Achieving Relaxor Ferroelectric-like Behavior in Nylon Random

Dec 1, 2017 - High dielectric constant polymers, exhibiting relaxor ferroelectric (RFE) behaviors (i.e., slim single and double hysteresis loops, SHLs...
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Achieving Relaxor Ferroelectric-like Behavior in Nylon Random Copolymers and Terpolymers Zhongbo Zhang, Morton H. Litt, and Lei Zhu* Department of Macromolecular Science and Engineering and Department of Chemistry, Case Western Reserve University, Cleveland, Ohio 44106-7202, United States S Supporting Information *

ABSTRACT: High dielectric constant polymers, exhibiting relaxor ferroelectric (RFE) behaviors (i.e., slim single and double hysteresis loops, SHLs and DHLs), are attractive for high energy density and low loss dielectric applications. Utilizing the principle of nanosized ferroelectric domains (nanodomains), this study has designed and developed novel RFE-like polyamides (PAs) based on 11-aminoundecanoic acid, 12-aminododecanoic acid, and N-methyl-11-aminoundecanoic acid (NM11) as an alternative to the high-cost and difficult-to-synthesize poly(vinylidene fluoride) (PVDF)based RFE polymers. In the first attempt, quenched and stretched (QS) PA(11-co-12) copolymers exhibited enhanced ferroelectricity as compared with either nylon-11 or nylon-12. Although relatively narrow hysteresis loops could be achieved at 75 °C, the hydrogen-bonding interaction was not weak enough to induce nanodomains in the nylon copolymers. To further reduce the hydrogen-bonding interaction and achieve nanodomains, a PA(11-co-12-co-NM11) 30/60/10 (molar ratio) terpolymer (terPA-NCH3) was synthesized. The NCH3 groups were expected to participate in the isomorphic crystals, blocking the formation of hydrogen bonds and inducing chain twists in the mesophase. Indeed, the RFE-like behavior with slim SHLs and high dielectric constant (60−70) was successfully achieved for the QS terPA-NCH3 at high temperatures (>75 °C). Pathways to achieve RFE-like behavior for nylon-based polymers are discussed and compared with those for PVDF-based polymers. The knowledge obtained from this study can inspire potential applications for nylon polymers in advanced electrical and power applications.

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(DHLs) in addition to the slim SHLs.1 Below, we review current understanding of the RFE behaviors for crystalline polymers. First, nanodomains are generated via repeat-unit crystal isomorphism or defect modification.1,2,16 Fundamentally, the size of ferroelectric domains in polymers is closely related to the chain conformation and intermolecular interactions in the crystals. When aligned chains adopt an all-trans (or zigzag) conformation, the intermolecular interaction is strong; therefore, ferroelectric domains are usually large. Consequently, broad hysteresis loops are observed for ferroelectric polymers, such as the β-form of poly(vinylidene fluoride) (PVDF) and P(VDF-co-trifluoroethylene) [P(VDF-TrFE)] below the Curie transition temperature (TC).17 Twisted chain conformations are needed to decrease the intermolecular interaction and achieve nanodomains. When a bulky termonomer is introduced into P(VDF-TrFE) crystals (note that the TrFE content should be above 20 mol % to induce the low-temperature ferroelectric phase17), twisted chain conformations can be induced because of the steric effect. The strong dipolar interactions between the aligned zigzag chains will be effectively reduced.18,19 At an

INTRODUCTION Relaxor ferroelectric (RFE) behavior in novel ferroelectric polymers,1,2 exhibiting a slim single hysteresis loop (SHL), is attractive for many potential applications, such as electric energy storage,2−5 electrostrictive actuation,6,7 and electrocaloric cooling.8−10 Detailed domain structures have not been unraveled for RFE polymers. However, the fundamental principle for RFE ceramics has been better understood. Basically, nanoscale ferroelectric domains (or nanodomains) are responsible for the RFE behavior. One example is Pb(Mg1/3Nb2/3)O3 (PMN), whose ordered polar domains have an average size of only 2−3 nm and are randomly distributed within the disordered cubic-phase matrix.11,12 Because of significantly weakened cooperative domain−domain interactions and the high mobility of the diffuse domain walls, narrow SHLs are achieved with ultrahigh dielectric constants (εr around several thousand).13,14 It is considered that a similar nanodomain principle should also govern the RFE behavior of polymers; however, RFE polymers could show unique characteristics different from RFE ceramics because of their long-chain nature. For example, the antiferroelectric (AFE) behavior in ceramics is different from the RFE behavior because AFE ceramics possess a well-defined AFE crystalline phase.15 Although crystalline polymers do not possess well-defined AFE phases, RFE polymers can also exhibit double hysteresis loops © XXXX American Chemical Society

Received: October 20, 2017 Revised: November 21, 2017

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DOI: 10.1021/acs.macromol.7b02243 Macromolecules XXXX, XXX, XXX−XXX

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(εr ∼ 30) and nylon-6 (εr ∼ 60) at 100 °C. This is attributed to the electric-field-induced reversible transitions between the paraelectric state with random hydrogen bonds and the ferroelectric state with more or less parallel hydrogen bonds (i.e., chains have a more twisted conformation). Although the slim DHL behavior has been achieved for evennumbered nylons at high temperatures, a question still remains: Is it possible to further tune the crystalline structure to achieve slim SHLs for nylons? Learning from P(VDF-TrFE)-based polymers, this work describes the structure−ferroelectric relationships in polyamide (PA) copolymers from 11-aminoundecanoic acid (AUA) and 12-aminododecanoic acid (ADA) [PA(11-co-12)] with different compositions. Consistent with previous studies,41,42 PA(11-co-12) copolymers can form isomorphic crystals with intermediate crystalline structures. As a result of the built-in chemical heterogeneity, enhanced ferroelectricity is observed for the copolymers. However, the hydrogen-bonding interaction still appears too strong, and relatively broad hysteresis loops are obtained. To further weaken the hydrogen-bonding, a termonomer, N-methyl-AUA (NM11), is used to copolymerize with AUA and ADA because the NCH3 groups can block hydrogen-bond formation and potentially induce more twists in the chains [note, the PA(11co-12-co-NM11) terpolymer is abbreviated as terPA-NCH3]. Indeed, slim SHLs were eventually achieved for the terPANCH3 terpolymer at elevated temperatures. Finally, the current understanding of various ferroelectric behaviors of nylon-based polymers, i.e., broad SHLs, slim DHLs, and slim SHLs, is discussed and compared with those of PVDF-based polymers.

appropriate composition (7−9 mol %),20 the bulky termonomers lead to the formation of nanodomains via crystal isomorphism (i.e., inclusion of the termonomer units in the crystal). Meanwhile, they also serve as pinning points to prevent the growth of large ferroelectric domains with the alltrans conformation during electrical poling. When the pinning effect is weak such as in P(VDF-TrFE-CFE) terpolymers (CFE is 1,1-chlorofluoroethylene, dipole moment, μ = 1.8 D),1 the twisted conformation can transform into an all-trans conformation under a high enough poling field, forming larger ferroelectric domains. Upon bipolar poling, the reversible transformations between the twisted and the all-trans conformations result in DHLs. When the pinning effect is strong such as in P(VDF-TrFE-CTFE) terpolymers (CTFE is chlorotrifluoroethylene, μ = 0.64 D),21 the twisted conformation is stable, and slim SHLs are achieved. Electron beam irradiation of P(VDF-TrFE)1 generates chemical cross-linking within the crystals, leading to permanent pinning and a stable twisted chain conformation. As a result, slim SHLs are also observed. Because the synthesis of P(VDF-TrFE)-based terpolymers requires special facilities (explosion-proof autoclaves) and infrastructures (i.e., on-site production of TrFE), it is difficult to reduce its high cost for near-term commercialization. It is desirable to search for viable and cheaper alternatives that can also exhibit slim SHL and DHL behaviors. Polyamides (or nylons) were chosen because of their lower cost, easier synthesis, and better potential to tune the crystalline structure. For example, ferroelectricity has been demonstrated in oddnumbered nylons, such as nylon-11.22−29 However, because of a higher rigid dipole moment of 3.7 D for the non-hydrogenbonded amide group30 compared to that (2.1 D31) of the VDF unit, nylons exhibit their own uniqueness. First, nylons become nonferroelectric when the chains adopt an all-trans conformation with strong hydrogen-bonding interactions. An example is the α phase nylon-11.32,33 A quenched smectic-like mesophase phase with twisted chain conformations34 is needed to weaken the hydrogen-bonding interaction in order to achieve ferroelectricity, whether for odd- or even-numbered nylons.35 However, ferroelectric behaviors for odd- and even-numbered nylons are different and highly conformation-dependent. Because of the stable all-trans conformation with parallel amide groups, odd-numbered nylons tend to favor large ferroelectric domains upon electric poling, leading to broad hysteresis loops. On the contrary, the field-induced ferroelectric domains with parallel amide dipoles in even-numbered nylons are not stable, and thus their hysteresis loops appear to be slimmer than those of the odd-numbered nylons. Note that the ferroelectricity in even-numbered n-nylons largely originates from the crystalline mesophase, rather than from the amorphous phase as previously reported.36−39 Ferroelectricity in mesomorphic even-numbered nylons is different from that of PVDF-based polymers because it is induced by the electric field during the dynamic poling process. The ferroelectric domains are metastable. Upon removal of the poling field and annealing at high temperatures, ferroelectric domains largely disappear.40 Second, the hydrogen-bonding interaction at room temperature (RT) is still too strong to achieve nanodomains for mesomorphic nylons with twisted chains. High temperature is needed to induce more twisted chain conformations and further weaken hydrogen bonding in order to achieve nanodomains. From our recent study,40 slim DHLs with relatively high dielectric constants were observed for nylon-12

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EXPERIMENTAL SECTION

Materials. Methylamine (40 wt %, aqueous solution), 11bromoundecanoic acid (BUA, 99%), 11-aminoundecanoic acid (AUA, 97%), 12-aminododecanoic acid (ADA, 95%), and nylon-12 pellets were purchased from Sigma-Aldrich (St. Louis, MO). AUA and ADA were recrystallized from boiling deionized water for three times before use. Nylon-11 resin (Arkema Rilsan Besno TL) was purchased from PolyOne (Avon Lake, OH). All nylon resins were used as received without further purification, except for thorough drying in a vacuum oven (∼1 mmHg) at 70 °C for 3 days before hot-pressing. Synthesis of NM11. CH3NH2 aqueous solution (40 wt %) was added dropwise into a NaOH solution (30 wt %), and the resulting CH3NH2 gas was collected in a liquid nitrogen trap. 40 g (151 mmol) of BUA was dissolved in 100 mL of methanol, after which the solution was cooled down in liquid nitrogen with dry N2 purge. The liquid CH3NH2 (5.62 g, 181 mmol) was added to the solution, and the reaction vial was immediate sealed using a Teflon-lined cap. The reaction was carried out at room temperature for 5 days, after which the product was precipitated by adding 800 mL of acetone. The precipitate was washed with diethyl ether and recrystallized from hot water (90 °C) twice. Yield: 25.9 g (80%). Proton nuclear magnetic resonance (1H NMR) spectroscopy (600 MHz, D2O) δ: 2.90 (t, NHCH2, 2H), 2.58 (s, NCH3, 3H), 2.05 (t, HOOCCH2, 2H), 1.56 (m, NHCH2CH2, 2H), 1.43 (m, HOOCCH2CH2, 2H), 1.18 (m, other CH2 units, 12H). Copolymerization of AUA and ADA. Mixtures of AUA and ADA were polymerized at 220 °C for 6 h under dry N2 protection. The following monomer molar ratios of AUA:ADA were used for the random copolymerization with the sample designation given in parentheses: 87:13 (coPA-1), 67:33 (coPA-2), 57:43 (coPA-3), 50:50 (coPA-4), 47:53 (coPA-5), 33:67 (coPA-6), and 13:87 (coPA7). After cooling to RT, the PA(11-co-12) copolymers were obtained by carefully breaking the reaction ampules. Terpolymerization of AUA, ADA, and NM11. AUA, ADA, and NM11 were mixed together and then polymerized at 190 °C for 6 h with a monomer molar ratio of 0.3:0.6:0.1. After cooling to RT, the B

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Figure 1. First heating DSC thermograms for (A) quenched (Q) and (B) quenched and annealed (QA) nylon samples. The heating rate is 10 °C/ min. (C) and (D) show Tm and crystallinity as a function of the mol % of 12-aminododecanoic acid (ADA) in the nylon samples. Solid and dashed lines are used to show the trends for the Q and QA samples, respectively. Electric displacement−electric field (D−E) loop measurements were performed using a Premiere II ferroelectric tester (Radiant Technologies, Inc., Albuquerque, NM), in combination with a Trek 10/10B-HS high-voltage amplifier (0−10 kV ac, Lockport, NY). The applied voltage had a bipolar sinusoidal or triangular waveform in the frequency range of 10−1000 Hz. Silver (Ag) electrodes (2.5 mm diameter and ca. 50 nm thick) were evaporated onto both sides of the film samples using an EvoVac Deposition System (Angstrom Engineering, Inc., Kitchener, ON, Canada). The metallized films were immersed in silicone oil (Fisher 460-M3001) during D−E loop tests to avoid corona discharge. The temperature was controlled by an IKA RCT temperature controller (Wilmington, NC). A home-built sample fixture with high-voltage cables was used for connecting the electrodes on films to the interface of the Radiant ferroelectric tester.

PA(11-co-12-co-NM11) terpolymer was obtained by carefully breaking the reaction ampule. Film Fabrication and Processing. Four types of film samples were fabricated via melt-processing: (1) Quenched (Q) samples. Using aluminum (Al) foils, nylon samples were melted at a temperature about 30−40 °C above their melting temperatures (Tm), e.g., 210 °C for nylon-11 and nylon-12. After hot-pressing, the Al foil-sandwiched samples were immediately quenched into an isopropanol/dry ice bath at about −78 °C. (2) Quenched and annealed (QA) samples. The Q samples were annealed at 140 °C for 20 h in a vacuum oven. (3) Quenched and stretched (QS) samples. The Q samples were uniaxially stretched at RT to an extension ratio of ca. 300% using a home-built stretching apparatus. (4) Quenched, stretched, and annealed (QSA) samples. The QS samples were annealed either at 100 °C for 5 min (denoted as QSA@100 °C, no fixed ends) or at 140 °C for 20 h (with fixed ends). Instrumentation and Characterization. 1H NMR spectroscopy was performed on a Varian Mercury 600 MHz spectrometer with tetramethylsilane (TMS) as the internal reference. A trifluoroethanol/ CDCl3 (3/1 v/v) mixture was used as the solvent for the nylon polymers. Fourier transform infrared (FTIR) spectra were recorded using a Nicolet IS50R FTIR spectrometer (Thermo Fisher Scientific, Waltham, MA) in the transmission mode. The scan resolution was 2 cm−1 with 64 scans. Differential scanning calorimetry (DSC) was performed using a TA Instruments Q100 DSC at a scanning rate of 10 °C/min. Around 1.5 mg of sample was used to avoid any thermal lag. Two-dimensional (2D) wide-angle X-ray diffraction (WAXD) experiments were carried out using a Rigaku MacroMax 002+ equipped with a Confocal Max-Flux optic and a microfocus X-ray tube source operating at 45 kV and 0.88 mA. The X-ray wavelength was 0.1542 nm (Cu Kα). The WAXD patterns were collected using a Fujifilm image plate scanned by a Fujifilm FLA-7000 scanner at a resolution of 50 μm/pixel. One-dimensional (1D) WAXD curves were obtained by integrating the corresponding 2D WAXD patterns using the Polar software developed by Stonybrook Technology and Applied Research, Inc.

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RESULTS AND DISCUSSION

Chemical Heterogeneity-Stabilized Mesophase in Nylon Random Copolymers. According to our previous study,35 Q and QS nylon films can exhibit a mesomorphic crystalline phase. Several features are noted for this defective structure in Q/QS nylons. First, the chain conformation contains multiple twists per repeat unit; this is different from the cases of stable α or γ phases in n-nylons, where the trans (or zigzag) conformation dominates (note that there is a twist in the amide bond for the γ form). As such, nylon chains are poorly packed with enlarged interchain distances. Second, although the crystalline structure is more or less ordered along the chain direction (i.e., smectic-like), hydrogen bonds are randomly organized and perpendicular to the chains. Consequently, hydrogen bonding is weakened, leading to feasible ferroelectric switching upon high-field electric poling. Third, the quenched mesophase and the field-induced ferroelectric domains are metastable. Upon thermal annealing, the C

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Figure 2. FTIR spectra for (A) Q and (B) QA nylon films at room temperature.

defective due to the built-in chemical heterogeneity. In addition to the Tm and crystallinity results, DSC also provided information on the glass transition temperature (Tg). For the Q samples (Figure 1A), small hysteresis peaks in DSC were observed (see the magnified curves in the inset) because the samples were stored in a desiccator at RT for a while before the first heating at 10 °C/min (i.e., the physical aging effect45). The Tg values for most samples were around 38 °C. For the QA samples (Figure 1B), the hysteresis peaks largely disappeared, and the Tg values somehow increased to 45−50 °C. This could be attributed to better hydrogen bonding in the amorphous phase, crystallinity increase, and possibly a higher content of the rigid amorphous fraction (RAF)47 of the nylon films after thermal annealing at 140 °C for nylon films. The type 1 isomorphism in the PA(11-co-12) copolymers was further studied by FTIR. Figure 2A shows FTIR spectra for the Q nylon films. Since all Q samples should be in the mesomorphic phase, the absorption bands were broad and no obvious difference could be identified for various nylon samples. For example, the CH2 (CO and N vicinity) wagging or twisting bands around the amide III region (i.e., 1287−1192 cm−1)48 were broad and fairly similar for all copolymers. The only difference was the −CONH− in-plane vibration bands of the α-form nylon-11 at 938 cm−1 and the γform nylon-12 at 946 cm−1. With increasing the ADA content, this band gradually shifted from 938 to 946 cm−1. After annealing at 140 °C, the QA films exhibited sharper absorption bands (Figure 2B). In particular, typical absorption bands for the γ nylon-12 were observed at 1192, 946, and 627 cm−1,48 and those for α nylon-11 were seen at 1420, 939, 686, and 586 cm−1.49 The transition from the α-like nylon-11 to the γ-like nylon-12 was monitored as the mol % of ADA in the copolymers was increased. For example, the CH2 bending band [1466 cm−1 (N vicinity) for γ and 1416 cm−1 (CO vicinity) for α], the CH2 (CO and N vicinity) wagging/twisting bands (1287−1192 cm−1), the −CONH− in-plane vibration (∼940 cm−1), amide V (∼700 cm−1), and amide VI (∼600 cm−1) bands displayed transitions from the α-like nylon-11 to the γlike nylon-12 (Figure 2B). From this study, it was clear that the PA(11-co-12) copolymers, especially those with the ADA content around 40−70 mol %, seemed to have more disordered chain conformations due to the chemical heterogeneity in the isomorphic crystals. To study the ferroelectric properties, the Q nylon samples were uniaxially stretched to an extension ratio of 300% at RT,

hydrogen bonding in the mesophase is strengthened and the ferroelectricity is largely lost. To retain the disordered hydrogen bonds and thus the fieldinduced ferroelectricity for nylons, we propose to introduce chemical heterogeneity into nylons.40 It has been reported that crystalline isomorphism was found for many nylon copolymers.43 In this study, copolymerization of AUA and ADA was carried out at 220 °C at different molar ratios, following a previous report.41 From the 1H NMR results (e.g., see the 1H NMR spectrum of coPA-6 later), the copolymers had relatively high molecular weights because the end-group peaks of H2NCH2− at 2.91 ppm and HOOCCH2− at 2.16 ppm largely disappeared. We expected that chemical defects should be able to stabilize the mesophase by mismatching the amide groups in the crystals. The defective crystalline structures of Q and QA nylon polymers were first studied by DSC. The results are shown in Figure 1A,B, and the thermal data are summarized in Table S1 of the Supporting Information. From Figure 1C, the melting temperatures (Tms) were depressed for the nylon copolymers and reached a minimum with an ADA composition between 50 and 70 mol %. This is similar to the DSC results of Johnson et al.,41 and we attributed it to the defective crystal structure from type 1 repeat-unit isomorphism (i.e., a mixed crystal structure).43 The estimated crystallinities for Q and QA samples are summarized in Figure 1D. Here, the average heat of fusion (ΔH0f ) from 100% perfect nylon-11 (244 J/g) and nylon-12 (245 J/g) crystals is used for crystallinity calculation.44 However, we bear in mind that the ΔH0f for isomorphic crystals should be smaller than those of respective homopolymers because of the chemical heterogeneity in the crystalline structure. For example, the ΔH0f values for α PVDF and PTrFE are 104.5 and 66.3 J/g, respectively.45 The ΔH0f value for the isomorphic P(VDF-TrFE) (VDF/TrFE ∼ 70/30 mol./ mol.) was only 42 J/g.46 Because the crystalline peaks are broad and overlap with the amorphous halo in the WAXD profiles (see results later), it is difficult to obtain accurate crystallinity from the WAXD profiles for nylon samples. Therefore, the average ΔH0f was used for the estimation of crystallinity for PA(11-co-12) copolymers. From Figure 1D, thermal annealing increased the crystallinity for nylon homopolymers and PA(11co-12) copolymers with either AUA or ADA being the major component. For the copolymers with ADA between 40 and 70 mol %, thermal annealing did not increase the crystallinity. This result suggested that these isomorphic crystals were intrinsically D

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Figure 3. 1D WAXD profiles for the (A, B) QS and (C, D) QSA nylon films at RT along (A, C) the meridional and (B, D) the equatorial directions in the corresponding 2D WAXD patterns (see Figure S1 in the Supporting Information).

Figure 4. Continuous bipolar D−E loops at room temperature for QS (A) nylon-11, (B−G) PA(11-co-12) copolymers, and (H) nylon-12 films. The poling frequency is 10 Hz with a sinusoidal waveform.

exhibited a broad single reflection peak around 14.7 nm−1, corresponding to the pseudohexagonal mesophase with random hydrogen bonds. It was observed that the nylon copolymer films exhibited a slightly larger d-spacing (i.e., 0.425 nm) of the equatorial reflection than nylon-12 (i.e., 0.421 nm). This suggests that nylon copolymers adopted a more twisted chain conformation than nylon-12, leading to a slightly larger interchain distance. The 1D WAXD profiles of the QSA films along the meridional and the equatorial directions are shown in Figures 3C and D, respectively. The 2D WAXD patterns are shown in Figure S1B. As shown in Figure 3D, nylon-11 exhibited double reflection peaks along the equator direction [reminiscent of the (100)α and (010)α reflections], indicating that the mesophase had partially transformed into the α-like structure after thermal annealing. Similarly, the coPA-2 also displayed the double equatorial peaks, while all the other nylon copolymers and nylon-12 exhibited a single equatorial reflection peak. The

and their crystalline structures were investigated by 2D WAXD (see Figure S1A). The integrated 1D WAXD profiles along the meridional and the equatorial directions are shown in Figures 3A and 3B, respectively. Two major reflections of aliphatic nylons were observed; the low-angle reflection in the meridional direction from the smectic-like structure of hydrogen-bonded amide groups (Figures 3A,C) and the highangle broad reflection in the equatorial direction from the interchain packing (Figures 3B,D). Note that the mesophase of nylon-12 is reminiscent of the monoclinic (or pseudohexagonal) γ phase with the chains along the b-axis [i.e., the (020) reflection], while the mesophase of nylon-11 is reminiscent of the triclinic α phase with the chains along the c-axis [i.e., the (001) reflection].35 Upon increasing the ADA content, the corresponding d-spacing of the smectic reflections along the meridional direction gradually increased, indicating again that the comonomers were included into the isomorphic crystals. In the equatorial WAXD profiles (Figure 3A), all QS films E

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electric domains were poorly developed. These results indicated that the amide dipoles were unable to rotate easily under this poling condition because of the relatively strong hydrogenbonding interaction due to chemical regularity in the main chains. In contrast, all nylon copolymers exhibited broad ferroelectric loops under the same poling conditions (Figures 4B−G), especially the copolymers having 40−70 mol % ADA (i.e., coPA-3 to coPA-6). The enhanced ferroelectricity for the QS copolymers could be attributed to the poorly hydrogenbonded amide groups as a result of chemical heterogeneity. The ferroelectric behavior of the QS nylon films was also studied at 50 °C under the same poling conditions (10 Hz and 185 MV/m) (see Figure S2). Again, ferroelectric switching was enhanced, especially for coPA-3 to coPA-6. This observation is consistent with the results from DSC, FTIR, and WAXD; namely, the QS copolymers with 40−70 mol % ADA had the most defective crystalline structure with weakened hydrogen bonding due to chemical heterogeneity. It is known that the high-temperature annealed nylon-11 film was nonferroelectric because the ferroelectric switching of the strongly hydrogen-bonded amide groups in the α-like sample was significantly suppressed.35 Even for the QSA nylon-12 film, which exhibited enhanced ferroelectricity compared with the QSA nylon-11 film due to the twisted amide bond in the γ-like structure, there were no ferroelectric switching at 185 MV/m (Figure 5A). When the poling field increased to 265 MV/m, slightly opened loops were observed (Figure 5C). This should not be attributed to the ferroelectricity from the amorphous phase (e.g., proposed earlier for the Q nylon-636−39) because the QSA nylon-11 film did not show broadened loops under similar poling conditions, as we reported previously.35 Note that the amorphous phase of the QSA nylon-11 film supposedly should be nearly the same as that of the QSA nylon-12 film. Instead, the slightly broadened loops for the QSA nylon-12 under high poling fields (265 MV/m) should be attributed to the switching of poor ferroelectric domains in the mesophase. In contrast, the QSA copolymers exhibited much enhanced ferroelectricity because chemical heterogeneity stabilized the mesophase structure. For example, after annealing, the maximum D (Dmax) at 185 MV/m decreased from ca. 60 mC/m2 for the QS coPA-5 (Figure 4E) to ca. 35 mC/m2 for the QSA coPA-5 (Figure 5B). The decreased Dmax could be ascribed to the enhanced hydrogen bonding and improved crystalline structure due to thermal annealing. However, upon poling at 265 MV/m, the Dmax returned to nearly 60 mC/m2 (Figure 5D). It is inferred that the chemical heterogeneity helped to retain the mesophase structure to a large extent for enhanced ferroelectricity in the QSA copolymers. Achieving RFE-like Behavior in a Nylon Terpolymer at High Temperatures. Although enhanced ferroelectric behaviors have been demonstrated, it is still difficult to achieve the RFE-like behavior with slim SHLs at RT for the QS and QSA copolymers. Here, the design rules to achieve the RFE-like behavior are proposed. First, hydrogen bonding in the mesophase of PA(11-co-12) copolymers is still too strong to prevent the formation of large ferroelectric domains. According to our recent work,40 high temperature could be used to weaken hydrogen bonding by introducing more twisted chain conformations and to decrease the domain size for evennumbered nylons. Indeed, much narrower D−E loops could be achieved for the QS nylon copolymers at high temperatures (e.g., see the D−E loops at 75 °C for the QS coPA-6 in Figure 10E later). However, judging from the loop shape, the electrical

pseudohexagonal mesophase of the nylon samples became better packed due to thermal annealing, as reflected by the higher q values of the equatorial reflections for the QSA samples (Figure 3D). Similar to the QS nylon films, all QSA copolymers exhibited a larger interchain spacing (0.418 nm) than that of nylon-12 (0.414 nm), further confirming the defect-induced twisted chain conformation (which should be stable upon thermal annealing at 140 °C). For nylon-12 and coPA-7, a shoulder peak appeared at 3.85 nm−1 [i.e., the position for the (020)γ reflection; see Figure 3C], suggesting that a fraction of the samples had transformed into the γ phase after thermal annealing. On the basis of the above studies of the chain conformation and crystalline structure of PA(11-co-12) copolymers, several conclusions are summarized. First, QS and QSA copolymers contain chemical defects in the isomorphic crystals, which reduces the overall strength of hydrogen bonding. Second, the chain conformations are disordered with more twists in the backbones, which lead to a slightly larger interchain distance for the copolymers. As a result, the hydrogen-bonding interaction is further weakened. We expect that the dipole and domain flipping upon electric poling should become easier, and thus ferroelectricity is expected to enhance for PA(11-co-12) copolymers. Enhanced Ferroelectric Switching in PA(11-co-12) Copolymers. Ferroelectric behavior of the QS nylon films was studied by bipolar D−E loop tests. Continuous loops were used to minimize the remanent polarization from a previous loop.50 Supposedly, there should be hardly any ferroelectric domains in the unpoled QS samples because of the randomly arranged hydrogen bonds in the quenched mesomorphic phase. By applying five continuous loops at a relatively high frequency (i.e., 10 Hz) under a moderate field (i.e., Tg), ac electronic conduction became significant for the terPANCH3 film due to the enhanced chain motion. The ac electronic conduction needed to be subtracted from the experimental loops to reveal the dipole and domain switching (the subtraction procedure is given in Figure S5). Upon increasing the temperature above the Tg, the hysteresis loops

became slimmer (Figures 9B−D). Again, this could be explained by the weakened hydrogen bonding due to increasingly twisted chains in the crystal and thus reduced ferroelectric domain size at high temperatures.40 The breadth of the D−E loops in Figures 9A−D is reflected in the Pr and EC values, summarized in Figure 9E. For the QS terPA-NCH3 film, the Pr and EC gradually decreased with increasing temperature. These values were lower than those of the QS coPA-6 film and were nearly the same as those of the RFE P(VDF-TrFE-CTFE) terpolymer reported in the literature (Figure 9E).21 The apparent dielectric constant (defined by the slope in the ±30 MV/m region for slim D−E loops) was as high as 70 at 75 °C (10 Hz) and 60 at 95 °C (50 Hz). Judging from the slim D−E loops and high εr values, the observed dielectric behavior of the QS terPA-NCH3 film at high temperatures should be ascribed to the RFE-like behavior, rather than the paraelectric behavior (εr = 28) reported for nylon-12 at 100 °C.40 Study of the RFElike behavior above 100 °C was prevented by easy dielectric breakdown due to high electronic conduction.55 Nonetheless, this RFE-like behavior of the QS terPA-NCH3 film is different from the genuine RFE behavior of P(VDFH

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Figure 9. (A) Two sets of five (2 × 5) continuous bipolar D−E loops for the QS terPA-NCH3 film at room temperature. The time interval between two sets of runs was ca. 30 s. The poling field was 185 MV/m, and the poling frequency was 10 Hz with a sinusoidal waveform. (B−D) Two continuous bipolar D−E loops for the QS terPA-NCH3 film at (B) 50 °C, (C) 75 °C, and (D) 95 °C, respectively. These loops were obtained by subtracting the ac electronic conduction from the raw data (the procedure is given in Figure S5). The poling frequency was 10 Hz for (B, C) and 50 Hz for (D) with a triangular waveform. (E) Pr and Ec as a function of temperature for terPA-NCH3 and coPA-6. (F) Two sets of five (2 × 5) continuous bipolar D−E loops for the QS terPA-NCH3 film at 75 °C. The time interval between two sets of runs was ca. 30 s. The maximum poling field was 100 MV/m at a poling frequency of 10 Hz with a sinusoidal waveform.

Figure 10. (A) Poling field profile for the lifetime test by using time-dependent D−E loop measurement. Films were prepolarized at 170 MV/m, 500 Hz, and 75 °C for 2 × 5 continuous cycles before the lifetime test. (H) −Pr as a function of time for QS coPA-6 and terPA-NCH3 films at 75 °C. Continuous bipolar D−E loops of the QS films of (C−E) terPA-NCH3 and (F−H) coPA-6 at 75 °C and different frequencies (a triangle waveform): (C, F) 10 Hz, (D, G) 100 Hz, and (E, H) 500 Hz. The poling field was 170 MV/m. All the loops were obtained by subtracting the ac electronic conduction. The detailed subtraction procedure is given in Figure S6.

domains were gradually generated as the poling field increased from 20 to 100 MV/m. The εr was only about 30 for the first linear loop, indicating an intrinsic paraelectric phase with no ferroelectric domain for the QS terPA-NCH3 film. During the second set of five (red) loops at 100 MV/m, nanodomains fully developed and the εr reached 70. Therefore, the RFE-like behavior of the QS terPA-NCH3 film at high temperatures was

TrFE-CTFE) terpolymers because the ferroelectric nanodomains were induced by the electric poling, rather than being an intrinsic property of the RFE phase in P(VDF-TrFECTFE) crystals that contain stable nanodomains without any external electric field. This was evidenced by the electric poling under a lower field of 100 MV/m at 75 °C (Figure 9F). During the first five continuous (black) loops, ferroelectric nanoI

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Figure 11. Understanding ferroelectricity in crystalline polymers by comparing PVDF- and nylon-based polymers, where intermolecular (i.e., dipolar or hydrogen bonding) interactions and chain (trans versus twisted) conformations play an important role. (A) Biaxially oriented PVDF,60 uniaxially stretched (B) P(VDF-TrFE-CFE) and (C) P(VDF-TrFE-CTFE) films,2 (D) QA nylon-11 and QSA nylon-12 film,40 (E) QS nylon-12 film at RT,40 (F) QSA coPA-6 film at RT, (G) QS nylon-12 film at 100 °C,40 and (H) QS terPA-NCH3 film at 95 °C (50 Hz). The poling frequency for (A−G) was 10 Hz.

obtained (see Figures 10E,H). We consider that the fieldinduced ferroelectric domain size for QS terPA-NCH3 should be smaller than that of QS coPA-6 because its D−E loops were slimmer. After annealing the QS terPA-NCH3 film at 100 °C for 5 min followed by cooling to RT, the QSA@100 °C film exhibited much reduced ferroelectricity (see Figure S7) compared to the QS film (Figure 9A). This could be attributed to the better developed hydrogen bonding upon thermal annealing at 100 °C. However, after returning to 75 °C (or above), the RFE-like behavior could be recovered after several poling cycles at 100 MV/m or above. Similar results were observed for PA(11-co12) copolymers and also reported for the QSA@100 °C nylon12 film.40 Discussion of Ferroelectric Behaviors in PVDF- and Nylon-Based Polymers. Combining the understanding of the ferroelectric behaviors observed in PVDF- and nylon-based polymers, the fundamental physics that governs ferroelectricity in long-chain polymer crystals can be rationalized (Figure 11). Basically, intermolecular interactions (i.e., dipolar or hydrogen bonding) in the crystals play an important role for the ferroelectricity of polar polymers. The strength of intermolecular interactions is closely related to the dipole moment of the polar groups and the chain conformation in polymer crystals. PVDF and its copolymers/terpolymers are used as first examples to illustrate this.17,56,57 In the β phase of PVDF, the rigid dipole moment is relatively high, about 2.1 D (note that the overall μ ∼ 3.0 D/repeat unit, when taking into account the reactive interaction between rigid dipoles and additional polarization of bonded electrons58,59). Because of the all-trans

also induced by the poling electric field and is not associated with any well-defined Curie transition. Lifetimes of field-induced, metastable nanodomains for QS terPA-NCH3 and coPA-6 films were evaluated by using timedependent D−E loop measurements at 75 °C. Before the lifetime test, film samples were prepolarized at 170 MV/m and 500 Hz for 2 × 5 continuous cycles. The time interval between the prepoling and the lifetime test was ca. 30 s (i.e., the time for the ferroelectric tester to set up the next run). The poling field profile for the lifetime test is shown in Figure 10A. After one cycle of sinusoidal poling at 500 Hz and 75 °C, the electric field was kept at zero for additional 20 ms. From Figure 10B, the Pr values for QS coPA-6 and terPA-NCH3 were −16.4 and −9.4 mC/m2, respectively. The larger Pr for coPA-6 than that for terPA-NCH3 was ascribed to its higher ferroelectricity. After 10 ms, both values decreased to ca. −1.4 mC/m2, indicating the lifetime for both samples should be around 10 ms at 75 °C. Such a short lifetime had a significant impact on the hysteresis of D−E loops under different poling frequencies. As the poling frequency increased from 10 to 500 Hz, the D−E loops of both QS terPA-NCH3 (Figures 10C−E) and coPA-6 films (Figures 10F−H) became broader (subtraction of ac electronic conduction for these high-temperature loops is illustrated in Figure S6). Basically, when the frequency was low, the induced domains could relax as the poling field decreased to zero. As such, slim hysteresis loops were obtained (see Figures 10C,F). At high frequencies, especially when the poling time was similar to or shorter than the lifetime (∼10 ms), the fieldinduced domains could not relax before the poling field decreased to zero. Consequently, relatively broad loops were J

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enlarged. As a result, dipole switching upon high-field electric poling becomes feasible, and the ferroelectric behavior is obtained for both odd- or even-numbered nylons. An example of broad ferroelectric loops is shown for the QS nylon-12 (Figure 11E). Supposedly, the unpoled mesomorphic nylons do not contain any ferroelectric domains. The ferroelectric domains are actually induced by the high-field poling. For odd-numbered nylons such as nylon-11, the resulting polar crystal structure in the generated domains stabilizes the ferroelectric behavior.40 It is likely that the TC is above the Tm for nylon-11. On the contrary, for even-numbered nylons such as nylon-12, the generated ferroelectric domains are not stable due to the torsional stress in the chains when the neighboring amide groups are forced by the electric field to orient in the same direction. As a result, the ferroelectric domains for mesomorphic even-numbered nylons are somewhat short-lived. Basically, ferroelectricity in mesomorphic even-numbered nylons only should be dynamic ferroelectric behavior during electric poling and does not relate to any polar crystalline structures in the unpoled samples. Therefore, evennumbered nylons do not possess a well-defined, reversible Curie transition like that in PVDF-based copolymers and terpolymers. To further reduce the intermolecular interaction in mesomorphic even-numbered nylons, high temperature can be utilized because the chains in the nylon crystal become more twisted at high temperatures. 40 However, the twisted conformations are not very stable upon electric poling. At a high enough electric field, the twisted conformation in evennumbered nylons can transform into the more or less trans conformation, resulting in transient ferroelectric nanodomains. The reversible paraelectric-to-ferroelectric transitions during bipolar poling lead to DHLs; an example of the QS nylon-12 at 100 °C is shown in Figure 11G. Chemical defects can be introduced into nylon crystals by using random copolymers to stabilize the twisted chain conformation, such as in PA(11-co-12) copolymers. For example, the D−E loops for the QSA coPA-6 at RT are shown in Figure 11F. Note that this is contrary to the case in the QSA@100 °C nylon-12 film (blue loops in Figure 11D), which does not exhibit ferroelectric switching at 200 MV/m due to the strengthened hydrogen bonding after thermal annealing at 100 °C for only 5 min.40 To further decrease the intermolecular interaction in nylon copolymers, both terpolymerization with the bulky NN11 and high temperature are utilized. The NCH3 groups are expected to be included in the defective mesomorphic crystals, stabilizing the twisted conformation/pinning the crystals and making hydrogen bonding even weaker. As shown in Figure 11H, the QS terPA-NCH3 indeed exhibits slim SHLs at ≥75 °C, which is a characteristic feature for the RFE-like behavior. This RFE-like behavior again is attributed to the field-induced nanodomains in the defective terPA-NCH3 crystals with stable twisted chain conformations.

chain conformation in the crystal that maximizes the intermolecular dipolar interaction, β PVDF favors large ferroelectric domains and tight chain packing, which results in broad hysteresis loops (Figure 11A).60 After copolymerizing with comonomers with a smaller dipole moment, such as TrFE (μ = 1.05 D) or tetrafluoroethylene (TFE, μ = 0 D), the intermolecular dipolar interaction among the all-trans chains in the low-temperature ferroelectric phase is decreased. As a result, both P(VDF-TrFE) and P(VDF-TFE) random copolymers exhibit enhanced ferroelectricity with easier domain switching. However, the intermolecular interaction is still strong, and the ferroelectric domains are large enough to result in broad hysteresis loops. To further weaken the intermolecular interaction, twisted chain conformations need to be generated. On one hand, the twisted conformation can prevent neighboring dipoles from lining up into a long-range interaction. On the other hand, the twisted chain conformation can increase the interchain distance and facilitate the rotation of polymer chain segments in the crystal. The twisted chain conformation can be realized by copolymerizing a relatively large termonomer such as CFE or CTFE with VDF and TrFE. Because of the crystal isomorphism, the CFE and CTFE units are included in the crystal as defects. Because they disfavor the all-trans conformation due to the steric effect, twisted chain conformations are adopted. Meanwhile, they serve to pin other dipoles along the chain in the crystal. There are two scenarios observed for the twisted chain conformations. First, in the case of P(VDF-TrFE-CFE),1 the twisted chain conformation is not stable under high-field poling, and it can be polarized at a high enough electric field into a more trans conformation to form transient ferroelectric domains. In other words, the CFE units only have a weak physical pinning effect. Consequently, reversible transitions between the trans and the twisted conformations upon bipolar poling lead to DHLs, as seen in Figure 11B. Second, for P(VDF-TrFE-CTFE),21 stable twisted conformations can be obtained by the inclusion of the even larger CTFE groups with a fairly low dipole moment (μ = 0.64 D). In other words, the CTFE units have a strong physical pinning effect. Upon electric poling, the twisted conformation does not transform into the trans conformation, and thus slim SHLs are achieved (see Figure 11C). From the above discussion of ferroelectricity in PVDF-based polymers, the ferroelectric behaviors of nylon-based polymers can be understood. The intermolecular interaction in nylon crystals is much stronger than that in PVDF crystals, not only because the rigid dipole moment of a non-hydrogen-bonded amide group is 3.7 D30 but also because hydrogen bonding further enhances the intermolecular interaction (e.g., the dipole moment can increase up to ∼5.5 D depending upon the number of hydrogen-bonded amide groups). When the nylon chains adopt the all-trans or nearly all-trans conformation in the hydrogen-bonded sheets, e.g., in the α phase of nylon-11 or the γ phase of nylon-12 [with one twist in the amide −(C O)NH− bond], the sample becomes nonferroelectric. An example for the QA nylon-11 film is shown in Figure 11D, where linear D−E loops are observed. In order to achieve feasible ferroelectric behavior for nylons, the crystals must adopt twisted conformations to reduce the overall hydrogen-bonding strength. An easy way is to quench nylon samples to obtain mesomorphic crystals, which contain multiple twists in the chain conformation. Because of chain twists, hydrogen-bonded sheets are largely destroyed (i.e., disordered hydrogen bonds) and the interchain distance is

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CONCLUSIONS By studying the crystalline structure−ferroelectric property relationships, the principle of nanodomain-induced novel RFE behaviors with slim DHLs and SHLs was successfully demonstrated in nylon-based copolymers and terpolymers. First, PA(11-co-12) copolymers exhibited defect-stabilized mesomorphic phases with twisted chain conformations in the isomorphic crystals, especially when the ADA composition was between 40 and 70 mol %. Consequently, enhanced K

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(5) Baer, E.; Zhu, L. 50th Anniversary perspective: Dielectric phenomena in polymers and multilayered dielectric films. Macromolecules 2017, 50, 2239−2256. (6) Zhang, Q. M.; Bharti, V.; Zhao, X. Giant electrostriction and relaxor ferroelectric behavior in electron-irradiated poly(vinylidene fluoride-trifluoroethylene) copolymer. Science 1998, 280, 2101−2104. (7) Zhang, Q. M.; Huang, C.; Xia, F.; Su, J. Electric EAP. In Electroactive Polymer (EAP) Actuators as Artificial Muscles: Reality, Potential, and Challenges, 2nd ed.; Bar-Cohen, Y., Ed.; SPIE Press: 2004; Chapter 4, pp 89−139. (8) Neese, B.; Chu, B.; Lu, S. G.; Wang, Y.; Furman, E.; Zhang, Q. M. Large electrocaloric effect in ferroelectric polymers near room temperature. Science 2008, 321, 821−823. (9) Moya, X.; Kar-Narayan, S.; Mathur, N. D. Caloric materials near ferroic phase transitions. Nat. Mater. 2014, 13, 439−450. (10) Correia, T.; Zhang, Q. Electrocaloric Materials: New Generation of Coolers; Springer: New York, 2013. (11) Boulesteix, C.; Varnier, F.; Llebaria, A.; Husson, E. Numerical determination of the local ordering of PbMg1/3Nb2/3O3 (PMN) from high resolution electron microscopy images. J. Solid State Chem. 1994, 108, 141−147. (12) Yoshida, M.; Mori, S.; Yamamoto, N.; Uesu, Y.; Kiat, J. M. TEM observation of polar domains in relaxor ferroelectric Pb(Mg1/3Nb2/3)O3. Ferroelectrics 1998, 217, 327−333. (13) Samara, G. A. The relaxational properties of compositionally disordered ABO3 perovskites. J. Phys.: Condens. Matter 2003, 15, R367−R411. (14) Bokov, A. A.; Ye, Z. G. Recent progress in relaxor ferroelectrics with perovskite structure. J. Mater. Sci. 2006, 41, 31−52. (15) Tan, X.; Ma, C.; Frederick, J.; Beckman, S.; Webber, K. G. The antiferroelectric ↔ ferroelectric phase transition in lead-containing and lead-free perovskite ceramics. J. Am. Ceram. Soc. 2011, 94, 4091−4107. (16) Soulestin, T.; Ladmiral, V.; Domingues Dos Santos, F.; Ameduri, B. Vinylidene fluoride- and trifluoroethylene-containing fluorinated electroactive copolymers. How does chemistry impact properties? Prog. Polym. Sci. 2017, 72, 16−60. (17) Tashiro, K. Crystal structure and phase transition of PVDF and related copolymers. In Ferroelectric Polymers: Chemistry, Physics, and Applications; Nalwa, H. S., Ed.; Marcel Dekker: New York, 1995; pp 63−182. (18) Chung, T. C.; Petchsuk, A. Synthesis and properties of ferroelectric fluoroterpolymers with Curie transition at ambient temperature. Macromolecules 2002, 35, 7678−7684. (19) Xia, F.; Cheng, Z.; Xu, H.; Li, H.; Zhang, Q.; Kavarnos, G. J.; Ting, R. Y.; Abdul-Sadek, G.; Belfield, K. D. High electromechanical responses in a poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) terpolymer. Adv. Mater. 2002, 14, 1574−1577. (20) Klein, R. J.; Xia, F.; Zhang, Q. M.; Bauer, F. Influence of composition on relaxor ferroelectric and electromechanical properties of poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene). J. Appl. Phys. 2005, 97, 094105. (21) Yang, L.; Tyburski, B. A.; Domingues Dos Santos, F.; Endoh, M. K.; Koga, T.; Huang, D.; Wang, Y. J.; Zhu, L. Relaxor ferroelectric behavior from strong physical pinning in a poly(vinylidene fluoride-cotrifluoroethylene-co-chlorotrifluoroethylene) random terpolymer. Macromolecules 2014, 47, 8119−8125. (22) Litt, M. H.; Hsu, C. H.; Basu, P. Pyroelectricity and piezoelectricity in nylon-11. J. Appl. Phys. 1977, 48, 2208−2213. (23) Mathur, S. C.; Scheinbeim, J. I.; Newman, B. A. Piezoelectric properties and ferroelectric hysteresis effects in uniaxially stretched nylon-11 films. J. Appl. Phys. 1984, 56, 2419−2425. (24) Lee, J. W.; Takase, Y.; Newman, B. A.; Scheinbeim, J. I. Ferroelectric polarization switching in nylon-11. J. Polym. Sci., Part B: Polym. Phys. 1991, 29, 273−277. (25) Lee, J. W.; Takase, Y.; Newman, B. A.; Scheinbeim, J. I. Effect of annealing on the ferroelectric behavior of nylon-11 and nylon-7. J. Polym. Sci., Part B: Polym. Phys. 1991, 29, 279−286.

ferroelectricity was observed in the nylon copolymers. Nonetheless, the hydrogen-bonding was still not weak enough for the formation of nanodomains and the slim SHL behavior. To further reduce the hydrogen-bonding interaction, 10 mol % NM11 was copolymerized with 30 mol % AUA and 60 mol % ADA to form a terPA-NCH3 terpolymer. Structural analyses from DSC, FTIR, and WAXD showed more defective mesomorphic crystals for the QS terPA-NCH3 terpolymer than those for the QS coPA-6 copolymer, suggesting that the NCH3 groups were probably included in the crystals. Because the NCH3 groups blocked the formation of hydrogen bonds and induce more twists in the chain conformation, hydrogen bonding was further weakened. As a result, slim SHLs with a high dielectric constant of ∼60−70, a signature of the RFE-like behavior, was achieved for the terPA-NCH3 terpolymer at elevated temperatures (≥75 °C). Because nylon terpolymers are much easier to synthesize and cheaper, they are more advantageous than the P(VDF-TrFE)-based terpolymers for potential electrical applications requiring high dielectric constant and low hysteresis losses.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b02243. Thermal data for Q and QA nylon films, 2D WAXD patterns for the QS and QSA nylon films, continuous bipolar loops of the QS nylon films at 185 MV/m, synthesis of NM11, 2D WAXD patterns for terPA-NCH3 and coPA-6 films, subtraction of ac electronic conduction from experimental D−E loops for the QS terPA-NCH3 and coPA-6 films at high temperatures, continuous bipolar D−E loops for the QSA@100 °C terPA-NCH3 film (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Tel +1 216-368-5861 (L.Z.). ORCID

Zhongbo Zhang: 0000-0001-5294-444X Lei Zhu: 0000-0001-6570-9123 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by National Science Foundation (DMR-1402733). Z.Z. acknowledges financial support from China Scholarship Council.

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REFERENCES

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