Relaxor Ferroelectric Behavior from Strong Physical Pinning in a Poly

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Relaxor Ferroelectric Behavior from Strong Physical Pinning in a Poly(vinylidene fluoride-co-trifluoroethylene-cochlorotrifluoroethylene) Random Terpolymer Lianyun Yang,† Brady A. Tyburski,† Fabrice Domingues Dos Santos,‡ Maya K. Endoh,§ Tadanori Koga,§ Daniel Huang,† Yijun Wang,† and Lei Zhu*,† †

Department of Macromolecular Science and Engineering, Case Western Reserve University, Cleveland, Ohio 44106-7202, United States ‡ Piezotech S.A.S, Arkema-CRRA, Rue Henri Moissan, 69493 Pierre-Benite, Cedex, France § Department of Materials Science and Engineering, Stony Brook University, Stony Brook, New York 11794-2275, United States S Supporting Information *

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γ-ray) radiation is needed. For example, a 40−60 Mrad dose of high energy (e.g., 1.2 MeV) e-beam at 70 °C (ca. 1 Mrad/min) is required to achieve narrow SHL behavior for P(VDFTrFE).16 Such a high dose at high temperatures is practically difficult to implement even industrially. Second, high energy ebeam irradiation significantly damages P(VDF-TrFE) films via cross-linking and chain scission reactions from ionizing radiation. As a result, P(VDF-TrFE) films become brittle and easy to break after a high dose of e-beam irradiation. Third, a large amount of space charges, including electrons and ions, is generated in the sample, and thus electric breakdown strength decreases for e-beam-irradiated P(VDF-TrFE) films. Therefore, it is highly desirable to achieve narrow SHL behavior for P(VDF-TrFE)-based terpolymers, which are free of the abovementioned problems for ionizing radiation. Nonetheless, from our previous report,16 only the DHL behavior is observed for P(VDF-TrFE-CFE) terpolymers due to weak physical pinning by the CFE units in isomorphic crystals. Although a higher discharged energy density is observed for P(VDF-TrFE-CFE) (see Figure S3D in the Supporting Information), its hysteresis loss at high electric fields is significantly higher than that of the e-beamed P(VDF-TrFE) (see Figure S3E in the Supporting Information). Here, we ask a question: Is it possible to permanently pin the P(VDF-TrFE) crystals using a strong physical pinning in P(VDF-TrFE)-based terpolymers? To answer this question, a more suitable comonomer is needed to replace CFE in the terpolymer. As we know, CTFE has a larger size and a lower dipole moment (ca. 0.64 D17) than CFE (dipole moment of ca. 1.8 D18), indicating its potential as a better third comonomer for P(VDF-TrFE)-based terpolymers. In this work, a P(VDF-TrFE-CTFE) 62.2/30.2/7.6 (molar ratios) terpolymer is chosen to study whether the physical pining force of CTFE is strong enough to permanently pin the P(VDF-TrFE) crystals and achieve the narrow SHL behavior.

INTRODUCTION Novel ferroelectric behaviors in relaxor ferroelectric (RFE) polymers, such as narrow single hysteresis loop (SHL) and double hysteresis loop (DHL) behaviors, attract increasing research attention because of their application potentials in electric energy storage,1−4 electrostrictive actuation,5−7 and electrocaloric cooling.8,9 These novel ferroelectric behaviors have been achieved in electron beam (e-beam) irradiated P(VDF-TrFE) random copolymers,5,10,11 P(VDF-TrFE)-based random terpolymers,11−13 and P(VDF-TrFE-CTFE)-g-PS graft copolymers.14−16 Here, VDF is vinylidene fluoride, TrFE is trifluoroethylene, CTFE is chlorotrifluoroethylene, and S is styrene. It is essential to understand the origin of these novel ferroelectric behaviors because it will help to guide the design and development of new ferroelectric polymers with better electroactive performance. On the basis of our previous study,11 novel ferroelectric behaviors in a P(VDF-TrFE-CFE) 59.2/ 33.6/7.2 (molar ratios) terpolymer (CFE is 1,1-chlorofluoroethylene) and an e-beam-irradiated P(VDF-TrFE) 50/50 (molar ratio) copolymer are explained by physical and chemical pinning effects on isomorphic (i.e., defect-modified) crystals. Briefly, the large CFE comonomers in the terpolymer (physical pinning) and radiation-induced −CF3 side groups in the ebeam-irradiated copolymer (chemical pinning) effectively increase the interchain distance and thus reduce the ferroelectric (FE) domain size to nanoscales (i.e., nanodomains). Both crystal pinning by large structural defects and the existence of nanodomains can lead to the formation of the so-called RFE structure or phase. For P(VDF-TrFE-CFE), the physical pinning force is relatively weak and eventually disappears under a high poling field (>75 MV/m). As a result, reversible, electric-field-induced RFE ↔ FE conformation transitions produce the DHL behavior. On the contrary, the chemical pinning force in e-beamed P(VDF-TrFE) is strong, and thus no RFE ↔ FE conformation transitions are seen. As a consequence, the narrow SHL behavior is achieved. To keep low dielectric loss for real applications, the SHL behavior is more advantageous than the DHL behavior. However, e-beam-irradiated P(VDF-TrFE) is disadvantageous in several aspects. First, due to high chemical stability of fluoropolymers, a significant amount of high energy e-beam (or © XXXX American Chemical Society

Received: September 6, 2014 Revised: November 3, 2014

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RESULTS AND DISCUSSION Samples for this study were obtained by uniaxially stretching hot-pressed P(VDF-TrFE-CTFE) films (ca. 100 μm) with a drawing ratio of ca. 600% at room temperature. Immediately after mechanical stretching, the sample was characterized by differential scanning calorimetry (DSC) and two-dimensional (2D) wide-angle X-ray diffraction (WAXD). The corresponding results are shown in Figures S4A and S4B in the Supporting Information, respectively. From the DSC results in Figure S4A, two weak Curie transition temperatures (TCs) were observed at 28.4 and 47.1 °C, respectively, upon the first heating process. However, no distinct (110/200)FE reflections for a separate FE phase were seen, except the (110/200)RFE reflections for the RFE structure, on the equator of the 2D WAXD pattern in Figure S4B (see detailed explanation for the 2D WAXD pattern later). This was different from the uniaxially stretched P(VDFTrFE-CFE) terpolymer films, which showed a separate FE phase in addition to the RFE structure after uniaxial stretching, as reported previously.11 Since the 2D WAXD pattern did not show a separate FE phase in the uniaxially stretched P(VDFTrFE-CTFE) film, the higher TC at 47.1 °C should be attributed to larger FE domains induced by mechanical stretching. Note that it was reported that the TC of P(VDFTrFE) copolymers could be largely affected by the FE domain size.19−21 To avoid the interference of mechanically induced larger FE domains, the stretched films were further annealed under tension at 80 °C (i.e., above the TC) for 2 h to eliminate larger FE domains and release internal stresses. After annealing, DSC results showed a single RFE → FE phase transition at 28 °C, together with a melting peak at 126 °C, during the first heating process (see Figure S5 in the Supporting Information). The crystalline structure in the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) film was characterized by WAXD and FTIR (see Figure 1A,B). One-dimensional (1D) WAXD profiles for the equator and meridian reflections were obtained by integrating the corresponding two-dimensional (2D) WAXD patterns. At −40 °C, a sharp reflection was seen on the equator, together with two weak ones at higher q values (see Figure 1A-a and the 2D WAXD pattern in Figure S6A in the Supporting Information). The relationship among these equator reflection peaks was 1:√3:√4, suggesting a pseudohexagonal packing in the crystal. Meanwhile, two broad reflections were centered on the meridian (see Figure 1A,b and the 2D WAXD pattern in Figure S6A in the Supporting Information). Given the single RFE → paraelectric (PE) transition in the DSC (see Figure S5 in the Supporting Information), this WAXD pattern could be explained by a single defective RFE crystal structure, which showed some similarity to that of the cooled monoclinic FE phase of P(VDFTrFE).19,22 Peak assignments are given in Figure 1A and the 2D WAXD pattern in Figure S6A in the Supporting Information. In this crystal structure, a small amount of long Tn (n > 4) sequences [see the weaker (001) reflection on the meridian] coexisted with major TmG and/or TmS (m ≤ 4) sequences [see the stronger (111/201) reflection on the meridian]. Here, T represents the trans conformation, G is the gauge conformation, and S is the skew linkage.19,22 Note that n = 4 is tentatively taken as the separation point between long and short T sequences (for a detailed explanation please refer to the discussion on the FTIR results later). This RFE structure in P(VDF-TrFE-CTFE) was different from that of P(VDF-TrFECFE) (see ref 11 and the 2D WAXD pattern in Figure S6B in

Figure 1. (A) 1D WAXD profiles for (a) equator and (b) meridian reflections of the uniaxially stretched and 80 °C annealed P(VDFTrFE-CTFE) terpolymer film with a draw ratio = 600%. Between −5 and 20 °C, a weak amorphous halo is centered at 17 nm−1, which is attributed to some condensed water on the window. (B) FTIR spectra for the P(VDF-TrFE-CTFE) terpolymer film. The right panel is from the uniaxially stretched and 80 °C annealed film (ca. 15 μm thick), and the left panel is from the solution-cast film (ca. 2 μm thick). Both WAXD and FTIR data are collected during a stepwise heating process using increments of 5 °C. At each temperature, the samples are held isothermally for 10 min before data collection. All curves are overlaid for clarity.

the Supporting Information). P(VDF-TrFE-CFE) had a similar 2D WAXD pattern as that of the high-temperature PE phase (rather than the cooled FE phase) of P(VDF-TrFE)19,22 because there was only one broad reflection on the meridian. In the crystal structure of P(VDF-TrFE-CFE), the RFE phase, therefore, should consist of more or less random TmG (m < 4) sequences. Upon increasing the temperature, the d-spacing for the (110/ 200) reflection gradually increased to 5.18 Å at 90 °C. Because of the larger size of the CTFE comonomer, the d-spacing for the (110/200) reflection in P(VDF-TrFE-CTFE) was always slightly larger than that for the (110/200) reflection of P(VDFTrFE-CFE) (e.g., 4.84 Å at −38 °C and 4.98 Å at 90 °C).11 Meanwhile, the (001) reflection, representing the long Tn sequences, gradually decreased its intensity and eventually disappeared above the TC at 28 °C. The chain conformation changes in the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) terpolymer film were further studied by Fourier transform infrared (FTIR) spectrosB

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Figure 2. (A) Real (εr′) and (B) imaginary (εr″) parts of relative permittivity as a function of temperature at different frequencies for the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) terpolymer film. (C) εr′ and (D) εr″ as a function of frequency at different temperatures for the same P(VDF-TrFE-CTFE) terpolymer film.

scan BDS in Figure 2C,D. Below 1 Hz, both εr′ and εr″ significantly increased for temperatures above 50 °C, and the slope of −1 in the double-logarithm plot of εr″ indicated that the low-frequency losses were originated from impurity ions.11,25 Note that the impurity ions were inherited from the suspension terpolymerization.3 Dipole relaxation took place at high frequencies, evidenced by a decrease in the εr′ (Figure 2C) and a relaxation peak in the εr″ (Figure 2D). Plateaus in the εr′ were observed at intermediate frequencies, and the maximum εr′ (∼53 at 1 Hz) was observed between 25 and 50 °C. This was consistent with the temperature-scan result in Figure 2A. From this BDS study, the RFE → PE Curie transition appeared fairly broad (Figure 2A), which is the signature of the RFE behavior,5,11 although the RFE structure contained a certain amount of long Tn sequences in the P(VDF-TrFE-CTFE) crystals. The ferroelectric behavior of the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) terpolymer film was studied by electric displacement−electric field (D−E) loop measurements. Temperature-dependent D−E loops under different poling frequencies are shown in Figure 3. At 0 °C, relatively broad hysteresis loops with a DHL shape were observed under all frequencies. The large FE hysteresis could be attributed to the presence of long Tn sequences in the RFE structure of P(VDF-TrFE-CTFE). These long Tn sequences would form FE domains, which might be similar to the kink bands in the cooled FE phase in P(VDF-TrFE).19,22 These FE domains then resulted in large FE hysteresis. In addition, high enough poling fields (e.g., above 75 MV/m) could further increase the length of long Tn sequences. Note that these increased long Tn

copy, and results are shown in Figure 1B. The spectra in the right panel were obtained from the uniaxially stretched and 80 °C annealed film, whose thickness was ca. 15 μm. However, this was too thick for the absorption bands above 950 cm−1, and significant peak saturation was observed. To avoid this, a thinner solution-cast film (ca. 2 μm thick) was used to obtain the spectra above 950 cm−1 (i.e., the left panel of the figure). Below the TC of 28 °C, the long Tn sequences were evidenced by absorption bands at 507 (T3), 848 (T>3), and 1287 cm−1 (T>4).23,24 Note that bands at 848 and 1287 cm−1 were either weak or absent for the uniaxially stretched and 70 °C annealed P(VDF-TrFE-CFE) terpolymer film below its TC,11 indicating that the T sequences in P(VDF-TrFE-CFE) might not exceed 4. Upon increasing the temperature to 30 °C, the intensity for bands at 507 and 848 cm−1 decreased and the band at 1287 cm−1 disappeared. Above 30 °C, a new band associated with the short T sequence (T 4) sequences and a major population of random TmG or TmS (m ≤ 4) sequences. This structure has a certain similarity to the cooled FE phase of P(VDF-TrFE).19,22 On the contrary, the hexagonal-like RFE structure of P(VDF-TrFE-CFE) looks more or less like the high-temperature PE phase of P(VDF-TrFE)19,22 and contains mostly TmG (m ≤ 4) conformation. As a result, P(VDF-TrFECTFE) exhibits a higher εr (70) than that (55) of P(VDFTrFE-CFE) in the linear region of its D−E loops (−25 to 25 MV/m) at the cost of larger remanent D or a broader loop shape (see the right panel of Scheme 1). Given the fairly similar comonomer compositions for both samples, the exact reason for this difference is still unclear from this study. We speculate that it may relate to different comonomer sequences during terpolymerization due to different reactivity ratios for CTFE and CFE (currently, the data is not public). For example, if CTFE copolymerizes into short blocks, the VDF/TrFE sequence between neighing CTFE units will be long. As a consequence, long VDF/TrFE sequence will favor the long Tn sequence. On the contrary, CFE may randomly copolymerize into the terpolymer, leading to short VDF/TrFE and thus short Tn sequence. This hypothesis should be detectable by highresolution 19F nuclear magnetic resonance (NMR) spectroscopy. However, this is currently not possible for our samples because they are polymerized to high conversions to avoid wasting expensive comonomers. At such high conversions, composition drift in the terpolymers cannot be avoided, and this will complicate the 19F NMR analysis of comonomer sequences. In the future, we intend to control the conversion below 10% to avoid composition drift and then use 19F NMR to determine the comonomer composition and sequence in order to understand why the RFE structure in P(VDF-TrFECTFE) is more ferroelectric than that in P(VDF-TrFE-CFE). Alternatively, P(VDF-TrFE-CTFE) can also be obtained by partially reducing P(VDF-CTFE) using tri-n-butyltin hydride, 27,28 and a more uniform comonomer sequence distribution could be achieved. From both approaches, it is possible that an optimum CTFE composition and sequence can be found to achieve the SHL behavior with even narrower D−E loops for P(VDF-TrFE-CTFE), like those observed from ebeamed P(VDF-TrFE) (see Figure S3C in the Supporting Information).5,11 Note that the inhomogeneities in composition and sequence in the terpolymers will not affect our conclusion because CTFE will have a strong physical pinning effect no matter what composition and sequence in the P(VDF-TrFECTFE) terpolymer.

stronger physical pinning by CTFE than by CFE in the terpolymers. To further study the strong physical pinning by CTFE, D−E loop measurements were conducted beyond 150 MV/m. Figure 4 shows continuous D−E loops at maximum poling fields ranging from 200 to 400 MV/m under 10 Hz at 22 °C. At poling fields up to 400 MV/m, only narrow SHLs were observed. This clearly demonstrated that the physical pinning force by CTFE persisted up to a poling electric field of 400 MV/m. The calculated εr values in the linear part of the SHLs remained nearly constant at 70 for all poling fields above 150 MV/m, suggesting the saturation of polarizable dipoles and nanodomains in the sample. Upon increasing the maximum poling field, the SHLs gradually opened up. A careful inspection of the experimental D−E loop data in Figure 4E showed slight up-shifts along the D-axis. This could be attributed to dc conduction as we reported in a recent work.16 Using an equivalent circuit by assuming a constant resistance of the sample,16 the dc conduction loops were obtained as the horizontal elliptical loops above D = 0 shown in Figure 4E. After subtracting these dc conduction loops from the experimental SHLs, narrow SHLs, which were symmetric around D = 0, are clearly seen in Figure 4F. These SHLs should demonstrate the intrinsic property (i.e., free of dc conduction) of strongly pinned RFE crystals in P(VDF-TrFE-CTFE). From the above result, the RFE phases can be achieved for both P(VDF-TrFE-CFE) and P(VDF-TrFE-CTFE) terpolymers by incorporating ca. 7−8 mol % large third comonomers (see Scheme 1). This is attributed to the physical pinning effect and nanodomain structure in the P(VDF-TrFE) crystals, as we have explained in a previous report.11 Briefly, the interchain distance in PVDF is pre-expanded by comonomer TrFE,26 namely, l2 > l1, where l1 and l2 are the interchain distances in PVDF and P(VDF-TrFE) crystals, respectively. Subsequently, the larger third comonomers such as CFE and CTFE can be incorporated into the isomorphic crystalline structure. For a P(VDF-TrFE-X) terpolymer (X = CFE or CTFE), the interchain distance of P(VDF-TrFE) crystals is further expanded by the comonomer X, namely, l3 or l4 > l2, where l3 and l4 are the interchain distances in the P(VDF-TrFE-X) terpolymer crystals (l3 for CFE and l4 for CTFE). Because of the expansion of interchain distance and pinning from the larger third comonomer units (note that the schematic in Scheme 1 is somewhat idealized. In reality, the large third comonomer units may not align together; however, they should still pin the crystals), the P(VDF-TrFE) dipoles in between the pinning points can freely rotate without touching neighboring chains. At an optimum amount of pinning points along the chains (i.e., an optimal third comonomer content), nanodomains may form from the long Tn sequences. As a result, the RFE behavior with narrow D−E loops, rather than normal FE behavior with rectangular shaped D−E loops, will be observed. However, the RFE behavior of P(VDF-TrFE-CTFE) is different from that of P(VDF-TrFE-CFE). The first and the most significant difference is the physical pinning force from CFE and CTFE (Scheme 1). As we mentioned above, CTFE has a larger size (as evidenced by l4 > l3 from WAXD results) and a lower dipole moment (0.64 D) than CFE (dipole moment of 1.8 D). As a result, CTFE has a stronger physical pinning force than CFE. Upon high enough electric field poling, CFE will rotate to increase the long Tn sequences while CTFE cannot. As shown in the right panel of Scheme 1, reversible increase and decrease in the long Tn sequences upon

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CONCLUSIONS In summary, we have compared the RFE structures in P(VDFTrFE-CTFE) and P(VDF-TrFE-CFE) terpolymers with similar CTFE or CFE composition (i.e., 7−8 mol %). WAXD and FTIR results showed that the RFE structure in P(VDF-TrFECTFE) contained more long Tn (n > 4) sequences than P(VDF-TrFE-CFE), and thus P(VDF-TrFE-CTFE) appeared F

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(13) Xia, F.; Cheng, Z.; Xu, H.; Li, H.; Zhang, Q.; Kavarnos, G. J.; Ting, R. Y.; Abdul-Sedat, G.; Belfield, K. D. Adv. Mater. 2002, 14, 1574−1577. (14) Guan, F.; Wang, J.; Yang, L.; Tseng, J. K.; Han, K.; Wang, Q.; Zhu, L. Macromolecules 2011, 44, 2190−2199. (15) Guan, F.; Wang, J.; Yang, L.; Guan, B.; Han, K.; Wang, Q.; Zhu, L. Adv. Funct. Mater. 2011, 21, 3176−3188. (16) Yang, L.; Allahyarov, E.; Guan, F.; Zhu, L. Macromolecules 2013, 46, 9698−9711. (17) Hilczer, B.; Markiewicz, E.; Pogorzelec-Glaser, K.; Polomska, M.; Pietraszko, A. Ferroelectrics 2011, 417, 124−135. (18) Dipole moment (μ) is calculated from the MOPAC program in Chem3D Ultra, μCFE = 1.73 D. This is verified by the dipole moment of CH2FCl of 1.82 D. See: Nelson, R. D.; Lide, D. R.; Maryott, A. A. Selected Values of Electric Dipole Moments for Molecules in the Gas Phase; National Bureau of Standards: Washington, DC, 1967. (19) Tashiro, K. Crystal structure and phase transition of PVDF and related copolymers. In Ferroelectric Polymers: Chemistry, Physics, and Applications, 1st ed.; Nalwa, H. S., Ed.; Marcel Dekker: New York, 1995; pp 63−182. (20) Li, G. R.; Kagami, N.; Ohigashi, H. J. Appl. Phys. 1992, 72, 1056−1061. (21) Barique, M. A.; Ohigashi, H. Polymer 2001, 42, 4981−4987. (22) Tashiro, K.; Takano, K.; Kobayashi, M.; Chatani, Y.; Tadokoro, H. Polymer 1984, 25, 195−208. (23) Kobayashi, M.; Tashiro, K.; Tadokoro, H. Macromolecules 1975, 8, 158−171. (24) Tashiro, K.; Kobayashi, M.; Tadokoro, H. Macromolecules 1981, 14, 1757−1764. (25) Kremer, F.; Schönhals, A. Broadband Dielectric Spectroscopy; Springer: New York, 2003. (26) Lovinger, A. J. Radiation effects on the structure and properties of poly(vinylidene fluoride) and its ferroelectric copolymers. In Radiation Effects on Polymers; Clough, R. L., Shalaby, W., Eds.; American Chemical Society: Washington, DC, 1991; pp 84−100. (27) Lu, Y.; Claude, J.; Neese, B.; Zhang, Q.; Wang, Q. J. Am. Chem. Soc. 2006, 128, 8120−8121. (28) Lu, Y. Y.; Claude, J.; Zhang, Q. M.; Wang, Q. Macromolecules 2006, 39, 6962−6968.

more ferroelectric than P(VDF-TrFE-CFE). Meanwhile, narrow SHL behavior was achieved for the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) terpolymer film above room temperature. This was different from the DHL behavior reported for the uniaxially stretched and 70 °C annealed P(VDF-TrFE-CFE) terpolymer film. This difference was explained by the stronger pinning force by CTFE because of its larger size and smaller dipole moment (0.64 D) relative to CFE (dipole moment of 1.8 D). On the basis of this study, it is possible to further fine-tune the comonomer composition and sequence in P(VDF-TrFE-CTFE) in order to decrease the amount of long Tn conformation in the isomorphic crystal structure and achieve narrow SHLs, which are similar to those in the e-beam-irradiated P(VDF-TrFE) copolymers.5,11

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ASSOCIATED CONTENT

* Supporting Information S

Experimental section, molecular weight and molecular weight distribution data, 1H and 19F NMR spectra and peak assignments, DSC and 2D WAXD results of the uniaxially stretched P(VDF-TrFE-CTFE) film, DSC results of the uniaxially stretched and 80 °C annealed P(VDF-TrFE-CTFE) film, and 2D WAXD patterns for the uniaxially stretched and high temperature annealed P(VDF-TrFE-CTFE) and P(VDFCTFE-CFE) films. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Tel 01-216-3685861 (L.Z.). Present Addresses

B.A.T.: Department of Chemistry, Central Michigan University, Mount Pleasant, MI 48859. D.H. and Y.W.: Hudson High School, Hudson, OH 44236. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by National Science Foundation (DMR-0907580 and DMR-1402733). REFERENCES

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