisms in dielectric materials. In general, five types of polarization exist in dielectric materials.1 Electronic and atomic polarizations, which relax around infrared and optical frequencies, are ubiquitous. For power electronics and radio-frequency applications, both types of polarization are the most desired because the dielectrics will Received: August 7, 2013 Revised: November 29, 2013 Published: December 11, 2013 9698
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normal ferroelectricity will be achieved and cooperative dipole/ domain switching will result in significant hysteresis in the electric displacement−electric field (D−E) loop (Scheme 1D). This is not desired for electric energy storage. Current research efforts focus on ferroelectric polymers having narrow double (Scheme 1B) or single (Scheme 1C) hysteresis loops, which are novel ferroelectric behaviors connecting between the dielectric (Scheme 1A) and ferroelectric (Scheme 1D) behavior, as we reported recently.18 Because of the semicrystalline nature of ferroelectric polymers, both internal crystal structure and crystal− amorphous interaction play important roles in achieving these novel ferroelectric behaviors, as shown in Schemes 1B,C. This is different from 100% crystalline ceramics. By utilizing crystal isomorphism, novel ferroelectric behavior can be obtained for relaxor ferroelectric polymers, such as electron-beam (e-beam)irradiated P(VDF-co-trifluoroethylene) [P(VDF-TrFE)] random copolymer19−21 and P(VDF-TrFE)-based random terpolymers.22−25 We attribute these novel ferroelectric behaviors to the formation of nanodomains in the crystalline structure, pinned by either the cross-links in e-beam-irradiated P(VDFTrFE) crystals (chemical pinning) or the larger third monomers in P(VDF-TrFE)-based terpolymers (physical pinning).18 The chemical pinning will result in narrow single loops, whereas the physical pinning will result in double hysteresis loops (DHLs) due to reversible relaxor ferroelectric ↔ ferroelectric phase transitions induced by the poling electric field. This DHL behavior is different from the antiferroelectric phase/behavior for ceramic26 and liquid crystalline27 materials because antiferroelectric phases are difficult, if not impossible, for crystalline polymers. We consider that this is attributed to the long chain nature of crystalline polymers. For example, even-numbered nylons have alternating hydrogen-bonded sheets in the crystalline structure. However, these alternating hydrogen-bonded sheets cannot be polarized into the same direction because even-numbered alkane spacers adopt a zigzag conformation in the crystal. In addition to the above effect of crystal internal structure (i.e., isomorphism), DHLs can also be achieved by fine-tuning the crystal−amorphous interaction, i.e., the competition between the polarization and depolarization fields in a semicrystalline ferroelectric polymer.28,29 Recently, we introduced a nanoconfinement effect to achieve the DHL behavior for PVDF-based graft copolymers [e.g., polystyrene (PS) grafts].29−31 After PVDF crystallization, an ultrathin dielectric PS layer will wrap around the PVDF crystal to effectively reduce the compensation polarization and thus the polarization field to be lower than the depolarization field inside the crystal. As a result, DHLs are achieved and can persist to high electric fields at room temperature. As we can see in Scheme 1, the DHL behavior is more desirable for electric energy storage because the discharged energy density is the highest if similar maximum polarization can be achieved. In this study, we continued to investigate the crystal orientation and temperature effects on the DHL behavior in the P(VDF-TrFE-CTFE)-g-PS(14%) graft copolymer (CTFE represents chlorotrifluoroethylene). Intriguingly, crystal orientation has a profound effect on the electrical behavior of the graft copolymer. For the nonoriented film, dielectric rather than ferroelectric behavior is observed. For orientated films, DHLs are observed up to 75 °C. Above 100 °C, significant dc conduction and ion migrational losses result in large hysteresis in the D−E loops. By using broadband
have extremely low losses. Current state-of-the-art polymer dielectrics, such as biaxially oriented polypropylene (BOPP), utilize both types of polarization in order to meet the stringent requirement of low losses in practical applications.6 If a polymer contains permanent dipoles, dipolar orientational polarization can take place in a frequency range between 1 Hz and several MHz, depending upon the mobility of permanent dipoles (e.g., amorphous versus crystalline dipoles). If a polymer contains ions, ionic polarization will happen below a couple hundred Hz. Finally, if the polymer system contains multiple immiscible components, Maxwell−Wagner−Sillars (MWS) interfacial polarization will result as long as there are large enough contrasts between dielectric constant and/or electronic conductivity.7 Because ionic and interfacial polarizations take place under low frequencies (below a few hundred Hz), they are not suitable for electric energy storage applications with fast repetition rates (i.e., μs to ms discharge time). Several strategies are presently pursued to achieve high dielectric constant and low loss polymer dielectrics. First, new polymers with enhanced electronic and/or atomic polarizations are proposed. For example, on the basis of the first-principles calculations, Si- and Ge-containing polymers exhibit higher electronic and atomic polarizability than C-based polymers.8,9 Current research challenge lies in a delicate balance between electronic polarizability and bandgap because a high electronic polarizability often results in a low bandgap, which will decrease the insulation property of polymers. Second, polymer-based composites, including polymer/ceramic nanocomposites (or nanodielectrics)2,10 and poly(vinylidene fluoride) (PVDF)based multilayer polymer films,11−14 are actively investigated. The idea is to combine the high dielectric constant property of ceramics and PVDF with the high insulating property of dielectric polymers. For polymer nanodielectrics, the current challenge focuses on the uniform dispersion of ceramic nanoparticles while maintaining (or even enhancing) the breakdown strength.2,10 For PVDF-based multilayer films, research efforts are dedicated to enhancing energy density and breakdown strength by using MWS interfacial polarization, while reducing impurity ion migrational loss and dc conduction at high temperatures.11−17 Recently, we propose to utilize the dipolar orientational polarization for high dielectric constant and low loss dielectrics, if dipolar orientational relaxation can be pushed to high frequencies (i.e., above the GHz range).1 By introduction of permanent dipoles into a nonpolar dielectric polymer (Scheme 1A), the dielectric nonlinearity will gradually increase. If a large number of permanent dipoles form large ferroelectric domains, Scheme 1. Schematics of (A) Dielectric, (B) Double Hysteresis Loop, (C) Relaxor Ferroelectric, and (D) Normal Ferroelectric Behavior for Polymersa
a
From left to right, the dielectric nonlinearity increases. The shaded areas represent the discharged energy density. 9699
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Scheme 2. Synthesis of P(VDF-TrFE-CTFE)-g-PS Graft Copolymer from a Precursor P(VDF-TrFE-CTFE) Random Terpolymer Using ATRP, Followed by Dechlorination Using nBu3SnH to Avoid Thermal Cross-Linking during Later Melt Processing
Figure 1. Bipolar D−E hysteresis loops at room temperature for the P(VDF-TrFE-CTFE)-g-PS(14%) graft copolymer with different extension ratios: (A) 100%, (B) 200%, (C) 280%, and (D) 450%. The poling frequency is 10 Hz with a triangular waveform, and the poling field is increased at increments of 25 MV/m. The corresponding 2D WAXD patterns are shown as insets. The Herman’s orientation factor (f) is indicated for each sample.
dielectric spectroscopy (BDS) study and theoretical fitting, the ion migrational loss is understood after subtracting the dc conduction from the experimental D−E hysteresis loops.
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thermally polymerized PS homopolymer), and dried to constant weight in a vacuum oven for 3 days. From the 19F NMR results,29,31 all CTFE resonance peaks disappeared after the graft copolymerization, indicating that they were successfully grafted with PS side chains. Film Fabrication and Processing. Thick films with a thickness around 80−100 μm were prepared by hot-pressing the graft copolymers at 240 °C, followed by immediate quenching into an ice/water bath. The hot-pressed thick films were uniaxially stretched at 110 °C to different extension ratios at a speed of 12.7 mm/min using a home-built stretching apparatus. After stretching, the films were further annealed under tension at 110 °C for 30 min to release internal stresses and enhance crystallinity. The final thickness of stretched films was around 25 μm. Instrumentation and Characterization. Two-dimensional (2D) wide-angle X-ray diffraction (WAXD) was performed on a Rigaku WAXD instrument (MacroMax 002+) with an 18 kW rotating anode X-ray generator operating at 45 kV and 0.88 mA. The X-ray wavelength (λ) was Cu Kα 0.1542 nm. The d-spacing was calibrated by using silver behenate with the first-order reflection at a scattering vector [(4π sin θ)/λ, where θ is the half scattering angle] of 1.076 nm−1. Additional WAXD experiments were also carried out at the synchrotron beamline X27C, National Synchrotron Light Source
EXPERIMENTAL SECTION
Materials. The P(VDF-TrFE-CTFE)-g-PS(14%) graft copolymer was synthesized by atom-transfer radical polymerization (ATRP) from a P(VDF-TrFE-CTFE) (88.8/7.7/3.5) precursor terpolymer (Scheme 2).29,32 Briefly, the P(VDF-TrFE-CTFE) (88.8/7.7/3.5) was obtained by partial dechlorination of P(VDF-CTFE) 88.8/11.2 (3 M Dyneon 31508 with 11.2 mol % CTFE). Short PS side chains with an average degree of polymerization of 3.0 were grafted to the random terpolymer main chain by ATRP. The PS weight fraction was 14 wt %. The P(VDF-TrFE-CTFE)-g-PS(14%) sample was completely dechlorinated using nBu3SnH to remove all −Cl end groups on the grafted PS side chains to avoid further cross-linking during storage and later film processing processes. Finally, the crude product was dissolved in N,N′dimethylformamide (DMF) and was purified by passing through silica gel columns several times to remove the copper catalysts as much as possible. The purified sample was precipitated in 1:1 (vol/vol) deionized water/methanol mixture, washed by toluene (to remove any 9700
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(NSLS) at Brookhaven National Laboratory (BNL). The wavelength of the incident X-ray was 0.1371 nm. A MAR charge-coupled device (CCD) was used as detector, and the typical data requisition time was 60 s. One-dimensional (1D) XRD curves were obtained by integrating the corresponding 2D patterns radially. Fourier transform infrared (FTIR) was carried out on a Nexus 870 FTIR spectrometer in transmission mode. The scan resolution was 2 cm−1 with total of 32 scans. Differential scanning calorimetry (DSC) was performed on a TA Q100 DSC at a scanning rate of 10 °C/min. Around 5 mg of sample was used. The D−E hysteresis loop measurements were carried out on a Premiere II ferroelectric tester from Radiant Technologies, Inc., in combination with a Matsusada AMT-20B10-B high voltage amplifier (0−10 kV ac). The applied voltage had a bipolar triangular or sinusoidal waveform at 10 Hz. Aluminum (Al) electrodes (ca. 4.5 mm2) with a thickness around 30 nm were evaporated onto both sides of film samples using a Denton benchtop Turbo III carbon/metal evaporator. The film sample with Al electrodes was immersed in a silicone oil (Fisher 460-M3001) bath for D−E loop measurements, and the temperature was controlled by using an IKA RCT temperature controller (Wilmington, NC). Low temperatures were achieved by cooling the silicone oil bath with a dry ice/isopropanol bath. A homebuilt fixture was used for connecting the electrodes on the film to the Radiant ferroelectric tester using high voltage cables. Stored and discharged energy densities (Ue) were obtained by integrating the corresponding charging and discharging curves, Ue = ∫ E dD. The unreleased percentage of the stored energy density was defined as unreleased% = 100(1 − Ue,released/Ue,stored). Detailed experimental setup and measurement principles can be found in our previous report.29 Frequency-sweep broadband dielectric spectroscopy (BDS) measurements were performed on a Novocontrol spectrometer (Hundsangen, Germany) under a vacuum of 40 mTorr. The applied voltage was 1 Vrms with frequency changing from 0.001 Hz to 100 kHz, and the test temperature varied from 25 to 125 °C using a home-built temperature control unit. Temperature-sweep BDS measurements were performed on a Novocontrol Concept 80 dielectric spectrometer with frequency ranging from 1 Hz to 1 MHz at 1 Vrms. Al or Ag electrodes were sputtered onto both surfaces of the film with a thickness around 30 nm, and the electrode area was 0.786 cm2 for BDS measurements.
strongest reflection peak in the diffraction pattern, and the results were f = 0, 0.54, 0.62, and 0.72, respectively. For uniaxially stretched films, the crystals were orientated with their c-axes parallel to the drawing direction because the (020)α and (110)α/(110/200)β reflections were concentrated on the equator. The film crystallinity was determined by peak-fitting of the crystalline reflections and the amorphous halo in the 1D WAXD profiles integrated from Figure 1. The results were 32.3%, 33.8%, 34.0%, and 34.6%, respectively, for films with different orientation factors, f = 0, 0.54, 0.64, and 0.72. Apparently, the crystallinity did not change much as the degree of crystal orientation increased. Therefore, it should have little influence on the DHL behavior. For the nonstretched film, the D−E loops appeared fairly linear (i.e., dielectric-like) up to 130 MV/m in Figure 1A (note that electric breakdown happened above 130 MV/m). As the Herman’s orientation factor f increased to 0.54, slight DHLs were observed above 100 MV/m (Figure 1B). With further increasing the degree of crystal orientation, more obvious DHLs were observed (see Figures 1C,D). This could be explained by the long-range interactions among ferroelectric domains composed of aligned P(VDF-TrFE) dipoles. In a recent report,27 we have discussed the long-range interactions among ferroelectric domains in semicrystalline PVDF (see Figures 13A,B in ref 27). These interactions are realized through the compensation polarization outside the ferroelectric domain and are anisotropic in nature, when an external electric field is applied. Namely, when two neighboring ferroelectric domains are aligned parallel to the external field, the interaction is the strongest. When they are aligned perpendicular to the external field, the interaction is the weakest. Currently, these interactions cannot be directly evidenced using available characterization tools, such as piezoresponse force microscopy (PFM), primarily owing to the small domain sizes (possibly nanometer to tens of nanometers) in the sample. Nonetheless, the anisotropic domain−domain interactions have already been reported in BaTiO3 single crystals,33 and the concept of the long-range interactions among ferroelectric domains has been widely accepted in the ferroelectric community.34 In the nonstretched film, chains and dipoles in the P(VDFTrFE) crystal were randomly oriented, and only a small fraction of the dipoles could respond to the applied field. Meanwhile, these crystalline dipoles must be randomly distributed in the sample, and the long-range interactions among them were weak. As a result, little dipole flipping was observed up to 130 MV/m, although we speculated that dipoles might flip at high enough fields. For stretched films with crystals c-axes parallel to the drawing direction, more crystalline dipoles were capable of orienting along the field direction. These polarizable dipoles/ domains must have increased long-range interactions, and the strong interactions (or high compensation polarization) resulted in ferroelectric hysteresis at high enough electric fields (>100 MV/m). Upon increasing the extension ratio, more dipoles/domains in the P(VDF-TrFE) crystals became polarizable, and more significant broadening in the hysteresis loops was observed (Figures 1B−D). Nonetheless, due to the nanoconfinement effect from the PS side chains outside the P(VDF-TrFE) crystal, all hysteresis loops appeared as DHLs, as we reported before.29 To quantify the changes in hysteresis loops for films with different degrees of crystal orientation, the maximum and remanent electric displacements (Dmax and Drem) were plotted as a function of the electric field (Figure 2). In Figure 2A, the
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RESULTS AND DISCUSSION Effect of Crystal Orientation on DHL Behavior. Recently, the DHL behavior was reported for a uniaxially stretched P(VDF-TrFE-CTFE)-g-PS(14%) film (extension ratio = 600%), and it was attributed to the nanoconfinement effect by the PS side chains wrapping around P(VDF-TrFE) crystals.29 In this study, the effect of crystal orientation was investigated for P(VDF-TrFE-CTFE)-g-PS(14%) films with different extension ratios. The corresponding bipolar hysteresis loops are shown in Figure 1. From Figures 1A−D, the extension ratios were 100%, 200%, 280%, and 450%, respectively. For the nonstretched sample (extension ration = 100%), the α-form was observed for P(VDF-TrFE) crystals as evidenced by the (020)α and (110)α reflection rings (see the inset WAXD pattern in Figure 1A). For stretched samples, a small amount of β-form crystals gradually developed, and its content slightly increased upon increasing the extension ratio. From our previous report, the α and β contents were ca. 31.5 and 6.5 wt %, respectively, for the film with an extension ratio of 600%, and the β crystals did not significantly affect the DHL behavior.29 The degree of crystal orientation for the combined (110)α and (110/200)β reflections at 12.95−14.25 nm−1 (see arrows in the insets of Figure 1) was determined from the 2D WAXD patterns in the insets. The orientation factor, Herman’s factors ( f), for the four patterns was calculated based on the 9701
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Figure 2. (A) Maximum electric displacement (Dmax) and (B) remanent electric displacement (Drem) as a function of the electric field for P(VDFTrFE-CTFE)-g-PS(14%) films with different Herman’s orientation factors.
Figure 3. (A) Stored and (B) released energy densities and (C) unreleased% as a function of the poling electric field for the P(VDF-TrFE-CTFE)-gPS(14%) films with different degrees of crystal orientation. (D) The stored (solid symbols) and released (open symbols) energy densities as a function of orientation factor at different poling fields.
Dmax increased linearly with electric field for the nonstretched film, indicating a typical dielectric behavior with an apparent dielectric constant of 9.7 (consistent with the value of 9.2 at 100 Hz in our previous report29). Upon increasing the degree of crystal orientation, more deviation from the linear behavior was observed, suggesting enhanced ferroelectricity above 75 MV/m. In Figure 2B, the Drem increased almost linearly for the nonstretched film. For stretched films, the nonlinear deviation in the Drem occurred only above 100 MV/m. This was attributed to the onset of dipole switching above 100 MV/m, as seen in Figures 1B−D. Meanwhile, there was no significant difference in the Drem values for different degrees of crystal orientation, suggesting that dipoles in these films were reversible.35
The stored and released energy densities (Ues) and the unreleased% energy density were calculated from the D−E loops in Figure 1. The corresponding results are shown in Figure 3. From Figures 3A,B, the stored and released Ues increased upon increasing the electric field and the orientation factor. It is noted that the released Ue was much higher than that of BOPP. For example, at 125 MV/m, the released Ues for the nonstretched film and the stretched film with f = 0.72 were 3.25 and 5.7 times that of BOPP. The unreleased% is shown in Figure 3C with two regimes observed. First, below 75 MV/m, the unreleased% decreased with increasing the orientation factor up to 0.62, and finally the unreleased% for the film with f = 0.72 showed similar values as those for the film with f = 0.62. This could be attributed to the crystal orientation effect on 9702
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Figure 4. (A) Repeated bipolar poling at 150 MV/m for the stretched P(VDF-TrFE-CTFE)-g-PS(14%) film with an extension ratio of 450% at room temperature. The poling frequency is 10 Hz with a triangular waveform. (B) 1D WAXD profiles and (C) FTIR spectra for the stretched P(VDF-TrFE-CTFE)-g-PS(14%) film before and after 300-cycle electric poling to 150 MV/m. Corresponding 2D WAXD patterns are shown as insets in (B). The weak reflection peak from aluminum electrode coating is seen at 26.6 nm−1 in (B).
Figure 5. (A−G) Bipolar D−E hysteresis loops (centered) for a stretched (extension ratio = 500%) P(VDF-TrFE-CTFE)-g-PS(14%) film at different temperatures: (A) −40, (B) 0, (C) 25, (D) 50, (E) 75, (F) 100, and (G) 125 °C. (H) shows raw data loops at 125 °C. The poling frequency is 10 Hz with a triangular waveform, and the poling field has increments of 25 MV/m until electric breakdown at each temperature.
dielectric loss, as we reported before.29 For the nonstretched film, the majority of α crystals randomly oriented, and part of them would orient with their chain axes more or less parallel to the field. It is the dipole switching along the α-form chains that causes significant dielectric loss even at low fields (see the αc relaxation peak around 10 Hz in Figure 8C).28 With increasing the extension ratio, more crystals could orient with their chain axes along the drawing direction (or perpendicular to the field direction). As a result, the dielectric loss would decrease upon increasing the orientation factor (see disappearance of the αc relaxation peak around 10 Hz in Figure 8F). Finally, after the chains fully oriented, the decrease in dielectric loss would cease. Second, above 100 MV/m, the oriented films had much higher unreleased% values (e.g., 26−29% at 150 MV/m) than the nonstretched film (14% as extrapolated at 150 MV/m). This was attributed to the ferroelectric loss from dipole switching induced by high poling fields. Figure 3D summarizes the effect of crystal orientation on the stored and released Ues. Below 75 MV/m, the stored and released Ues did not change much for samples with different orientation factors because dipoles did not switch at all. At high electric fields (>75 MV/m), the difference among different samples became more obvious, and
the stored and released Ues increased upon increasing the orientation factor. From these results, we conclude that the parallel crystal orientation is beneficial for high energy density and low loss applications. Nonetheless, the hysteresis at high electric field shall be minimized for better performance. This may be achieved by utilizing more reversible dipoles in P(VDFTrFE) with a higher TrFE content than 25 mol %. The DHL behavior was also observed for P(VDF-TrFE) with high TrFE contents.36−41 However, it would disappear after several poling cycles or by annealing at high temperatures. The DHLs for stretched P(VDF-TrFE-CTFE)-g-PS(14%) films were fairly stable. Figure 4A shows that DHLs in the P(VDFTrFE-CTFE)-g-PS(14%) film persisted even after poling to 150 MV/m for 300 cycles at room temperature. For the first 10 cycles, the Dmax had the most significant changes. Afterward, it only increased slightly. Meanwhile, the Drem also increased slightly as the Dmax increased. To understand the reason for slight Dmax and Drem increases, WAXD and FTIR experiments were carried out for the stretched P(VDF-TrFE-CTFE)-gPS(14%) film before and after 300-cycle electric poling at 150 MV/m. As shown in Figure 4B, the integrated 1D WAXD profiles before and after 300-cycle electric poling largely 9703
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Figure 6. (A) Real (εr′) and (B) imaginary (εr″) permittivity as a function of temperature for hot-pressed nonoriented P(VDF-TrFE-CTFE)-gPS(14%) film (70.7 μm). (C) First cooling and second heating DSC curves for P(VDF-TrFE-CTFE)-g-PS(14%). The scanning rate is 10 °C/min.
in the main chains. To understand the D−E loop changes around 75 °C, we carried out temperature-sweep BDS at various frequencies and DSC studies. In both experiments, nonoriented samples were used. From the BDS and DSC results in Figure 6, several thermal events were observed. First, the glass transition temperature (Tg) was seen at −40 °C at 1 Hz in BDS and −35 °C in DSC (see inset of Figure 6C). Note that BDS is sensitive to dipole mobility whereas DSC is sensitive to heat capacity. The lower Tg by BDS indicated that dipoles started to become mobile before any heat capacity change could be detected by DSC (note that the heat capacity change at Tg in DSC was fairly weak). Around 12 °C, a shoulder peak was observed in the imaginary permittivity (εr″) plot in Figure 6B, and a weak peak was observed in the heating curve in DSC (inset of Figure 6C). These weak peak could be attributed to the dipole relaxation in the α-form crystals (i.e., the αc relaxation).12,42 Above 50 °C, the onset of crystal melting was observed in the DSC heating curve in Figure 6C. Correspondingly, both real permittivity (εr′) and εr″ started to increase, as seen in Figures 6A,B. We called the region between 50 and 100 °C as the premelting stage, and it shall be responsible for the D−E loop change around 75 °C. Around 120 °C, a premelting peak was observed in the DSC heating curve, causing significant increases in both εr′ and εr″, which could be attributed to impurity ion and dc conduction losses. Detailed analysis will be discussed below. Finally, the major crystal melting peak was observed around 165 °C in DSC. Note that no significant premelting peak was observed for the parent P(VDF-TrFE-CTFE) (88.8/7.7/3.5) terpolymer (data not shown). The appearance of the premelting peak around 120 °C could be attributed to the grafting of short PS side chains on the P(VDF-TrFE-CTFE) main chains, which disrupted the crystallization capability and crystallite sizes. The low-field elliptical/rounded loops at 100 and 125 °C could not be simply attributed to the dipole switching because of the lack of dipole saturation (i.e., the “S” shape in the loop). First, on the basis of our recent report, migrational loss from even a trace amount of impurity ions (e.g., 0.1 μM level) could significantly broaden the D−E hysteresis loop,11,12 especially at elevated temperatures when the ion mobility were high. Although we have stringently purified the P(VDF-TrFECTFE)-g-PS(14%) sample by passing through silica gel columns several times after ATRP, a trace amount of copper ions was still possible and the increased ion mobility at high temperatures could broaden the hysteresis loops. Second, the D−E loops shown in Figures 5A−G were presented after
overlapped, indicating no noticeable changes in crystallite size after 300-cycle electric poling. However, after 300-cycle electric poling, the FTIR spectrum for the stretched P(VDF-TrFECTFE)-g-PS(14%) film showed intensity changes for absorption peaks at 841 (increase), 532 (decrease), and 472 cm−1 (increase) (see Figure 4C). The absorption peaks with increased intensity at 841 and 472 cm−1 were attributed to the β-form crystal, whereas the absorption peak with decreased intensity at 532 cm−1 was assigned to the α-form crystal. This result clearly indicated that there was a minor α → β phase transition in P(VDF-TrFE) crystals after extensive electric poling. It was noted that WAXD was less sensitive than FTIR in identifying the α → β phase transformation for PVDF copolymers because X-ray was sensitive to interchain spacing whereas FTIR was sensitive to intrachain conformation. On the basis of these results, we conclude that the gradual increase of both Dmax and Drem upon electric poling should be largely attributed to slight electric poling-induced α → β phase transition in P(VDF-TrFE) crystals. Effect of Temperature on the DHL Behavior. The thermal stability of the DHL behavior in the stretched P(VDFTrFE-CTFE)-g-PS(14%) film was studied by bipolar D−E loop tests, and the corresponding results are shown in Figure 5. The overall DHL shapes did not change much from −40 to 75 °C. Under low electric fields (
