Understanding of Relaxor Ferroelectric Behavior of Poly (vinylidene

Apr 13, 2015 - Guanyao Wang , Xingyi Huang , and Pingkai Jiang. ACS Applied Materials ... Feihua Liu , Zeyu Li , Qing Wang , Chuanxi Xiong. Polymers f...
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Understanding of Relaxor Ferroelectric Behavior of Poly(vinylidene fluoride−trifluoroethylene−chlorotrifluoroethylene) Terpolymers Matthew R. Gadinski, Qi Li, Guangzu Zhang, Xiaoshan Zhang, and Qing Wang* Department of Materials Science and Engineering, The Pennsylvania State University, University Park, Pennsylvania 16802, United States ABSTRACT: Relaxor ferroelectric poly(vinylidene fluoride) (PVDF) based terpolymers are attracting tremendous interest because of their potential applications in advanced energy harvesting and storage devices. Fundamental understanding of the ferroelectric behaviors of poly(vinylidene fluoride) (PVDF) based terpolymers has proved elusive. Current research suggests that the existence of different hysteresis loops results from physical pinning of the ferroelectric domains by the bulky defect monomers and that the size of the defect monomer determines the ferroelectric behavior. In this study, a poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene) random terpolymer is processed using a variety of methods and found to exhibit normal ferroelectric, single hysteresis loop (SHL), and double hysteresis loop (DHL) behaviors depending on the processing method. This indicates that the ferroelectric behavior of the terpolymer is related to not only the size of an individual defect unit but also how they are arranged within the relaxor ferroelectric phase. The results show that DHL behavior is a result of paraelectric domains that are promoted by long crystallization times, while the SHL behavior stems from a more random dispersion of these defects.

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INTRODUCTION Since the discovery of relaxor ferroelectric behavior in e-beam and γ-irradiated poly(vinylidene fluoride-co-trifluoroethylene), P(VDF−TrFE),1−3 and terpolymers of P(VDF−TrFE) with bulky comonomers such as chlorofluoroethyelene (CFE)4 and chlorotrifluoroethylene (CTFE),5,6 increasing interest has been placed on understanding the fundamental mechanisms behind these behaviors due to their relevance to applications such as electrostrictive actuators,1−3 electric energy storage,4,7−9 and electrocaloric cooling.10,11 The attractiveness of these materials arises from their higher dielectric constant, e.g., K > 50, relative to most polymeric materials (K ∼ 2−3), and slimmer hysteresis loops when compared to normal ferroelectric materials.12 These polymers are found to exhibit both narrow single hysteresis loops (SHLs) and double hysteresis loops (DHLs) depending on the nature of the bulky defect introduced into the polymer structure.13,14 While SHL and DHL behaviors have been observed in ceramic materials for decades, and their fundamentals are well documented and understood, the mechanism for these behaviors in PVDF based ferroelectric polymers is quite different. More specifically, the differing ferroelectric behaviors in ceramics can be attributed to the presence of specific crystalline phases and atomic displacement within the lattices,15 while for ferroelectric polymers only a single crystalline phase is observed for both behaviors with the polarization arising from defect enabled mobility of molecular dipoles along the polymer backbone.13 From current understanding, the introduced defects must be included in the crystal phase to expand the interchain spacing and reduce the steric hindrance to dipole motion in order to achieve a relaxor ferroelectric state in these © 2015 American Chemical Society

polymers. An initial expansion of this dimension is achieved through introduction of the TrFE defects, which enables further expansion induced by the third bulky defect.3,13 Without the TrFE defect, the copolymers of PVDF and CTFE, for example, only show limited CTFE inclusion and normal ferroelectric behavior.16,17 The second requirement is that the bulky defect must then reduce the size of ferroelectric domains to nanoscales to reduce the cooperative polarization that leads to the large remnant polarization shown in the P(VDF−TrFE) copolymers.13 In essence, what is formed is a highly defected nonpolar crystalline phase with dispersed, discrete regions of all-trans chain segments induced by the presence of TrFE defect and broken up by the larger defect. The formation of a SHL or DHL, however, is dependent on the nature of the third bulky defect. For example, e-beam and γirradiated P(VDF−TrFE) copolymers exhibit a narrow SHL.1,13 This behavior is thought to be achieved through internal crosslinking of the crystalline phase. Further evidence that the crosslinking occurs within the crystals is the re-emergence of ferroelectric behavior after recrystallization of irradiated films.13,18 In this way the cross-links expand the lattice dimension and reduce the size of ferroelectric domains but also chemically “pin” the ferroelectric domains.13 This allows for easy rotation of the dipoles on the polymer chains between these pin points with an application of electric field and promotes their relaxation to an unpolarized state upon cessation of applied field. As a result, a narrow SHL is observed. Received: January 27, 2015 Revised: March 17, 2015 Published: April 13, 2015 2731

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Figure 1. High field P/E loops of (a) TerQ, (b) TerAC, (c) TerSCAn, (d) TerQS, (e) TerACS, and (f) TerSCAnS. TerQ, TerAC, and TerSCAn are applied at a field up to 200 MV/m with the stretched versions shown up to 300 MV/m due to differences in breakdown strength. All measurements were carried out at 10 Hz.

In comparison, P(VDF−TrFE−CFE) terpolymers have been found to exhibit DHL behavior.13 Similarly to above, a relaxor ferroelectric state is induced in the material through the incorporation of CFE into the crystalline phase enabled by the prior expansion by the TrFE defect. In this case, however, the CFE defects result in a purported physical pinning of the

ferroelectric domains compared with the permanent, chemical pinning resulting from cross-linking. As the pinning in P(VDF−TrFE−CFE) is physical in nature and the CFE defects themselves also have a dipole, though weaker than in the ferroelectric domains, they will also orient under a sufficient applied field.13 The DHL is then a consequence of the relaxor 2732

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Macromolecules ferroelectric state being initially stable at low fields, but at higher applied fields a reversible transformation occurs to a ferroelectric state resulting from large ferroelectric regions nucleating due to the CFE rotation. DHL behavior has also be reported in polystyrene side chain grafted P(VDF−TrFE− CTFE) terpolymers,19 which is ascribed to the nanoconfinement effect due to wrapping of the crystals by the poorly polarizable side chains effectively reducing the cooperative polarization of the ferroelectric domains.13−19 It was originally thought that this DHL behavior would then be inherent of all the P(VDF−TrFE) based terpolymers due to the physical nature of the pinning. More recently, however, P(VDF−TrFE− CTFE) terpolymers have been reported to also exhibit SHL behavior.14 This has been attributable to the larger CTFE in comparison to CFE which acts as a superior physical pin preventing the formation of the ferroelectric phase even at high fields.14 Ultimately, the most recent explanations suggest that only the nature (chemical/physical) and size of the third defect determine the hysteresis loop behavior of the relaxor terpolymer. In this study, in order to shed light on the mechanisms of ferroelectric behavior, a P(VDF−TrFE−CTFE) terpolymer subjected to a variety of processing methods has been fully characterized by DSC, XRD, FTIR, and dielectric spectroscopy. The polarization of the terpolymers has been measured as a function of the applied electric fields. It was found that the terpolymer crystallized from melt could exhibit both normal ferroelectric and DHL behaviors depending on the time given for crystallization. The melt-crystallized films were then uniaxially stretched and observed to exhibit SHL behavior. To our knowledge, no P(VDF−TrFE) terpolymer has been reported to exhibit all of these behaviors as a result of processing though most literature only examines these materials from a limited processing standpoint, i.e., effect of annealing temperature or time, stretching, and then annealing samples, often solution cast. Our finding suggests that it is not only the nature and size of the bulky defect that determine the behavior of these relaxors but also how these defects arrange within the crystalline phase determined by processing. By comparing the structural differences that arise from these processing methods in terms of defect inclusion, crystalline phase and dimension, chain conformation distributions, thermal properties, and orientation, a better understanding of the role of the bulky defect in determining hysteresis loops in these terpolymers will be achieved, which will help to design and tailor this unique class of electroactive materials for advanced devices.

0.023 C/m2 at 200 MV/m (see Figure 2). With longer crystallization time, the loops of TerAC and TerSCAn become

Figure 2. Comparison of maximum polarization and remnant polarization of melt pressed and stretched samples at 200 MV/m.

slimmer and display the DHL behavior. These will often be referred to as the melt crystallized samples. All stretched samples show narrow SHLs characteristic of strongly pinned relaxor ferroelectrics but are clearly affected by the starting films and become progressively more narrow from Figure 1d−f. It is important to note that the processing of the initial film affects the polarization behavior of the stretched film. Because of the reduced modulus of these materials in comparison to the homopolymer or copolymers, these materials are stretched at room temperature or at slightly elevated temperature. Consequently, there is limited local melting of crystallites so that a remnant of the original structure will be retained. The maximum polarization and remnant polarization of each sample are compared in Figure 2 at 200 MV/m applied field. In both the melt crystallized and stretched sets of data the quenched samples are found to have the highest remnant polarization followed by the air-cooled samples with the lowest exhibited by the slow cooled and annealed samples. Interestingly, the trend of the maximum polarization to remnant polarization differs on whether the sample exhibits either SHL or DHL behavior. For melt crystallized samples, as remnant polarization is decreased the maximum polarization is found to increase. Whereas, in the SHL systems the maximum polarization seems to reach a maximum in TerACS but begin to drop as the remnant polarization is decreased further in TerSCAnS. The difference in this trend is an indication that the DHL behavior is a result of a different mechanism than that in the SHL behavior. This provides valuable insight into the nature of these behaviors and defects that they result from and will be further elaborated on in the structural characterization. To further study these behaviors, temperature-dependent dielectric spectroscopy was employed. Plots of the real (ε′) and imaginary (ε″) permittivity are shown in Figure 3a−f for each of the samples. TerQ of Figure 3a shows two main features. The first is a low-temperature peak seen at −14 °C in the ε″ plot at 100 Hz that is assigned to the glass transition (Tg) of the polymer, which progressively shifts to higher temperature and becomes broader with increasing frequency.13,14 The second is a high-temperature peak that is more obvious in the ε′ data set due to the peak being obscured by conduction at low frequencies and merging with the Tg peak at high frequencies in ε″. This peak is a result of the ferroelectric−paraelectric (FE−PE) Curie transition located at ∼40 °C.13,14 Because of

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RESULTS AND DISCUSSION High-Field Polarization Behavior and Weak-Field Dielectric Spectroscopy. Six samples were prepared from the P(VDF−TrFE−CTFE) (78.1/16.5/5.4 mol %) terpolymer using varied processing conditions. The first three included films that were quenched, air cooled, and slow cooled and annealed from melt. These were designated as TerQ, TerAC, and TerSCAn, respectively. Samples of each processing method were then uniaxially stretched, and were labeled as TerQS, TerACS, and TerSCAnS. The specific processing conditions of each sample are described in the Experimental Procedures. The polarization behavior of each sample was characterized by polarization versus field (P/E) loops as shown in Figure 1a− f. As we can see from Figure 1a, the quenched terpolymer exhibits a normal ferroelectric behavior characterized by a broad polarization loop and high remnant polarization, i.e., D = 2733

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Figure 3. ε′ and ε″ as a function of temperature for (a) TerQ, (b) TerAC, (c) TerSCAn, (d) TerQS, (e) TerACS, and (f) TerSCAnS.

transition relaxation, whose onset is the same as in the quenched sample, however, with a lower temperature Curie transition that is attributed to the relaxor ferroelectric− paraelectric (RFE−PE) transition located at ∼20 °C.4,5,13,14 These relaxations merge at higher frequencies and result in the rapid increase in permittivity, but can be seen at 100 Hz and 1 kHz in the ε″ of Figure 3c. In Figure 3f, however, the lowtemperature peaks shifts to ∼−5 °C, likely resulting from an

the reduction in the ferroelectric domain size, the temperature of this transition is lowered from 60 to 80 °C as is found in the P(VDF−TrFE) copolymer (depending on composition).13,21 The primary feature of this peak is the relative invariance with frequency. On the opposite end of the spectrum both TerSCAn and TerSCAnS samples exhibit the reported relaxor behavior as shown in the irradiated P(VDF−TrFE) copolymer and terpolymers.1,13,22 This behavior is a result of the glass 2734

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at low temperature (≤∼40 °C) and another at high temperature (>125 °C). TerQ shows a single relatively sharp peak at low temperature, whereas TerAC shows the same main peak shifting to slightly lower temperature and the emergences of a shoulder off of this main peak centered at ∼30 °C. This peak, similar to the temperature-dependent dielectric spectroscopy, is a result of the FE−PE Curie transition.13 With increased crystallization time, this peak begins to shift to lower temperature, and with sufficient time, i.e. TerSCAn, this peak becomes very broad, decreases significantly in intensity, is centered around 20 °C, and is attributed to the RFE−PE transition.13 Concurrent with this shift in Curie transition temperature is the growth of a shoulder which begins to form off the main crystalline melting peak centered at 130 °C in the melt crystallized samples. This shoulder which progressively shifts to be centered at 140 °C in TerSCAn could be a result to two crystalline phases, though, as discussed below in the XRD data, is not supported. The lower temperature peak has previously been attributed to crystals that formed during the cooling process which is likely the case here.23 However, the higher temperature shoulder has in the past been attributed to crystals formed during annealing, which does not occur in either TerQ or TerAC, and yet the presence of the shoulder is observed though clearly weaker than in TerSCAn.23 For this terpolymer, the formation of this higher temperature shoulder occurs even when the film is quenched from melt, which is promoted with longer crystallization times. The stretched samples all show the presence of the FE−PE transition. Other than becoming broader, there is observed very little difference between the original films and the stretched versions for TerQS or TerACS in terms of this transition. It is also noted that neither TerQS nor TerACS shows any indication of the formation of RFE−PE transition. TerSCAnS still retains the presence of the RFE−PE transition though also is observed to develop a weak FE−PE peak. In terms of melting peaks, all the stretched samples show no indication of the hightemperature shoulder and become sharper, indicating that there is a reduced distribution of the crystalline sizes in the stretched polymers in comparison to the melt crystallized films. Table 1 contains the values for the melting and Curie transitions of the above samples along with the enthalpy values and calculated crystallinities. The crystallinity is found to increase with increasing crystallization time. After stretching, there is also a slight increase in crystallinity over the unoriented films. It is important to note that the dielectric relaxations of Figure 3 increase with increasing crystallinity, indicative of this behavior arising from defects in the crystalline phase as opposed to be a result of defects in the amorphous phase. To further evaluate the crystalline phase, 2-D XRD was employed

increased Tg due to orientation of the amorphous phase and an earlier merging of the relaxation peaks. The permittivity peak is also observed to shift to lower temperature. TerAC, TerQS, and TerACS samples are intermediates in behavior between the extremes of the quenched and annealed samples. TerQS and TerACS samples show very similar spectra and seem to have the same features as TerQ. The primary difference is that the Tg peak shifts to higher temperature, and permittivity, instead of plateauing as in TerQ, shows a rapid increase similar to the relaxor ferroelectric spectra before peaking at the obvious FE−PE transition. This behavior bears a striking resemblance to uniaxially stretched P(VDF−TrFE− CFE) though their polarization behaviors differ and will be furthered discussed below.13 Finally, TerAC is found to be somewhat unique in comparison to the other samples. Similar to TerQS and TerACS samples, the Tg peak and FE−PE transition are obvious, but instead of permittivity rapidly increasing, it is found to increase to more gently in comparison to what it expected for the RFE−PE transition. To explain these differences and determine the role of bulky defect in the observed behaviors, structural characterization was performed. Structural Characterization. The thermal transitions of the samples were examined by differential scanning calorimetry (DSC). Figure 4a−c shows the melting endotherms between

Figure 4. Melting endotherms of (a) TerQ and TerQS, (b) TerAC and TerACS, and (c) TerSCAn and TerSCAnS. The heating rate was 10 °C/min for all samples.

−35 and 150 °C. Samples TerQ, TerC, and TerSCAn are the first to be compared. These samples show two main peaks: one

Table 1. Curie Transition Temperatures, Melting Temperatures, Enthalpies, and Crystallinity of Terpolymer Samples

a

sample

TFE−PE (°C)

TerQ TerQS TerAC TerACS TerSCAn TerSCAnS

43.4 42.3 40.4 41.9 43.1

TRFE−PE (°C)

Tm (°C)

ΔHFE−PE (J/g) 4.7 3.3 2.9 2.9

17.7 14.3

125.2 130.0 127.7 130.8 126.6 128.5

0.9

ΔHRFE−PE (J/g)

ΔHm (J/g)

crystallinitya (%)

0.7 0.6

14.6 15.0 16.8 18.6 19.5 23.1

34.8 35.7 40.0 44.3 46.4 55.0

Crystallinity determined from ΔHf0 = 42.0 J/g.23 2735

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Macromolecules to study not only crystalline phase and dimension but also orientation. Figure 5a−f shows the 2-D diffraction spectra of the terpolymer samples from XRD analysis. TerQ, TerAC, and

Figure 6. 1-D XRD spectra of terpolymer samples integrated from 2-D spectra.

and ferroelectric crystalline phase. In TerAC, there is a substantial increase in the amount of the relaxor ferroelectric phase which dominates in TerSCAn. Interestingly, all the stretched samples exhibit the same 1-D spectra as their unstretched counter-parts only with broadened peaks. Table 2 summarizes the lattice spacing, crystalline phase content, and crystallite size calculated from the 1-D XRD spectra.

Figure 5. 2-D XRD spectra of (a) TerQ, (b) TerAC, (c) TerSCAn, (d) TerQS, (e) TerACS, and (f) TerSCAnS. Parts d−f indicate the stretch direction in relation to the patterns.

TerSCAn (Figure 5a−c) exhibit continuous diffraction rings, indicative of random orientation of crystallites in these samples as would be expected of melt crystallization. Samples TerQS, TerACS, and TerSCAnS due to the orientation induced by the uniaxially stretching show diffraction arcs centered about the equator of the samples. Being located at the equator is indicative of the chain direction of the crystallites being oriented with the stretch direction, which is shown in the figures.24 The orientation of the stretched samples was quantified using Herman’s orientation parameter ( f). This parameter is described by eq 1 where ⟨cos2(ϕ)⟩ is the average angle between a crystalline reflection and a reference angle in this case the stretch direction (eq 2), and Ihkl is the intensity of a given crystalline reflection ((110/200) reflection of relaxor ferroelectric phase) as a function of ϕ.25 f=

3⟨cos2(ϕ)⟩ − 1 2 π

⟨cos2(ϕ)⟩ =

Table 2. Crystalline Phase Content, Lattice Spacing, and Crystallite Size of Terpolymer Samples crystalline phase content RFE (%)

FE (%)

lattice spacinga,b (A)

crystallite sizeb,c (A)

TerQ TerAC TerSCAn TerQS TerACS TerSCAnS

59.5 82.2 100 57.3 81.3 100

40.5 17.8

4.8 4.8 4.8 4.8 4.8 4.8

148 194 334 98 131 174

42.7 18.7

a Calculated from Bragg’s law. bCalculated from (200, 110) PE phase reflection. cCalculated from Scherrer’s equation.

(1)

In Table 2, it is seen that the ferroelectric phase content of the melt crystallized samples decreases with increasing crystallization time up to TerSCAn where no ferroelectric phase is present. After stretching, there is seen to be no change in the phase comparing with the original sample, and there is no mechanically induced ferroelectric phase due to the uniaxially stretching. This has previously been observed in samples of P(VDF−TrFE−CTFE) and is contrary to the results in P(VDF−TrFE−CFE) where some ferroelectric phase is induced by stretching.13,14 However, the temperature of stretching is likely more important in this case than the bulky defect as the former example is stretched at a temperature above the FE−PE transition (as is this one) while the latter is stretched below.13,14 There is also no notable difference in the lattice spacing between the samples. An increase in the lattice spacing would indicate increasing levels of CTFE inclusion into the crystalline phase. It is surprising that, even with annealing, there is observed no considerable expansion in this dimension though similar results have been observed for P(VDF−TrFE− CFE).23 This also suggests that all samples appear to include similar amounts of CTFE, meaning the differences in the polarization behavior must result from how these defects are incorporated.

2

∫0 Ihkl(ϕ) cos (ϕ) sin(ϕ) dϕ π

∫0 Ihkl(ϕ) sin(ϕ) dϕ

sample

(2)

Herman’s orientation parameter can possess values between −0.5 and 1 depending on orientation with values of −0.5, 0, and 1 being indicative of perpendicular, random, and parallel orientation, respectively.25 Samples TerQS, TerACS, and TerSCAnS were found to possess f values of −0.127, −0.140, and −0.179, respectively. Since the amorphous contribution was not removed, these values also reflect an amount of unoriented amorphous phase, slightly reducing their values. Nonetheless, the samples exhibit similar levels of orientation with the highest degree exhibited by TerSCAnS. Figure 6 shows the 1-D XRD spectra integrated from the 2-D spectra. In both samples TerQ and TerQS, there are two peaks located at 18.4° and 19.5° 2θ. The higher 2θ peaks are attributed to the (200,110) plane of the ferroelectric phase.13,14,22,23 The lower peak arises from the same plane of the defected paraelectric phase, also referred to as the relaxor ferroelectric phase.13,14,22,23 In both quenched samples, there appears to be a fairly even amount of both relaxor ferroelectric 2736

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Macromolecules The primary difference between the samples is the difference in crystallite size indicated by the broadening and narrowing of the crystalline peaks most noticeably of the defected paraelectric phase. For the melt crystallized samples, the quenched sample is found to possess the smallest crystals due to the rapid temperature decrease from melt. With the slower temperature decrease of TerAC, there is found to be a slight increase in the crystallite size. There is found to be a substantial increase in this dimension upon annealing. After stretching, the crystallite size is decreased in comparison to their melt crystallized counterparts. Both TerQS and TerACS show an approximate 30% decrease in crystallite size with TerSCAnS showing a ∼50% decrease. Considering that no crystalline phase change occurs during stretching, and no change in lattice spacing, but a substantial decrease in crystallite size provides information into the primary effect of the stretching procedure. Because of the low-temperature stretching utilized, there is a limited amount of local melting of crystallites. Consequently, the existing crystallites in the samples are ripped apart. While not inducing the ferroelectric phase, this procedure may affect how the CTFE defects are distributed within the crystals. To study this, FTIR was utilized to examine the chain conformation distributions within the samples. Figure 7 shows the FTIR spectra of the terpolymer samples. Of note are the peaks located at 510, 614, and 1280 cm−1 which

Figure 8. Chain conformation distribution of terpolymer samples from FTIR.

increase with increasing crystallization time. As the TG chain conformation is characteristic of the paraelectric phase of the homopolymer, this could simply be a result of increasing defected paraelectric phase (i.e., relaxor ferroelectric) content, which is consistent with XRD data.26 However, the XRD shows that there is no change in relaxor ferroelectric phase content after stretching, though the TG content is found to decrease in TerAC and TerSCAnS. The effect of stretching for the quenched sample is slightly different and will be explained below. As there is no change in the crystalline phase, and just as importantly all the samples show similar CTFE inclusion by the constant lattice spacing, the difference in the TG content must then relate to how the CTFE defect is organized in the crystalline phase, and this determines the polarization behavior. For the melt crystallized samples, increasing crystallization time leads to increasing DHL character of the P/E loops. TerQ with little time to crystallize forms about equal amounts of relaxor ferroelectric and ferroelectric crystalline phases, determined by local amounts of the TrFE and CTFE; i.e., TrFE-rich regions form ferroelectric phase and CTFE-rich regions produce relaxor ferroelectric phase. Therefore, there is likely little inclusion TrFE in the paraelectric-like phase and vice versa for CTFE in the ferroelectric phase. Similar results have also been observed in P(VDF−TrFE−CFE) samples crystallized from solution at low temperature in terms of crystalline phase and dielectric spectrum.27 The application of field polarizes the ferroelectric domains whose coupling is not inhibited and irreversibly transforms the relaxor ferroelectric to the ferroelectric phases (with sufficient applied field) with normal ferroelectric behavior being observed. In TerAC this begins to change as the defected paraelectric (i.e., relaxor ferroelectric) phase begins to dominate the sample, though some remaining ferroelectric phase exists. With sufficient crystallization time, both TrFE and CTFE defects exist in the same crystalline phase with this beginning in TerAC and completed in TerSCAn. The coalescing of the crystalline phase also results in an increase in the crystallite size and total crystallinity. Along with this change, there is an increasing TG content in the samples with reduced Tm>4 and TTTG conformations. Concurrently, there is found to be a reduction in the TC beginning with formation of a low-temperature

Figure 7. FTIR spectra of terpolymer samples with the TTTG, TG, and Tm>4 conformations highlighted.

are characteristic of the TTTG, TG, and Tm>4 conformations (T = trans, G = gauche), respectively.23 These conformations are of interest as they are characteristic of the ferroelectric (Tm>4) and paraelectric (TG) phases and the manner of how the ferroelectric domains are broken up (G containing conformations).23 These can also be used to interpret how the defects are affecting the structure of the material as TrFE defects are known to stabilize the trans conformation and CTFE is known to induce G conformation in their respective copolymers with VDF.16,21 From inspection of Figure 7 the most noticeable difference is in the prominence of the TG peak. This peak increases with increasing crystallization time and is reduced in the stretched samples. Figure 8 shows the chain conformation distribution determined from the above peaks. In all samples, the higher trans containing conformations dominate and is expected due to the high TrFE content. The most notable difference, however, is the variation in the TG content in the samples. In the melt crystallized samples, the TG content is found to 2737

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this pinning mechanism, but to elaborate further on the nature of the physical pins. The results of this study indicate that DHL behavior is induced by the formation of TG dominated defected regions developing in the relaxor ferroelectric phase, indicated by DSC and FTIR. In essence, the pinning regions in this case are paraelectric-like defected regions which can undergo cooperative rotation with applied field. The SHL behavior is a result of the stretching procedure breaking up these regions and dispersing the pinning defects throughout the crystal, similar to how the defects are dispersed in the irradiated copolymer. This explanation is consistent with previous reports when considering the processing. Uniaxially stretched P(VDF− TrFE−CFE) shows the same temperature-dependent dielectric spectra as TerQS and TerACS, an undescribed amount of stretch induced ferroelectric phase, and a weak TG chain conformation peak.13 At room temperature and 10 Hz, the P/E loop resembles that of TerQS and increasing temperature to 50 °C results in a loop similar to TerACS.13 Once this film is annealed, there is an increased intensity of the TG chain conformation, the removal of the ferroelectric phase, and DHL comparable to TerSCAn is observed.13 The most recent literature report for the SHL observed in P(VDF−TrFE− CTFE) is stretched and annealed; however, the annealing is done under tension preventing the TG domain formation.14 Also, important to note that in the above comparison is that the terpolymers previously reported possess concentrations of TrFE and CTFE/CFE of approximately twice the concentration of the terpolymer reported here, indicating that the processing dependent relaxor ferroelectric properties are present over a wide range of compositions. Though more work is needed in understanding the effect of different defect monomers and the effects of regiodefects in the main chain to optimize these materials, this paper provides a general framework on the mechanism of these behaviors.

shoulder in TerAC and a complete shift by TerSCAn. These results suggest that the ferroelectric domains sizes are being reduced by the formation of paraelectric domains. This is supported by the presence of the high-temperature shoulder off the main melting peak with the paraelectric crystalline phase known to be a higher melting temperature phase compared with ferroelectric phase.26 Similar to TerQ this would be a result of local defect concentration differences within the crystal. The DHL behavior is then a result of a reversible, cooperative transformation of the paraelectric domains to ferroelectric domains, similar to the limited reversibility observed in P(VDF−CTFE) copolymers.16,28 Conversely, TerACS and TerSCAnS show a reduction in crystallite size and decrease in TG conformations, but no change in crystalline phase or lattice spacing. The stretching breaks up the large crystals formed during the melt crystallization, breaking up the paraelectric domains and inducing more trans containing conformations, in effect dispersing the defects more randomly throughout the crystal. As a result, there is no co-operative PE−FE transition and SHL behavior is observed. Since there is no change to the ferroelectric phase of TerQS, this behavior is likely just a result of reduced co-operative polarization due to the fact that crystallite size decreases as the loop does remain comparatively broad. This then explains the observed trends observed in Figure 2. For the DHL behavior in the melt crystallized samples, TerSCAn shows the highest maximum polarization as it contains the most paraelectric domains to undergo transformation to the ferroelectric state with the lowest remnant polarization as a result of improved reversibility of that transition. TerACS shows the highest polarization of the stretched samples being that it contains a larger amount of the ferroelectric chain conformations though the remaining ferroelectric crystalline phase results in higher remnant polarization than TerSCAnS. TerSCAnS TG content, though reduced in comparison to the nonstretched version, still contains a significant amount of this chain conformations, indicating that, possibly due to the large size of the crystals, stretching was less effective in dispersing the defects but sufficient to break up the paraelectric domains. As a result, there is a reduction in both maximum and remnant polarization due to no RFE−PE.

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EXPERIMENTAL PROCEDURES

Synthesis of Terpolymer. Synthesis of P(VDF−TrFE−CTFE) terpolymer was accomplished via suspension polymerization using a 300 mL stainless steel Parr reaction vessel. 100 mL of deionized water and 0.15 g of potassium peroxodisulfate initiator were added to the vessel which was subsequently sealed and degassed via vacuum pump and cooled using liquid nitrogen bath. Gaseous VDF, TrFE, and CTFE were pumped into the reaction vessel separately. VDF was pumped in at a rate of 523.6 cm3/min. CTFE and TrFE were pumped in at a rate of 385 cm3/min. The amount of monomer was controlled by varying the time each monomer was allowed to flow into the reaction vessel. After charging the vessel with monomer, the vessel was sealed. The vessel was then heated to 90 °C and stirred at 600 rpm for 24 h or until the vessel pressure became stable for 2 h. Once the reaction was complete, the polymer was washed by vacuum filtration with both distilled water and methanol and then dried at 90 °C for 24 h to yield about 15−20 g of which powder corresponding to ∼50% conversion. GPC results using a Viscotek TDA 302 with DMF (0.01 M LiBr) as eluent at 65 °C with a refractive index detector calibrated by universal calibration with PS standards show a Mn of ∼60 000 g/mol and a PDI of ∼1.9. The composition was determined as per previous report to be VDF/TrFE/CTFE 78.1/16.5/5.4 mol %.20 Film Processing. The polymer films were produced via the melt press method. The powder polymers were heated in hydraulic press to 200 °C. Once the desired temperature was achieved, the pressure was increased by 500 psi every 15 min up to 6500 psi. The films were left at pressure and temperature for a minimum of 2 h to ensure film uniformity and maximum thinness. Sample TerQ was immediately taken from the hot press and quenched in ice water bath. Sample TerAC was removed from the hot press and allowed to cool in air. Sample TerSCAn was left in press at 200 °C and cooled to 90 °C at a

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CONCLUSIONS Because of their relevance to several important applications, PVDF based relaxor ferroelectrics have been receiving increasing interest. Especially, in terms of understanding the fundamental mechanisms that result in this relaxor ferroelectric behavior and how to tune whether these materials will exhibit SHL or DHL behavior. This question is more of an issue with regards to terpolymers as compared with the irradiated P(VDF−TrFE) copolymers as these always show SHL behavior. The most recent publications explain this behavior by a pinning mechanism in which the size of the third bulky defect determines the polarization behavior through its effectiveness as a physical pin.13,14 Meaning the larger the defect, the more resistant it is to poling and prevents the double hysteresis loop behavior. In this paper, a P(VDF−TrFE−CTFE) terpolymer through processing was found to be able to exhibit normal ferroelectric, SHL, and DHL behavior. This indicates that the determination of these behaviors is dependent on more than the physical size of the third defect monomer. It is not our intention to dismiss 2738

DOI: 10.1021/acs.macromol.5b00185 Macromolecules 2015, 48, 2731−2739

Article

Macromolecules rate of 37 °C/h. The temperature was then maintained at 90 °C for 12 h. The film thickness varied between 15 and 20 μm. The films of each respective cooling method were then stretched over a 60 °C surface to a draw ratio of 5. Film thickness varied between 5 and 12 μm. These films were used in all structural characterization and dielectric measurements. For the dielectric measurements, the polymers were sputter coated with gold using a Denton Vacuum Desk IV sputter coater under an argon atmosphere at 50 mTorr with the instrument setting of 47% power for 125 s. The estimated electrode thickness was 30 nm. Characterization. Polymer compositions were determined by 19F and 1H NMR. 19F (1H) NMR spectra were obtained on Bruker CDPX-300 NMR (7T, 300 MHz) using a CFCl3 (TMS) internal standard. Samples were dissolved in deuterated DMSO and scanned 250 (32) times. The data were acquired using a 11.3 (12.1) μs pulse width, 1.0 (1.0) s relaxation delay, 7.4 (81.0) μs dwell time, 90° (90°) flip angle, 67567.57 (6172.84) Hz spectral window, 0.515 (0.094) Hz FID resolution, and 0.96998 (5.3085) s acquisition time. The size of processed data was 65536, which was set to half that of the total data. The spectral reference frequency was 282.13 (299.87) MHz. The spectra were processed utilizing no broadening factors and were both phase and baseline corrected. DSC curves were acquired using a TA Instruments model Q 100 DSC and a heating rate of 10 °C/min ramping from −40 to 150 °C. FTIR spectra were obtained on a Bruker V70 FTIR using an attenuated total reflectance mode with a Harrick MVP-Pro Star equipped with a diamond prism. 2-D XRD spectra were obtained on a Rigaku DMAX RAPID using a fixed geometry and exposing samples for 30 min. The radiation source was a Cu Kα source with a wavelength of 1.54 Å. 1-D spectra were obtained through integration of the 2-D spectra with respect to 2θ. Peak deconvolution was done using Jade XRD analysis software using either Gaussian or Pearson VII peaks, which ever resulted in the lowest ⟨R2⟩ value. The complex dielectric constant as a function of temperature was analyzed with a Hewlett-Packard 4284 LCR meter in conjunction with a Delta Design oven model 2300 equipped with liquid nitrogen cooling was utilized using a 2 V bias. High field polarization−electric field (P−E) loops were measured using a modified Sawyer−Tower circuit with a Trek Model 30/20 ± 30 kV high-voltage amplifier system. Measurements were performed in Galden HT insulation fluid using a triangular bipolar bias at 10 Hz.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (Q.W.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported in part by the Office of Naval Research Office. X.Z. acknowledges the fellowship provided by the China Scholarship Council (CSC).

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DOI: 10.1021/acs.macromol.5b00185 Macromolecules 2015, 48, 2731−2739