50th Anniversary Perspective: Dielectric Phenomena in Polymers and

Perspective pubs.acs.org/Macromolecules

50th Anniversary Perspective: Dielectric Phenomena in Polymers and Multilayered Dielectric Films Eric Baer and Lei Zhu* Center for Layered Polymeric Systems (CLiPS) and Department of Macromolecular Science and Engineering, Case Western Reserve University, Cleveland, Ohio 44106-7202, United States ABSTRACT: High dielectric constant and low dielectric loss are desirable electrical properties for next-generation polymer dielectrics that show promise for applications in pulsed power, power electronics, and printable electronics. Unfortunately, the dielectric constant of polymers is often limited to 2−5, much lower than that of inorganic dielectrics, because of the nature of hydrocarbon covalent bonds for electronic and atomic polarizations. It is essential to understand the fundamental physics of different types of polarization and the associated loss mechanisms in polymers. In this Perspective, we discuss the characteristics of each polarization and explain how to enhance the polarization using rational molecular designs without causing significant dielectric losses. Among various approaches for high dielectric constant and low loss polymers, the multilayer film technology is of particular interest because a multilayer film is a unique one-dimensional system with tailored material choices, layer thicknesses, and interfaces. By minimizing the disadvantageous polarizations and enhancing the advantageous polarizations, multilayer films hold promise as advanced dielectrics for future polymer film capacitors.

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trifluoroethylene) [P(VDF-TrFE)] have been used for transducer applications.12 The second type of electromechanical actuation utilizes Maxwell force from the oppositely charged electrodes. With prestraining, dielectric elastomers such as 3M very high bond tape (VHB, an acrylic elastomer) and silicones can exhibit more than 100% deformation.13 The third type of electromechanical actuation utilizes electrostrictive relaxor ferroelectric (RFE) polymers, such as e-beam-irradiated P(VDF-TrFE) and P(VDF-TrFE-X) terpolymers [X being 1,1-fluorochloroethylene (CFE) or chlorotrifluoroethylene (CTFE)].14 Different from soft dielectric elastomers, electrostrictive polymers exhibit a high induced stress with reasonably large deformations (4−10%). In addition to the highpermittivity, electric-field-induced transformation from the mixed TG/TTTG (T is trans and G is gauche) conformation in the RFE phase to the T conformation in the ferroelectric (FE) phase is considered to be a major contribution of electrostrictive actuation.15 The FE property of PVDF and P(VDF-TrFE) can be used for nonvolatile ferroelectric memory applications.16,17 The advantage is the low power consumption because only voltage is used to rewrite the FE pixels. Finally, highly efficient electrocaloric cooling (i.e., reverse pyroelectric effect) is discovered near the Curie transition temperatures (TCs) for P(VDF-TrFE) and RFE e-beamed P(VDF-TrFE) and P(VDF-TrFE-X) polymers.18,19 A temperature change as high as 28 °C is observed for P(VDF-TrFE) at a poling field of 180 MV/m and 50 °C.

ielectric polymers are important insulating materials, which find a broad range of electrical applications. Table 1 summarizes current and potential applications for various dielectric polymers, ranging from linear dielectrics to ferroelectrics. Because of high insulation, cross-linked polyethylene (XLPE) has become the standard in extruded high-voltage direct current (HVDC) cables for underground/undersea power transmission and distribution.1−3 Since 1990, polymer film capacitors have replaced paper-foil capacitors, owing to new metallization technology and the graceful failure mode.4,5 Compared to ceramic and electrolytic capacitors, polymer film capacitors exhibit extremely low losses, stable capacitance, and high-voltage characteristics and are capable of handling high ripple currents. Current state-of-the-art technology uses metallized biaxially oriented polypropylene (BOPP) films. When polarized, highly insulating polymers have been used as electrets.6 For example, electret microphones utilize polarized highly insulating polymer films such as fluorinated ethylene−propylene (FEP) copolymers and polytetrafluoroethylene (PTFE) as the diaphragm.7 Corona-charged nonwoven polypropylene (PP) fiber mats are used for highefficiency air filters, owing to the stable surface charges.8,9 Again, corona charging of cellular PP foams results in ferroelectret materials, which can exhibit a high piezoelectric coefficient, d33, up to 600 pC/N, much higher than that (ca. 200 pC/N) of piezoelectric ceramics such as lead zirconate titanate.10,11 Electromechanical actuation from electroactive polymers is promising for applications such as artificial muscles. The first type of electromechanical actuation is piezoelectric actuation, and poly(vinylidene fluoride) (PVDF) and poly(VDF-co© XXXX American Chemical Society

Received: December 9, 2016 Revised: February 7, 2017

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Macromolecules Table 1. Current and Potential Applications of Polymer Dielectrics applications HVDC cables film capacitors

properties

state-of-the-art polymers

refs

high insulation low loss, stable capacitance, high voltage/current, and graceful failure

XLPEa BOPPa

1−3 4, 5

piezoelectric property stable surface charges

FEP and PTFEa nonwoven PP fiber matsa

7 8, 9

piezoelectric property

PP cellular foams

10, 11

piezoelectric property

PVDF and P(VDF-TrFE)

12

large deformation (>100%) high stress and relatively large deformation (∼5−8%)

3M VHB and silicone e-beamed P(VDF-TrFE) P(VDF-TrFE-CFE), P(VDFTrFE-CTFE) P(VDF-TrFE) and PVDF P(VDF-TrFE), e-beamed P(VDF-TrFE), P(VDF-TrFECFE), P(VDF-TrFE-CTFE)

13 14

electrets microphones high-efficiency air filters cellular ferroelectrets electromechanical actuation piezoelectric transducers dielectric elastomers electrostrictive polymers nonvolatile memory electrocaloric cooling a

rewritable bistable polarization reverse pyroelectric property

16, 17 18, 19

Industrial standard. Those without note are still in the research stage.

Figure 1. (a) Electronic, (b) atomic (or ionic), and (c) total dielectric constant, εr, as a function of the band gap for hydrocarbon-based polymers, computed using density functional theory within the single chain approach. Reproduced with permission from ref 28. Copyright 2014 Elsevier.

should be no thicker than the normal thickness of current BOPP films, i.e., 2.5−3.0 μm. Although ultrahigh capacitance density has been achieved for molecular gate dielectrics such as self-assembled monolayers, technological challenges such as high pinhole density in self-assembled monolayers demand high-κ and low-loss spin-coatable gate dielectrics with ultralow pinhole density.26 To realize this goal, the following property parameters need to be achieved for next-generation gate dielectric polymers: εr > 15, capacitance density >15 nF/cm2, leakage current 300 MV/m, and high operation frequency up to 5 GHz. In this Perspective, we discuss various dielectric phenomena in polymers and their applications for the next-generation film capacitors and gate dielectrics.

Because of the rapid advances in electrical and power applications, high dielectric constant (or permittivity, εr or κ) and low loss polymer dielectrics have become increasingly attractive because they can enable high energy density and high power density at the same time, which lead to device miniaturization for better performance and higher efficiency. However, a knowledge gap exists between electrical engineering at the device level and fundamental polymer science at the material level. Currently, no single-component dielectric polymers exhibit a high κ (>6)/high energy density, high temperature tolerance (>150 °C), and low dielectric loss (e.g., dissipation factor, tan δ < 0.002) simultaneously. There is an urgent need to design and develop novel high-κ and low-loss dielectric polymers for the above-mentioned electrical applications. Because of limitation of space, this Perspective will only focus on the applications of next-generation film capacitors20−22 and gate dielectrics in field effect transistors (FETs).23−25 Although much has been reviewed for both applications in the past, challenges still exist. For example, nextgeneration film dielectrics need to simultaneously achieve high εr (>10)/high energy density (10−20 J/cm3), ultralow dielectric loss (tan δ < 0.005), high breakdown strength (Eb), graceful failure, upper temperature limit of at least 125 °C, and upper operating frequency of MHz. More importantly, the price of packaged capacitors needs to be less than 1/2 of current BOPP capacitors with the same capacitance. In other words, the new high-κ, high-Eb, and low-loss polymer films

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LIMITATIONS OF CURRENT STATE-OF-THE-ART DIELECTRIC POLYMERS There are four types of polarization in dielectric polymers: electronic, atomic (or vibrational), orientational (or dipolar), and space charge (electrons/holes and ions) polarizations.27 Electronic and atomic polarizations belong to the resonance regime, and they take place in the optical (1014−1016 Hz) and infrared frequencies (1011−1014 Hz), respectively. Because there is no loss in the power and radio frequencies ( 1000) need to be discovered. According to the space-charge-limited current theory,27 a higher εr will result in a higher bulk conductivity. There should be a balance between high εr and low bulk conductivity. Third, it is known that dielectric constant of FE ceramic particles (e.g., BaTiO3) significantly decreases when their size decreases to the nanometer scale.50−53 This phenomenon is attributed to the disappearance of 90° domains and domain walls when the grain size is below 1 μm.54,55 When the particle size decreases to below 20−30 nm, the FE phases are destabilized and the paraelectric (PE) cubic phase exhibits a low εr around 100. In recent studies, uniform BaTiO3 nanoparticles have been synthesized via the nanoreactor approach by utilizing star- and comb-like poly(acrylic acid)-b-polystyrene (PAA−PS) block copolymers.56,57 Because of the inclusion of PAA chains in the inorganic BaTiO3 nanocrystals, the FE phases are stabilized, and therefore high εr values of 300−600 are reported for 10−27 nm hybrid BaTiO3 nanoparticles.58,59 It will be interesting to see how these high-κ BaTiO3 nanoparticles perform in polymer nanodielectrics in the future. In addition to ceramic nanoparticles, conducting metallic and carbonaceous (carbon nanotubes and graphenes) nanoparticles are also used as fillers in polymer nanodielectrics. Ideally, the permittivity of conductive nanoparticles can be considered to be infinity. Because these nanoparticles are fairly conductive, the internal conduction loss can be neglected; therefore, conductive nanoparticles seem to be better than high-κ ceramic nanoparticles for polymer nanodielectrics. A number of publications report that high dielectric constants can be achieved for polymer/conducting particle nanodielectrics when the conducting filler content is just below the percolation threshold. However, there is a serious issue for polymer/ conducting particle nanodielectrics to be used under high electric fields. It is known that the local electric field in conducting particles is zero, no matter how high the applied electric field is, because the polarized interfacial charges in conducting particles can completely compensate the external field. Similar to the situation in polymer/ceramic particle nanodielectrics, these polarized interfacial charges can also significantly increase the local field in the polymer matrix between the aligned particles in the field direction (e.g., see

breakdown strength further reduces to 162 MV/m for the PP/ BaTiO3 80/20 film. In addition, the Weibull slope, which reflects the distribution of breakdown strength, decreases from 11.4 for the hot-pressed PP film to 5.7 for the PP/BaTiO3@ POSS 80/20 and 4 for the PP/BaTiO3 80/20 films, suggesting decreased film quality with increasing the BaTiO3 content in the PP matrix. The mechanism of decreased Eb for the PP/BaTiO3@POSS nanocomposites is understood using a coarse-grained molecular dynamics computer simulation.46 Figure 3B shows threedimensional (3D) local field enhancement (EL/E0, where EL and E0 are local and applied electric fields) in a random nanocomposite with a filler content of 20 vol %.45 The particles are 70 nm and have a dielectric constant of 125, as taken from experimental results. The matrix has a dielectric constant of 2.25. Because of the large permittivity contrast (125 for the particles vs 2.25 for the matrix), the local field is concentrated near the particles, especially around the two poles along the applied field direction. When more than three particles form a chain in the external field direction, the local field is greatly enhanced due to the strong dipole−dipole coupling effect.46 This is clearly seen in Figure 3B, and the maximum enhancement can be as high as EL/E0 ∼ 18−20. These locations with a concentrated local field can form hot spots, which are prone to dielectric breakdown in the sample. In the above simulation, space charges are not taken into account. In real nanocomposites, space charges cannot be ignored (particularly for those with a large contrast in conductivity, σ), and they play a significant role in the dielectric loss mechanism for polymer nanodielectrics. The effect of space charge polarization-induced dielectric loss mechanism is studied using electric displacement−electric field (D−E) loop tests. Figure 3C shows 10 Hz bipolar D−E loops (dashed lines) for the PP/BaTiO3@POSS 80/20 film at room temperature.45 When the poling field is beyond 30 MV/ m, a significant loop loss (or broadening) is observed. The figure also shows four continuous unipolar D−E loops (20 Hz, solid lines) for the nanocomposite film at room temperature. The maximum poling field is 150 MV/m. Intriguingly, the first unipolar loop exhibits a high loop loss, similar to that in the bipolar D−E loop, whereas the rest consecutive unipolar loops are fairly narrow with a much lower loop loss. These experimental results can be explained using the model of space charge polarization in Figure 3D.45 Because of the large conductivity contrast, 10−10 S/m for BaTiO3 and 10−17 S/m for PP, the concentration of thermally activated free electrons must be higher in the BaTiO3 particles than in the PP matrix. Note that the space charges are thermally activated electrons in this case, rather than impurity ions. Under electrical poling, these free electrons can be polarized to one side of the nanoparticle (to account for the neutrality of the particle, imaginary positive charges or holes are shown on the opposite site of the particle). When the external field alternates, these free electrons will be polarized back and forth, forming two conduction pathways: (I) boundary layer conduction and (II) bulk conduction (see red arrows). Because these free electrons are confined within the BaTiO3 particle, they form internal AC conduction. The large loop loss in the bipolar D−E loops in Figure 3C is a result of these internal AC conductions. Under the unipolar poling, the free electrons are polarized to one side of the particle and will not have a chance to relax back and switch polarity. Therefore, only the first unipolar loop shows a high loop loss and rest loops show much lower loss. E

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Figure 4. (A) Schematic of the frequency dependence of real (εr′) and imaginary (εr″) relative permittivity for pure water at room temperature and a polar glassy polymer. (B) Schematic of intermediate dipolar states between (a) the linear dielectric and (e) the normal ferroelectric behavior, including (b) dipolar glass, (c) paraelectric, and (d) relaxor ferroelectric behavior. Reproduced by permission from ref 20.

temperature.71 Present PE polymers include both polar polymer melts and crystals. The only crystalline type of PE polymers reported so far includes P(VDF-TrFE) and P(VDFco-tetrafluoroethylene) [P(VDF-TFE)] above their TC.72 Because of the lack of FE domains, PE polymers usually exhibit a narrow D−E loop (Figure 4B-d, but having a relatively low apparent dielectric constant of 30−40). Quenching these high-temperature PE polymers into the DG state has not been successful. Instead, a cooled FE phase is obtained upon quenching, where the chains in the FE crystal contains numerous gauche defects (or kinks).72 Polar polymer melts are considered as PE, and they include amorphous PVDF, polyamides (nylon), polyesters, and polyacrylates above their Tgs. By quenching below the Tg, these glassy polymers can be considered as DG polymers, and dipole rotation is realized via sub-Tg transitions73 although the segmental chain motion is still largely frozen. Note that the DG polymers with sub-Tg transitions are different from polar molecule-doped glassy polymers,74,75 where the doped dipolar molecules are usually too large to freely rotate below the Tg of the matrix polymer. At high doping concentrations, the dipolar molecules tend to macrophase separate from the polymer matrix. However, these are not issues for DG polymers because the dipolar side groups with small sizes are covalently attached to the polymer chains. The interactions among these dipolar groups are relatively weak due to the absence of translational motion in the glassy state and thus the chance to form FE domains. It is expected that their D−E loops should be narrow with relatively high εrs (Figure 4B-b). Because of the frozen chain dynamics below the Tg, DG polymers can exhibit low leakage and thus a low dissipation factor. DG polymers with relatively high εrs have already been studied, and some are even commercialized. Most of these DG polymers contain cyanoethyl (−CH2CH2CN) or cyanomethyl (−CH2CN) side groups. When −CH2CN groups are attached as the side chains in a bisphenol A polycarbonate (i.e., CN-PC), the dielectric constant at 1 kHz increased to 4.0 for CN-PC as compared to that of 2.9 for the neat PC.76 The dielectric loss was reasonably low, e.g., tan δ ∼ 0.005 at 130 °C and 1 kHz. However, the dielectric constant was low due to the low density of −CH2CN dipoles. Cyanoethylated poly(vinyl alcohol) (CNPVA) exhibited a dielectric constant of ca. 10 below its Tg of 25 °C.77 Cyanoethylated poly(2,3-dihydroxylpropyl methacrylate) (CN-PDPMA) exhibited a relatively high εr of ∼8 between the β (relaxation of the cyanoethyl dipoles at −60 °C) and the α (Tg at 25 °C) transitions at 500 Hz.78 However, the window

Figure 3B). A unique property for conducting materials under a strong electric field is the field electron emission.60 Note that the Coulomb-blockade effect only happens either at ultralow temperatures or for metallic particles smaller than 2 nm.61 It is the field electron emission that substantially decrease the electrical percolation threshold (ca. 12 vol %) than the physical percolation threshold (30−35 vol %) in polymer/conducting particle nanodielectrics observed experimentally.62 As a result of the local field enhancement, the dielectric breakdown strength substantially decreases upon increasing the conducting particle content. Although polymer/conducting particle nanodielectrics are not suitable for high field applications, they are good candidates for field-dependent electrical switches63,64 and piezoresistive-like sensors.65 For example, quantum tunneling composites (QTCs) utilize spiky metallic particles in a dielectric elastomer matrix to enhance the field electron emission, leading to a piezoresistive-like effect.65 Upon a small mechanical deformation, a large change in electrical resistivity is observed.

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ORIENTATIONAL POLARIZATION TO ENHANCE DIELECTRIC CONSTANT FOR ELECTRIC ENERGY STORAGE An important polarization mechanism in the relaxation regime is orientational polarization. The idea is to utilize dipolar groups to enhance the permittivity of polymers. Small molecules, e.g., N-methylformamide (εr ∼ 180)66 and water (εr ∼ 80),67 can exhibit high dielectric constants with reasonably low dielectric loss at room temperature. For example, the dipole relaxation peak for pure water at 25 °C is found at 20 GHz, and the dissipation factor at 100 MHz is as low as 0.005 (Figure 4A). However, polar polymers [e.g., cyano (CN)-containing polyimides68] exhibit dielectric constants of only 3−5, and the broad dipole relaxation peak directly covers the frequency range of interest (Figure 4A). Learning from H2O,67,69 it is highly desirable to enhance the dielectric constant and speed up the dipolar relaxation toward the gigahertz range for polar polymers. Recently, we proposed various intermediate dipolar states between a linear dielectric state and a nonlinear FE state (Figure 4B).20 They include dipolar glass (DG), PE, and RFE polymers. Dipolar Glass Polymers. Reminiscent of spin glasses in magnetic materials,70 a dipolar glass contains isolated and weakly interacting electric dipoles (Figure 4B-b). It can be considered as a quenched or frozen PE upon decreasing the F

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Figure 5. (A) Schematic representation of the multilayer coextrusion process. (B) AFM phase images of PC/PVDF 50/50 (vol/vol) multilayer films: (a) 2-layer, (b) 8-layer, (c) 32-layer, and (d) 256-layer. The total film thicknesses are around 12 μm. The individual layer thicknesses are 6000, 1500, 380, and 50 nm, respectively. Reproduced with permission from ref 105.

between the β and α transitions was narrow, only about 85 °C. Meanwhile, the Tgs for CN-PVA and CN-PDPMA are too low for high-temperature applications. The low Tg of cyano-ethylated polymers can be attributed to the relatively long side chains. To overcome this problem, ShinEtsu Chemical commercialized cyanoethylated pullulan (CEP, ∼90% functionality). The original application was for a high-κ coating for the phosphor grains in so-called thick-film electroluminescent lamps.79 Because of the rigid cellulose backbone structure, the Tg reached ∼110 °C,80 and the dielectric constant at room temperature was 13−18 at frequencies below 104 Hz.81 From the frequency-scan result, the dipole relaxation peak is just above 1 MHz at room temperature. We recently reported methylsulfonyl-containing DG polymers because the sulfone group has an even higher dipole moment (∼4.25 D) than that of the cyano group (3.9 D). For a new DG polymer, poly[2-(methylsulfonyl)ethyl methacrylate] (PMSEMA),82 both BDS and D−E loop results show a high εr of 11−12 at room temperature with a reasonably low dissipation factor of 0.02 at 10 Hz. At room temperature, the methylsulfonyl dipole relaxation occurs around 1 MHz. Compared with the theoretical polarization from the Langevin model,27 about 40% of the sulfone dipoles can rotate to give a high εr of 12. In the future, it is desirable to further increase the percentage of dipole rotation to achieve even higher εr values (e.g., >20) for DG polymers. High Dielectric Constant and Low Loss RFE Polymers. In general, PE polymer melts are not suitable for dielectric insulation because of high charge carrier mobilities and thus high DC conduction. On the contrary, FE polymers are also not suitable for electric energy storage because of their large FE hysteresis (Figure 4B-e), which is attributed to the large FE domains. For electric energy storage applications, it is desirable to minimize the hysteresis of FE polymers and achieve narrow single and double hysteresis loops (i.e., SHL and DHL), as shown in Figures 4B-c/d. Based on the state-of-the-art understanding of RFE ceramics,83,84 an effective strategy is to divide the large FE domains in normal ferroelectrics into nanodomains. For RFE ceramics such as Pb(Mg1/3Nb2/3)O3 (PMN), the domain size is only 2−3 nm, as revealed by high-

resolution transmission electron microscopy (TEM) study.85,86 From recent reports,20,87 repeat unit isomorphic (or defectmodified) crystals are effective to achieve this goal for RFE polymers. Basically, isomorphic P(VDF-TrFE) copolymers (ebeam irradiated)88,89 and P(VDF-TrFE-X) terpolymers (X is CFE or CTFE)90,91 have successfully achieved the SHL and DHL behavior. The large termonomers such as CFE and CTFE serve as several important roles. First, they are included in the isomorphic crystals and effectively divide the large FE domains into nanodomains by introducing gauche defects in the long trans sequence. Second, due to their large sizes, interchain distance is expanded as a result of crystal isomorphism (i.e., inclusion of the termonomer units in the crystals), which facilitates easier dipole switching under an alternating field. Third, they form effective pinning points for the dipoles between neighboring termonomer units along the chain, and this greatly helps to restore the poled dipoles to the original random state. As a result of these effects, narrow DHL and SHL can be achieved. Note that the crystal isomorphism mechanism for the e-beam cross-linked P(VDF-TrFE) copolymers is somewhat different from that for the P(VDF-TrFE-X) terpolymers. Because of the cross-linked nature, the detailed mechanism is still difficult to know. However, solid-state 19F NMR studies showed the formation of −CF3 side groups in the e-beam cross-linked P(VDF-TrFE).92,93 It is speculated that the TrFE units inside crystals can cause cross-linking, and some of them transform into large −CF3 containing units to pin the VDF/TrFE dipoles. From the processing point of view, terpolymers have more advantages than the e-beamed copolymers because the tedious high-energy e-beam irradiation can be avoided. Eventually, the apparent dielectric constants for e-beam-irradiated P(VDF-TrFE) 50/50, P(VDF-TrFE-CFE), and P(VDF-TrFE-CTFE) are about 30, 55, and 70, respectively, as determined by high-field D−E loop studies at 10 Hz.20 Using the concept of crystal isomorphism learned from P(VDF-TrFE-X) terpolymers, RFE n-nylon copolymers and terpolymers are currently being developed. Contrary to the conventional understanding that only odd-numbered nylons can exhibit ferroelectricity,94 we recently reported that evennumbered n-nylons could also exhibit ferroelectricity.95 The G

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Figure 6. (A) Real (εr′) and (B) imaginary (εr″) parts of the relative permittivity as a function of temperature for PC/PVDF 50/50 (vol/vol) 32layer films under different frequencies. Reproduced with permission from ref 20.

mechanism is attributed to the high electric field-induced ferroelectric domains in mesomorphic crystals quenched from the melt, no matter even- or odd-numbered nylons. By random copolymerizing different n-nylon comonomers, such as those for nylon-11 and nylon-12, structural defects with random hydrogen bonding are built-in inside the isomorphic crystals. Narrow SHLs with an apparent dielectric constant of ca. 60 can be obtained at elevated temperatures (>75 °C), which are similar to those of the RFE P(VDF-TrFE-CTFE) terpolymers.20 From this study, the concept of isomorphic crystals with nanodomains should be universal to transform normal FE polymers into RFE polymers with high εr and a reasonably low loop loss. Although high dielectric constants, which are close to that of water, have been achieved for P(VDF-TrFE-X) terpolymers, there are several drawbacks. First, the high εr is a result of nanodomain switching, which is not suitable for frequencies higher than 1 kHz. Our preliminary high-voltage BDS results show that when the poling frequency is above 1 kHz, nanodomains start to lag behind the fast switching electric field. This is similar to the case of PVDF.96 Consequently, the εr will significantly decrease. Second, polar polymers are easily contaminated with impurity ions, and even a sub-ppm concentration can cause high nonlinear dielectric losses, especially at elevated temperatures and low frequencies.97 Because of the ionic and electronic conductions, the dielectric loss is still fairly high. For example, the dissipation factor is around 0.01−0.02, and the hysteresis loop loss at high fields can be greater than 0.2.98 For most capacitor-type applications, this is not acceptable since the heat generation during the charge− discharge cycles will quickly melt the polymers. Third, PVDFbased polar polymers can undergo electrochemical degradation and emit HF during a long-term operation at high voltages.99,100 The emitted HF will corrode metal electrodes. For practical application, research needs to be carried out to mitigate these issues. In the following, we introduce multilayer polymer films, which is a viable approach to further decrease dielectric losses for high-κ polymers.

effect plays an important role in tailoring material’s properties.102,103 For dielectric applications,104 these multilayer films are usually comprised of alternating layers of a high-κ polymer and a linear dielectric polymer with low dielectric loss and high breakdown strength. The high-κ polymer can be PVDF (or its random copolymers) and nylons, which exhibit an εr greater than 10 above the Tg. The linear dielectric polymer can be PC, polysulfone (PSF), and poly(ethylene terephthalate) (PET), which possess high breakdown strengths and can adhere reasonably well with the high-κ polymer. The dielectric constants of these linear dielectric polymers are usually low, ranging from 2.9 to 3.6. Example atomic force microscopy (AFM) phase images for the PC/PVDF 50/50 multilayer films are shown in Figure 5B.105 In these films, all types of polarization take place, which can be divided into advantageous and disadvantage groups. The advantageous group includes deformational (i.e., electronic and atomic) polarization from both components, orientational polarization in the amorphous phase of the high-κ polymer, and interfacial polarization. The disadvantageous group includes hopping polarization of space charges (i.e., impurity ions and electrons) and ferroelectric switching of the crystalline phase of PVDF.106 For practical applications, viable strategies are needed to maximize the advantageous polarizations and minimize the disadvantageous polarizations. Orientational Polarization in the High-κ Polymer. Because of good commercial availability, semicrystalline PVDF and P(VDF-co-hexafluoropropylene) [P(VDF-HFP)] copolymers are often used as the high-κ polymer in multilayer films. In these polymers, the molten amorphous phase is PE and thus has a relatively high dielectric constant. The crystalline phase is FE, and ferroelectricity should be avoided for electric energy storage. Figures 6A,B show the temperature-scan low-field BDS results of relative permittivities, εr′ and εr″, for the PC/PVDF 50/50 (vol/vol) 32-layer (32L) film (Figure 5B-c) at various frequencies.20 From the εr′ and εr″ results at 1 Hz, various relaxation processes can be identified: γPC at −105 °C,73 αa,PVDF (i.e., Tg,PVDF) at −42 °C,99 αc,PVDF at 10 °C,99 ion relaxation in PVDF at 108 °C,99 and onset of αPC (i.e., Tg,PC) around 140 °C.73 These transitions gradually shift to higher temperatures upon increasing the frequency. The enhanced εr′ from 3.0 at −100 °C to 4.7 at 10 °C for the multilayer films at 103 Hz is primarily a result of the orientational polarization from the amorphous PVDF dipoles (i.e., αa,PVDF; see Figure 6A). Because the αa,PVDF depends on frequency, an appropriate combination of temperature and frequency is needed to avoid the

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MULTILAYER POLYMER FILMS FOR DIELECTRIC APPLICATIONS Using the state-of-the-art multilayer coextrusion technology at Case Western Reserve University (CWRU, see Figure 5A), multilayer polymer films have been fabricated with the individual layer thickness reaching as low as ca. 10 nm.101 In these multilayer films with thin layers, the nanoconfinement H

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Figure 7. Polarized FTIR spectra for uniaxially stretched (draw ratio of ∼300%) (A) PVDF and (B) PC/PVDF 50/50 32-layer films under different DC electric fields. (C, D) FTIR absorption as a function of the applied DC bias field for the β and the α/δ absorption peaks at 509 and 488 cm−1, respectively. Reproduced with permission from ref 105.

detrimental for electric energy storage. Therefore, it is highly desired to avoid FE switching in PVDF crystals. Multilayer films are effective in achieving this goal. According to the dielectric theory,27,110 electric field distribution in multicomponent systems is nonuniform. Namely, the high-κ domains bear a low local field and the low-κ domains bear a high local field. At room temperature, the dielectric constant of PVDF is ca. 12 whereas that for PC is around 2.9. Using a bilayer dielectric model,105 the local field in PVDF is only 1/4 of that in PC. It is because of the low local field in the PVDF layers that FE switching in the crystalline phase can be largely avoided. Meanwhile, the high breakdown strength of PC will make the PC layers survive the high local field. This hypothesis is testified by polarized FTIR experiments on uniaxially stretched (draw ratio of ∼300%) PVDF and PC/PVDF 50/50 32L films (ca. 10 μm).105 As shown in Figure 6A, both α and β phases are present in the stretched PVDF film. The intensities of the absorption bands at 509 and 488 cm−1 show clear electric field dependence. The 509 and 488 cm−1 peaks are assigned to the CF2 bending motions in the β and α/δ crystals, respectively. Upon the application of a bipolar electric field, both peak intensities follow a butterfly pattern (Figures 6C,D), suggesting the FE switching of FE domains in the stretched PVDF film. However, the stretched PC/PVDF 50/50 32-layer film only show the α phase (Figure 6B) with no sign of FE switching in PVDF up to 200 MV/m (Figure 6D). From this study, it is clear that multilayer films with alternating high- and low-κ layers are effective in preventing the FE switching in PVDF crystals. We should note that FE switching is still possible when the applied electric field is high enough, e.g., ≥600 MV/m (data not shown).

corresponding dielectric loss from the glass transition. For example, the temperature needs to be above 10 °C at 103 Hz, whereas it needs to be above 50 °C when the frequency is 105 Hz in order to achieve the minimum dielectric loss (Figure 6B). However, right above the αa,PVDF relaxation peak in Figure 6B, the αc,PVDF relaxation is observed, which originates from the wagging motion of CF2 dipoles along the c-axes of α PVDF crystals.107 The appearance of the αc,PVDF indicates certain unoriented α crystals in the PVDF layers (ca. 380 nm thick). Although the αc,PVDF further increases the εr′ to 5.1 at 50 °C (1 Hz), there is a significant loss associated with it. To minimize the dielectric loss from the αc,PVDF relaxation, several strategies can be pursued. First, confined melt-recrystallization in the 380 nm PVDF nanolayers can produce the edge-on crystals.105 Once the α crystals orient perpendicular to the applied electric field, the αc,PVDF loss will be reduced. Second, stretching is an alternative way to align the α crystals parallel to the film, which also reduces the αc,PVDF loss.108 Finally, converting the α crystals to the β crystals should decrease the αc,PVDF loss. Mechanical stretching, high-field electric poling, or blending with PMMA can induce the β phase for PVDF;109 however, the phase conversion usually cannot be 100%. To avoid the α phase, P(VDF-TrFE) with a high VDF content (e.g., 75 mol %) and thus a high TC can be a good choice because the isomorphic crystals adopt a zigzag conformation with a minimum amount of gauche defects below the TC.72 Above 100 °C, the ionic conduction loss becomes a major loss mechanism, and we will discuss strategies to minimize ionic conduction loss later. Orientational polarization in the crystalline phase of PVDF and its copolymers can lead to ferroelectricity, which is I

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Figure 8. (A) Dissipation factor (tan δ) as a function of frequency for the PC/PVDF 50/50 2L, 8L, 32L, 256L, and PVDF and PC control films at 75 °C. All films have a total thickness of ca. 12 μm. (B) Schematic of the nanoconfinement effect on the ion-hopping polarization in PVDF layers having different thicknesses. Reproduced with permission from ref 105. (C) Schematic of a monolithic slab (thickness a) containing impurity ions, sandwiched between two metal electrodes. The matrix has a dielectric constant of 10. The applied voltage V(t) = V0eiωt. (D) Simulated D−E loops for the monolithic slab having various thicknesses at 100 °C. The poling frequency is 10 Hz.

Ion-Hopping Polarization in the High-κ Polymer. For polar polymers, ionic mobility is usually higher compared to that in nonpolar polymers. By fitting the BDS frequency-scan results of εr′ and εr″ using the blocking electrode theory,111 the impurity ion concentration (n) and diffusion coefficient (D) can be obtained. For example, for PVDF at 100 °C, the impurity ion concentration is about n = 2 × 1021 /m3 and the diffusion coefficient is about D = 3 × 10−13 m2/s by fitting with the Sawada equations (the theory is for low ion concentrations, and the local field is not modified by the accumulated ions at the dielectric/electrode boundaries).112 Assuming the impurity ions are Na+ and Cl−, the ion concentration is just below 2 × 10−7 g/mL (note that ionic impurities in PVDF are originated from the nonionic surfactants used in suspension polymerization113). Because the Na+ size (2.32 Å) is smaller than that (3.34 Å) of Cl−, it is reasonable to consider that only Na+ cations move under the applied field, and the Cl− anions are more or less immobile. Effects of impurity ions on dielectric properties of the multilayer films are studied by BDS.105 Figure 8A shows the frequency-scan dissipation factor (tan δ) for the PC/PVDF 50/50 2L, 8L, 32L, and 256L films (see Figure 5B) at 75 °C, together with those for the PVDF and PC control films (note that all films are about 12 μm thick). In addition to the αc,PVDF relaxation peak around 700 Hz, there is a lowfrequency loss peak around 0.05 Hz, which can be attributed to the impurity ion-hopping polarization. Intriguingly, upon decreasing the PVDF layer thickness by increasing the number of layers, the loss from the ion-hopping gradually decreases. Eventually, the loss peak for the 256L film completely disappears. This can be explained by the nanoconfinement effect on ion hopping, as shown in Figure 8B. For thick PVDF layers, the ion hopping is more significant than that for thin PVDF layers. The nanoconfinement effect can be understood by theoretical calculation of the D−E loops45,97 for a monolithic slab containing mobile ions at 100 °C, as shown in Figure 8C.

In this case, we assume that single valence cations are mobile and the anions are immobile. The dielectric constant of the matrix is 10, and the impurity ion concentration and diffusion coefficient are taken from the experimental fitting results mentioned above. The total current density includes contributions from both the displacement current and the ionic conduction current. The calculated D−E loops are shown in Figure 8D, and the hysteresis loop loss is critically dependent upon the layer thickness. For layer thickness (a) above 200 nm, broad loops are observed, whereas narrow loops are seen when the layer thickness is below 100 nm. Especially, the hysteresis loop loss is negligible when a is 50 nm. This theoretical result fits well with the experimental observation. Namely, the ionic conduction loss at 0.05 Hz for the PC/PVDF 50/50 256L film (i.e., PVDF layer is ca. 50 nm) is negligibly small in Figure 8A. From this study, one strategy to reduce ionic conduction loss is to decrease the PVDF layer thickness. Interfacial Polarization in Multilayer Films. For phaseseparated multicomponent systems having large contrasts in dielectric constant and conductivity, interfacial polarization from space charges is ubiquitous. Because of the extended hopping polarization, space charges (impurity ions and electrons/holes) in the more conductive domains will accumulate at the interfaces, blocked by the more insulating domains. As a result, these interfacial space charges will modify the effective local fields in the high-κ and high-conductivity domains. In this sense, the interfacial polarization can reduce electronic/ionic conductions and increase the breakdown strength for polymers.114 Dielectric breakdown strength is studied for the PC/PSF 32L and the PC/PVDF 32L films using the needle-plane method at room temperature. The needle tip diameter is ca. 40 μm, and each data point is the average of at least 10 samples. Although the needle-plane geometry deviates from the parallel plane geometry and thus the apparent electric field appears not very accurate, this method gives a qualitative trend for the J

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Figure 9. Dielectric breakdown strength as a function of PC composition for (A) the PC/PSF 32L and (B) the PC/P(VDF-HFP) 32L films measured by the needle-plane method at room temperature. Each data point is an average of at least 10 samples. Reproduced with permission from ref 115. Copyright 2009 IOPscience. (C) Weibull analysis of dielectric breakdown strength for the PSF/PVDF 30/70 32L and 256L films measured by the plane−plane method (electrode area =1.77 cm2) at room temperature. (D) Schematic of the interfacial polarization from both impurity ions and free electrons for the PSF/PVDF multilayer films with either thick or thin PSF layers. Reproduced with permission from ref 116. Copyright 2014 Elsevier.

breakdown strength of polymers. For the PC/PSF 32L films, there are nearly no contrasts in dielectric constant and conductivity for PC (εr,PC = 2.9 and σPC = 10−17 S/m) and PSF (εr,PSF = 3.1 and σPSF = 10−17 S/m). Therefore, no effect of interfacial polarization on dielectric breakdown strength is observed for the PC/PSF 32L films (Figure 9A). For the PC/ P(VDF-HFP) 32L films at room temperature, there are large contrasts in dielectric constant and conductivity for PC and P(VDF-HFP) (εr,P(VDF‑HFP) = 10−12 and σP(VDF‑HFP) = 10−14 S/ m). Enhanced dielectric breakdown strengths are observed for the PC/P(VDF-HFP) multilayer films compared to the individual components and the linear average values (Figure 9B).115 For comparison, melt-extruded PC/P(VDF-HFP) blend films with a total thickness around 12 μm are also studied. All the blend films exhibit lower dielectric breakdown strengths than the linear average values. This result indicates that both interface geometry and orientation are important for the dielectric insulation of multicomponent systems. Only when the flat interfaces align perpendicular to the applied field direction, such as in multilayer films, the dielectric insulation will be enhanced. In immiscible polymer blends, curved interfaces are randomly oriented, and those more or less parallel to the applied field serve as conductive pathways. This is also the reason why polymer nanodielectrics with aggregated spherical and rodlike particles exhibit low dielectric breakdown strengths45 because connected interfaces serve as effective conductive pathways. Only for polymer nanodielectrics having oriented nanoplatelets (similar to the multilayer film geometry), enhanced dielectric breakdown strength is observed.49

In multilayer films, the linear dielectric polymer serves as blocking electrodes, and their insulation is important for the ultimate dielectric properties. Figure 9C shows the dielectric breakdown strengths for the PSF/PVDF 30/70 32L and 256L films measured by the plane−plane method at room temperature.116 Because the electrode area of 1.77 cm2 is much larger than that of the needle-plane method, the breakdown values appear to be lower due to a higher content of defects in a larger sample area. Although both films have the same composition and a similar total thickness (12 μm), the 256L film shows a consistently lower breakdown strength than the 32L film. This can be attributed to the poorer insulating property of the thinner PSF layers (ca. 28 nm) in the 256L film than the thick PSF layers (ca. 220 nm) in the 32L film. At room temperature, the ionic conductivity in PVDF is negligibly low,99,116 and thus the interfacial polarization can be primarily attributed to the free electrons (Figure 9D), which are most likely thermally activated in PVDF due to its high dielectric constant and the high content of impurities. Because PVDF is an insulating material, there should not be mobile holes. To account for the neutrality of the PVDF layer, positive charges are drawn in Figure 9D to represent imaginary holes. For thick PSF layers, the insulation is good and free electrons in PVDF will be accumulated at the PVDF/PSF interfaces upon electric poling. These interfacial electrons can modify the local electric field and prevent the injected electrons from the cathode from passing through the sample. On the other hand, thin PSF layers have a poor insulation, and certain interfacial electrons can tunnel/penetrate through them. As a result, the overall insulation for the 256L film becomes poorer than that for the K

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Figure 10. (A) Schematic of biaxial stretching of PET/PMMA/P(VDF-HFP) multilayer films using a Bruckner Karo IV lab-scale stretcher. (B) Maximum discharged energy density (Ue) at breakdown as a function of the PMMA content for as-extruded and biaxially oriented PET/PMMA/ P(VDF-HFP) 65L films (4.5 × 4.5). (C) D−E loops for as-extruded and biaxially oriented PET/PMMA/P(VDF-HFP) 65L films with 0% and 8% PMMA, respectively. Reproduced with permission from ref 108.

32L film due to fewer interfacial electrons. Not only the layer thickness affects the insulation, the Tg of the linear dielectric polymer also plays an important role; the higher the Tg, the better the insulation for multilayer films. For example, PSF has a higher Tg (185 °C) than PC (145 °C). Therefore, the PSF/ PVDF multilayer films should be more insulating than PC/ PVDF multilayer films. In addition to the dielectric breakdown strength study, lifetimes of the PC/P(VDF-HFP) 32L and 256L films with various compositions are also investigated using Weibull analysis when the samples are held under a constant electric field.117 It is observed that the lifetimes of the PC/P(VDFHFP) 32L and 256L films are higher than the linear average values. Especially, the lifetimes for the 70/30 films are nearly the same as that of the PC control film. Furthermore, the 32L films exhibit slightly higher lifetimes than those of the 256L films, which again can be attributed to the higher interfacial polarization in the 32L films than in the 256L films. For 10 μm thick PC/P(VDF-HFP) 50/50 multilayer films, the optimal lifetime is found for the 65L films (i.e., both sides are PC).118 Currently, thermally stimulated depolarization current (TSDC) experiments are underway to provide direct evidence for the electronic interfacial polarization in PVDF-based multilayer films. Besides interfacial electrons, impurity ions can also form interfacial polarization upon electric poling (Figure 9D), especially when the ion mobility is high at elevated temperatures. Similar ionic interfacial polarization is also reported for monolithic biaxially oriented PVDF films, as detected by TSDC.99 Because the concentration of impurity ions is much higher than that of free electrons in PVDF, the ionic interfacial polarization can significantly modify (or reduce) the local electric field. A direct result of the ionic interfacial polarization (i.e., EDL) is the decrease of the electric field in the bulk of the PVDF layers, and energy storage from the orientational polarization of amorphous PVDF dipoles will be lost. It is desirable to diminish this effect by driving impurity ions into the PSF or PC layers. It is known that small ions are still mobile even when the temperature is below the Tg of a polymer.119 By electric poling at a temperature slightly below the Tg of PSF or PC, it is possible to drive the impurity ions in PVDF into the PSF or PC layers. After locking the impurity ions in the glassy PSF or PC layers, the dielectric loss from the ion-hopping polarization will be reduced. Currently, experiments are underway to achieve this goal.

Note that low loss dielectric applications demand an ultralow level of impurities or contaminants. For example, capacitor grade PP resins with an extremely low content of catalysts are required for BOPP films. As we can see from the above discussion of the multilayer film technology, commodity polymer resins with a small amount of impurity are still tolerable, if relatively thin PVDF layers are used and the impurity ions can be driven into the PC or PSF layers. Therefore, the multilayer film technology is more flexible in polymer choices for demanding dielectric applications. Interfacial Adhesion in Multilayer Films. Interfacial adhesion between the immiscible polymer layers is important for the ultimate dielectric properties of multilayer films. PC/ PSF and PVDF adhere reasonably well in multilayer films due to their polar chemical structures. However, the adhesion still needs further improvement. For example, dielectric breakdown of multilayer films can initiate layer delamination in the vicinity of the breakdown sites.118,120,121 The cavities from layer delamination are prone to further dielectric breakdown due to the low breakdown strength of gases.27 To mitigate this problem, a tie layer of PMMA is added between the PC and the P(VDF-HFP) layers.122 It is observed that the dielectric breakdown strength can be further enhanced and the hysteresis loss can be reduced when the PMMA is about 8 vol % (or 25 nm thick). The enhanced breakdown strength is attributed to the improved interfacial adhesion because PMMA is miscible with PVDF and compatible with PC.123 The result is later confirmed by a theoretical simulation,124 where modified interfaces with a tie layer having an intermediate permittivity can improve dielectric breakdown strength for multilayer films. In the future, it is desirable to further improve the interfacial adhesion by using PC-g-PMMA graft copolymers, which are yet to be synthesized. Biaxial Stretching To Further Enhance Dielectric Properties. Most commercial dielectric films, no matter semicrystalline or amorphous, need biaxial stretching to achieve uniform and thinner films ( Tg in this case). In a study on PET/P(VDF-TFE) multilayer films,128 the stretching temperature is selected as 105 °C, which is above the Tg of PET (75 °C) and just below the Tm (ca. 125 °C) of P(VDF-TFE). Simultaneous biaxial stretching can be performed on these multilayer films at a stretching rate of 100%/s and a stretching ratio up to 4.5 × 4.5. The fast stretching rate can induce PET crystallization and the formation of the rigid amorphous phase.129 The large stretching ratio can orient crystallites in both PET and P(VDF-TFE). All these factors are beneficial for reducing the charge carrier mobility and thus increasing the dielectric breakdown strength for multilayer films. In a recent report, the interfacial tie-layer strategy and biaxial stretching are combined in the PET/PMMA/P(VDF-HFP) multilayer system.108 Using similar biaxial stretching condition as those for the PET/P(VDF-TFE) multilayer films, the PET/ PMMA/P(VDF-HFP) multilayer films are successfully stretched into ca. 10 μm films with a stretching ratio of 4.5 × 4.5 and a stretching rate of 100%/s (Figure 10A). The PMMA tie layers improve the interfacial adhesion between PET and P(VDF-HFP). An optimal PMMA content is also found at 8 vol % (i.e., the PMMA layer is 25 nm thick); see Figure 10B. The biaxial stretching induced highly oriented edge-on crystallites in both PET and P(VDF-HFP) layers. Because of the formation of the rigid amorphous PET and the improved interfacial adhesion, the biaxially oriented PET/PMMA/ P(VDF-HFP) 65L films exhibit much enhanced discharged energy density (Ue) at breakdown (Figure 10B). Meanwhile, the hysteresis loop loss of the biaxially stretched PET/PMMA/ P(VDF-HFP) 65L film with 8 vol % PMMA is reduced compared to that of the as-extruded PET/P(VDF-HFP) 33L film (Figure 10C). The reduced dielectric loss is attributed to the formation of rigid amorphous phase in PET and the edgeon β PVDF crystals in P(VDF-HFP). Similar to the multilayer film concept, graft130−132 and block133 copolymers can also be used as potential dielectrics, and promising properties such as relatively high dielectric constants have been reported. The challenges lie in the largescale production and high costs of specific graft and block copolymers. Recently, layered nanofillers are incorporated in polymers, and enhanced dielectric properties such as increased breakdown strength are reported.49,134 The reason for enhance breakdown strength could be similar to that for multilayer films, i.e., interfacial polarization of space charges. However, largescale exfoliation and specific orientation of nanoplatelets in thin polymer films are still challenging tasks. Meanwhile, introduction of impurities should be strictly avoided when incorporating inorganic platelets into polymers. Finally, both approaches face a technological challenge when competing with BOPP films. Namely, it is difficult, if not impossible, to achieve biaxially

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CONCLUDING REMARKS AND OUTLOOK In summary, enhancing dielectric constant while keeping low dielectric loss for polymers has been challenging for practical electrical and power applications because increasing the overall polarization in polymers is often accompanied by a substantial increase of the dielectric loss in the samples. In this Perspective, we discuss various types of polarization in single-component dielectric polymers and their characteristics. Among all types of polarization, the orientational polarization is a promising approach to increase the dielectric constant of polymers close to that (εr = 80) of water with a reasonably low dielectric loss. The strategy is to enable various types of intermediate dipolar states between linear dielectrics and nonlinear ferroelectrics. The first promising candidate is a dipolar glass polymer, where dipolar groups are confined within the frozen free volume. As a result, the interaction among the confined dipolar groups is relatively weak and no FE domains can form. Relatively high dielectric constants (εr = 5−25) may be obtained without any hysteresis loss. The second promising candidate is a crystalline PE polymer, rather than a polar polymer melt, because the polar polymer melt exhibits high dielectric loss from dipolar switching and high leakage currents. Current crystalline PE polymers are only P(VDF-TrFE) and P(VDF-TFE) copolymers, and their dielectric constant can reach 30−40. The third promising candidate is an RFE polymer, where nanodomains in the isomorphic crystals are achieved by incorporating a large termonomer in P(VDF-TrFE). Narrow DHL and SHL are obtained as a result of the enlarged interchain distance and the effective pinning effect from the large termonomer units included in the crystals. Consequently, the apparent dielectric constant can reach 50−70. Although high dielectric constants have been achieved for polar polymers, the dielectric loss (tan δ ∼ 0.02) is still significantly higher than that (tan δ = 0.0002) of BOPP. To mitigate the high dielectric loss, the multilayer dielectric film technology is being pursued. The multilayer films contain alternating high κ (e.g., PVDF or the RFE terpolymers) and high breakdown/low loss polymers (e.g., PC and PSF). All types of polarization take place in multilayer films and can be divided into advantageous and disadvantageous groups. The advantageous polarizations include deformational (electronic/ atomic) polarization, dipolar polarization from the amorphous phase, and interfacial polarization (from electrons and ions). The disadvantageous polarizations include dipolar polarization from FE crystals and hopping polarization from impurity ions and free electrons. It is highly desirable to minimize the disadvantageous polarizations and promote the advantageous polarizations for multilayer films. First, the multilayer film is a good approach to minimize FE switching in PVDF crystals because the local field in high-κ PVDF layers is effectively reduced in multilayers having large permittivity contrast. Second, dielectric loss from impurity ion hopping polarization can be reduced by decreasing the PVDF layer thickness as a result of the nanoconfinement effect on ion transport. Third, dipolar polarization from the amorphous PVDF phase is advantageous to enhance the overall dielectric constant of multilayer films. In this sense, the RFE P(VDF-TrFE-X) M

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Macromolecules terpolymers are better than simple PVDF. Fourth, interfacial polarizations from both ions and electrons can modify or reduce the local electric field in the multilayer films. Consequently, enhanced dielectric breakdown strength and lifetime are observed. Fifth, improved interfacial adhesion by adding a tie layer in multilayer films can further enhance the ultimate dielectric performance. Finally, biaxial stretching is beneficial to achieve thin and high quality polymer films. With optimized biaxial stretching conditions, enhanced dielectric breakdown strength and reduced dielectric loss can be achieved for multilayer films. In the future, research effort should focus on the development of more affordable RFE polymers for practical applications, because P(VDF-TrFE-X) terpolymers are still expensive (>$10000/kg). Currently, nylon copolymers and terpolymers are being developed to achieve the RFE behavior with much lower prices. In addition to orientational polarization, high-κ polymers utilizing large atomic polarization are also promising to achieve intermediate dielectric constant with low dielectric losses.32,33 Finally, multilayer dielectric films show promise for translation into future products, provided the large-scale processing and manufacturing challenges are adequately addressed at the capacitor level.

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Lei Zhu is a Professor of Macromolecular Science and Engineering at CWRU. He received his Ph.D. degree in Polymer Science in 2000 from the University of Akron. He is Associate Director for Research for CLiPS. His research interests include high κ/low loss dielectric and ferroelectric polymers for electric energy storage (in the area of pulsed power and power electronics) and electromechanical actuation, polymer-grafted inorganic nanoparticles and nanocomposites, and supramolecular self-assembly of liquid crystalline polymers. He is recipient of National Science Foundation (NSF) Career Award, 3M Nontenured Faculty Award, and DuPont Young Professor Award. He serves on the editorial board of Macromolecules/ACS Macro Letters and Polymer.

AUTHOR INFORMATION

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Corresponding Author

*E-mail: [email protected] (L.Z.).

ACKNOWLEDGMENTS The work on novel ferroelectric polymers is supported by NSF Grants DMR-0907580 and DMR-1402733. The work on fundamental understanding of orientational and impurity ionhopping polarizations in multilayer dielectric films is supported by NSF through the Science and Technology Center (CLiPS) under Grant DMR-0423914. The work on multilayer film capacitors for pulsed power applications is supported by Office of Naval Research (N00014-16-1-2170). The authors thank Dr. Elshad Allahyarov for theoretical and simulation studies, Professor Donald Schuele for insightful discussion, and students/postdoctoral fellows for their original work.

ORCID

Lei Zhu: 0000-0001-6570-9123 Notes

The authors declare no competing financial interest. Biographies

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ABBREVIATIONS Ar-PTU aromatic polythiourea AFM atomic force microscopy BOPP biaxially oriented polypropylene Eb breakdown strength XLPE cross-linked polyethylene TC Curie transition temperature CEP cyanoethylated pullulan DG dipolar glass DHL double hysteresis loop D−E loop electric displacement−electric field loop EDL electric double layer [EMIM][TFSI] 1-ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide FE ferroelectric FET field effect transistor FEP fluorinated ethylene−propylene Tg glass transition temperature h-BN hexagonal boron nitride HVDC high-voltage direct current PMN lead manganese niobate [Pb(Mg1/3Nb2/3)O3]

Eric Baer is the Herbert Henry Dow Professor of Macromolecular Science and Engineering at CWRU. He earned his doctoral degree in Engineering from Johns Hopkins University. He is Director of the National Science Foundation (NSF) Science and Technology Center of Layered Polymeric Systems (CLiPS). His research interests include multilayered polymer films for barrier, optical, and electrical applications, structure−property relationships of hierarchical polymer and biological systems, mechanical properties of soft connective tissue, and damage and fracture analysis of polymers and their composites. He is the recipient of numerous professional awards, including Distinguished University Professor of CWRU, ACS Award in Applied Polymer Science, and the James L. White Innovation Award of the Polymer Processing Society. N

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PE PAA PC PDPMA

paraelectric poly(acrylic acid) polycarbonate poly(2,3-dihydroxylpropyl methacrylate) PEO poly(ethylene oxide) PET poly(ethylene terephthalate) PMMA poly(methyl methacrylate) PMSEMA poly[2-(methylsulfonyl)ethyl methacrylate] POSS polyoctahedral selsisquioxane PS polystyrene PSF polysulfone PTFE polytetrafluoroethylene PVA poly(vinyl alcohol) PVDF poly(vinylidene fluoride) P(VDF-HFP) poly(vinylidene fluoride-co-hexafluoropropylene) P(VDF-TFE) poly(vinylidene fluoride-co-tetrafluoroethylene) P(VDF-TrFE) poly(vinylidene fluoride-co-trifluoroethylene) P(VDF-TrFE-CFE) poly(vinylidene fluoride-co-trifluoroethylene-co-1,1-chlorofluoroethylene) P(VDF-TrFE-CTFE) poly(vinylidene fluoride-co-trifluoroethylene-co-chlorotrifluoroethylene) QTC quantum tunneling composite RFE relaxor ferroelectric SHL single hysteresis loop TEM transmission electron microscopy VHB very high bond

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DOI: 10.1021/acs.macromol.6b02669 Macromolecules XXXX, XXX, XXX−XXX