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Tuning Phase Composition of Polymer Nanocomposites towards High Energy Density and High Discharge Efficiency by Non-Equilibrium Processing Jianyong Jiang, Xin Zhang, Zhenkang Dan, Jing Ma, Yuanhua Lin, Ming Li, Ce-Wen Nan, and Yang Shen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b07963 • Publication Date (Web): 15 Aug 2017 Downloaded from http://pubs.acs.org on August 15, 2017

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ACS Applied Materials & Interfaces

Tuning Phase Composition of Polymer Nanocomposites towards High Energy Density and High Discharge Efficiency by Non-Equilibrium Processing

Jianyong Jiang,† Xin Zhang,† Zhenkang Dan,† Jing Ma, † Yuanhua Lin,† Ming Li,† Ce-Wen Nan,†,* and Yang Shen†,*

† State Key Lab of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, 100084, China *Corresponding authors. E-mail: [email protected]

Abstract Polymer nanocomposite dielectrics with high energy density and low loss are major enablers for a number of applications in modern electronic and electrical industry. Conventional fabrication of nanocomposites by solution routes involves equilibrium process which is slow and results in structural imperfections, hence high leakage current and compromised reliability of the nanocomposites. We propose and demonstrate that a non-equilibrium process, which synergistically integrates electrospinning, hot-pressing and thermal quenching, is capable of yielding nanocomposites of very high quality. In the non-equilibrium nanocomposites of poly(vinylidene fluoride-co-hexafluoropropylene) (P(VDF-HFP)) and BaTiO3 nanoparticles (BTO_nps), an ultrahigh Weibull modulus β of ~ 30 is achieved, which is comparable to the quality of the bench-mark biaxially-oriented polypropylene (BOPP) fabricated with melt-extrusion process by much more sophisticated and expensive industrial apparatus. Favorable phase composition and small crystalline size are also induced by the non-equilibrium process, which leads to concomitant enhancement of electric displacement and breakdown strength of the nanocomposite hence a high energy density of ~ 21 J/cm3. Study on the polarization behavior and phase transformation at high electric field indicates that BTO_nps could facilitate the phase transformation from α to β polymorph at low electric field.

Keywords: non-equilibrium process, phase tuning, nanocomposite dielectrics, energy density, discharge efficiency.

Introduction Owing to their high effective dielectric permittivity, high breakdown strength and hence high energy density, ACS Paragon Plus Environment


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polymer nanocomposite dielectrics are emerging as critical energy storage media in electrostatic capacitors for a number of modern electronic and electrical applications,1-7 such as pulsed power system,8 smart grid9 and electric (or hybrid) vehicles.10,11 With ever increasing needs for miniaturization and lightweight of power modules, high energy density dielectrics are highly desirable for reducing the size and weight of electrostatic capacitors. For an electrostatic capacitor with a dielectric layer between two electrodes, the discharge energy density (Ue) is determined simultaneously by electric displacement (D) and electric field (E) as5,12 =

(1)

The maximum Ue is then determined by the breakdown strength (Eb) which signifies the highest electric field applicable to the dielectrics before dielectric failure. With D related to the dielectric permittivity (εr) as: D =

ε0εrE,13 ultrahigh Ue could only be achieved by concomitant enhancement of Eb and εr, which has been the Holy Grail for dielectrics. Hence, in addition to having a high Eb, a high electric displacement is also desirable for achieving high energy density. For this reason, despite of their high Eb, pure polymer dielectrics exhibit rather low Ue as a result of their low dielectric permittivity (εr ~ 1-2). For instance, owing to its low dielectric permittivity (εr ~ 2-3), biaxial-oriented polypropylene (BOPP), the bench-mark polymer dielectric of current use, bears a low Ue of 1-3 J/cm3 albeit its ultrahigh Eb (>700 MV/m).2,14 In stark contrast, inorganic dielectrics, such as ferroelectric ceramics (BaTiO3,15 Pb(Zr,Ti)O3,16 PbMg1/3Nb2/3O3-PbTiO317 et.al.,) exhibit high dielectric permittivity (εr ~ 102-104). But their maximum Ue is substantially compromised by the low Eb as a result of the structural imperfections introduced during processing.2,18 A straight forward and promising approach towards dielectrics of high Ue is to fabricate polymer nanocomposites composed of polymer matrix (with high Eb) and inorganic nanometric fillers (with high εr).9 Among all the dielectric polymers explored so far, poly(vinylidene fluoride) (PVDF) and its random co-polymers, such as poly(vinylidene fluoride-co-hexafluoropropylene) (P(VDF-HFP)) or poly(vinylidene fluoride-co- chlorotrifluoroethylene) (P(VDF-CTFE)), are considered as promising candidates for ultrahigh energy density for their high intrinsic polarizability (i.e., high εr ~ 10-20) and high Eb (~ 600-800 MV/m).19-23 However, their discharge energy density and discharge efficiency are substantially compromised by their high ferroelectric loss, e.g., a large portion of the electrical energy stored during the charging process may not be recovered but stored in the form of remnant polarization during the discharging process.21 To realize the true potential of PVDF for high energy storage performances, tremendous efforts have been made to tune the phase composition and crystalline orientations in order to suppress the ferroelectric loss. In principle, there are four polymorphs of PVDF and its copolymers with different chain conformations, i.e., the paraelectric (non-polar) α-phase (TG+TG-) and the ferroelectric (polar) β (all-trans), γ (TTTG+TTTG-), and δ-phase (TG+TG-), where T (trans) and G± (gauche±) refer to the torsional bond arrangements with substituents at 180° and ±60° to each other, respectively.24,25 The primary

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ACS Applied Materials & Interfaces dipolar moments in these polymorphs are induced by the alternating –CH2 and –CF2 groups along the polymer chains. In paraelectric α-phase, the neighboring dipoles are in the opposite direction and cancel each other, resulting in no net dipole in the unit cell hence of low polarizability and low remnant polarization. As for the β, γ, and δ-phases, the alignment of the neighboring dipoles gives rise to net dipole in the unit cell hence high polarizability of these ferroelectric polymorphs. It is also worth noting that these primary dipolar moments are perpendicular to the polymer chains. When c-axis of the PVDF crystal, which is normally the same direction as the polymer chains, is perpendicular to the electric field, all the dipolar moments could be easily aligned in the direction of the electric field by either crank-shaft motions or chain flipping,26 giving rise to high polarizability and high stored electric energy.27 Tuning the orientation of the PVDF crystals is thus considered a viable approach toward higher energy density of PVDF and its co-polymers. Along this line, Zhu et.al., explored the implications of crystal orientation on the dielectric and energy storage performances of P(VDF-HFP). Anisotropic polarizability is indeed observed in P(VDF-HFP) of different orientations, e.g., higher polarizability when c-axis is perpendicular to the electric field and vice-versa. Also, it is worth noting that there is a lower limit of coercivity about 50 MV/m during the ferroelectric switching of copolymer.28 A number of fabrication and processing techniques are employed to modulate the crystal orientation of the P(VDF-HFP), such as solution-casting, uniaxial stretching, melt-pressing and thermal quenching.27 However, the effects of nanoinclusions, such as nanoparticles or nanofibers, on the crystal orientation of PVDF and its co-polymers still remain ambiguous. To answer this question, the very first step has been to obtain homogeneous dispersion of nanoinclusions in polymer matrix. Since T.J. Lewis first proposed the concept of “nanometric dielectrics” (or polymer nanocomposite dielectrics) in 1994,29 a number of different techniques have been employed for the fabrication of polymer nanocomposites, such as hot-pressing,30 solution casting31-34 and in-situ polymerization with presence of nanoinclusions,35,36 etc. Among all these methods, solution casting is the most widely used technique for the fabrication of polymer nanocomposite films because it is much easier to achieve homogeneous dispersion of nanoparticles in the low-viscosity solution of polymers.37 The typical procedure of solution casting technique is illustrated in Scheme 1a for the fabrication of BaTiO3 nanoparticles (BTO_nps for short hereafter)/P(VDF-HFP) nanocomposites. Certain amount of BTO_nps is first dispersed by ultrasonic treatment in the solution of P(VDF-HFP) in organic solvent. BTO_nps/P(VDF-HFP) suspension is obtained with the aid of magnetic stirring. Polymer nanocomposite films are then cast from suspension on glass substrates. The as-cast film is subject to heat treatments for the evaporation of the solvents and the condensation of polymer. We note that the solution casting process is equilibrium in thermodynamics. During the slow drying process for the evaporation of solvents, the nanoparticles of high surface energy are allowed to migrate and coalesce into large agglomerates.38,39 Sedimentation of the nanoparticles during the drying process is another issue with the solution casting process, which results in inhomogeneity along the out-of-plane directions of the composite films. To improve the compatibility between ACS Paragon Plus Environment


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nanoparticles and the polymer matrix hence suppress the coalescence of nanoparticles, surface modification of the nanoparticles has been performed by using a number of surface modifiers, such as dopamine,4,40-42 ethylene diamine,33,43 H2O244 and phosphonic acid,38 etc. Unfortunately, despite of the homogeneous dispersion of nanoparticles, these surface layers also induce inhomogeneity in the distribution of local electric field because of their low dielectric permittivity as compared with nanoinclusions and polymer matrix.31 The concentrated local electric field then leads to breakdown of the nanocomposites at electric fields well below the intrinsic Eb of the polymer matrix.45 Moreover, residual surface modifier molecules could also migrate or even be ionized into mobile charge carriers under high electric fields, which may results in high leakage current, hence high loss, and further compromise Eb of nanocomposites. In this contribution, a non-equilibrium process is proposed and demonstrated to achieve homogeneous dispersion of BTO nanoparticles in P(VDF-HFP), and more importantly, to tune the phase composition of P(VDF-HFP) in the nanocomposites. As illustrated in Scheme 1b, P(VDF-HFP) nanofibers embedded with BTO nanoparticles are prepared by modified electrospinning and are assembled into BTO_nps/P(VDF-HFP) fibrous mats. In this process, the non-equilibrium fast evaporation of solvents during the electrospinning of the P(VDF-HFP) nanofibers gives rise to homogeneous dispersion of BTO nanoparticles in P(VDF-HFP) nanofibers. In the meantime, these P(VDF-HFP) nanofibers are also uniaxially stretched along the direction of the electric field by the large electrostatic force during electrospinning. The field-assisted stretching is in favor for the orientation of c-axis along the in-plane directions of BTO_nps/P(VDF-HFP) nanocomposite films, which is highly desirable for high polarizability of the nanocomposites. The following hot-pressing treatment transforms the porous fibrous mats into dense polymer nanocomposite films of high quality and high structural integrity. In the last step, a non-equilibrium thermal quenching process leads to the formation of paraelectric α-phase. Through this non-equilibrium (NEQ for short) technique, we are able to achieve unprecedented high quality of polymer nanocomposites by a low-cost solution process. As a result, the BTO_nps/P(VDF-HFP) nanocomposite films exhibit a rather high β-parameter of ~ 35, which is the key parameter that reflects the uniformity and quality of dielectrics. This is the highest

β-parameter ever achieved for nanocomposite dielectrics, which is even comparable with the β-parameter of BOPP films (β ~ 40) fabricated through melt-extrusion & stretching technique by highly complicated and expensive apparatus in atmosphere-controlled clean rooms.46 Furthermore, this non-equilibrium process could be well extended to other nanocomposite systems with different kinds of nanofillers. For instance, a nanocomposite film with parallel BaTiO3 nanofibers (BTO_nfs) is shown in Scheme 1b.

Experimental Section Materials. Chemicals were obtained from the following commercial sources and used without further purification: BaTiO3 ACS Paragon Plus Environment

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ACS Applied Materials & Interfaces (BTO, ~50 nm, seen in Figure S1a), poly(vinylidene fluoride-co-hexafluoropropylene) (P(VDF-HFP), Arkema, France, Kynar Flex 2801, with 10 wt.% HFP). All other solvents were purchased from China National Chemicals Corporation Ltd. and used without further purification. Fabrication of polymer nanocomposite films with solution-casting method. P(VDF-HFP) powder was thoroughly dissolved in a mixed solvent of N,N-dimethylformamide (DMF) and acetone. With the aid of ultrasonic treatment, BTO_nps were then dispersed into the P(VDF-HFP) solution with different volume fractions (1 vol.%, 5 vol.% and 10 vol.%). After magnetic stirring for 12 hrs, a stable composite suspension was obtained. Then, the suspension was dropped onto a glass substrate and cast into films with a doctor blade. After being dried in vacuum at 40 °C, the composite films were heated at 200 °C for 7 min followed by quenching in icewater. The thickness of the final composite films was ~ 13 µm. Fabrication of polymer nanocomposite films with non-equilibrium process. The electrospinning solution was the same as those for casting mentioned above and then transferred into a syringe. The modified electrospinning process with constantly magnetic stirring was performed at an applied field of 1.0 kV/cm with a flow rate of 1.0 mL/h for 3 hrs.4 Meanwhile, a lateral movement of the syringe was carried out. The nanofibers were collected and assembled into a fibrous mat on a rolling receiver, rotating at 300 rpm. The fibrous mats were then transformed into dense body by hot-pressing at 200 °C for 1 hr under a pressure of 10 MPa. The as-pressed composite films were annealed at 200 °C for 7 min followed by ice-water quenching. The thickness of the final composite films was ~13 µm. Instrumentation and Characterization The morphology of P(VDF-HFP)_nfs, BTO_nps@P(VDF-HFP)_nfs (which means the BTO nanoparticles are dispersed in P(VDF-HFP) nanofibers) and nanocomposite films were characterized with scanning electron microscopy (SEM, ZEISS MWRLIN compact). Wide-angle X-ray diffraction (WAXD, Rigaku SmartLab), Fourier transform infrared spectroscopy (FTIR, Nicolet 6700), differential scanning calorimetry (DSC, TA Q-2000) were performed for the nanocomposite films. Dielectric properties were measured by employing a HP 4294A precision impedance analyzer (Agilent Technologies) at room temperature within the frequency range of 102 to 108 Hz at 1

V . Electric breakdown tests were performed with dielectric withstand voltage test system (Beijing

Electro-mechanical Research Institute Supervoltage Technique) at a ramping rate of 200 V/s and a limit current of

5 mA. Unipolar displacement-electric field (D-E) hysteresis loops were collected at 10 Hz with a Premier II ferroelectric test system (Radiant Technologies, Inc.).

Results and Discussion Morphology of nanocomposite films. The pure P(VDF-HFP) and BTO_nps@P(VDF-HFP) nanofibers were prepared by a modified electrospinning ACS Paragon Plus Environment


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process. As shown in Figure 1a, the pure P(VDF-HFP) nanofibers are fairly smooth and straight. These morphologies are still retained after incorporating with various volume fractions of BTO_nps (Figure 1(d, g, j)). Also observed in the SEM images is the homogeneous dispersion of BTO_nps within the P(VDF-HFP) nanofibers. After hot-pressing treatment, the fibrous mats are transformed into dense nanocomposite films of high structural integrity. The surface (Figure 1(b, e, h, k)) and cross-sectional (Figure 1(c, f, i, l) & Figure S2 in the supporting information) SEM images reveal that both the pure P(VDF-HFP) film and the BTO_nps/P(VDF-HFP) nanocomposite films are dense and free of structural imperfections, such as pores or voids, and the BTO_nps without surface modification are homogeneously dispersed in P(VDF-HFP) matrices without obvious aggregation. As shown in Figure S2, BTO_nps are homogeneously distributed throughout the whole cross sections. The advantages of the fast non-equilibrium process lie in the following two aspects: i) the rapid evaporation of solvents during the electrospinning process, hence the substantially increased viscosity of the composite slurry, prevents the coalescence of BTO_nps into large agglomerates in the BTO_nps@P(VDF-HFP) nanofibers, giving rise to homogeneous distribution of BTO_nps in the in-plane directions of the nanocomposite films; ii) The BTO_nps@P(VDF-HFP) fibrous mats are assembled in a layer-by-layer manner, which ensures homogeneous dispersion of BTO_nps along the out-of-plane direction of the nanocomposite films, as illustrated in Scheme 1b. For comparison, BTO_nps/P(VDF-HFP) nanocomposite films are also prepared by conventional equilibrium process (or the solution-casting method, denoted as EQ for short here and after). As evidenced in the SEM images of the top and bottom surfaces (Figure 2a&2b) and cross-section (Figure 2c&2d) of the P(VDF-HFP) nanocomposites loaded with 5 vol.% BTO_nps, large agglomerates of BTO_nps are readily distinguished close to the bottom surface of the BTO_nps/P(VDF-HFP) nanocomposites, which is in stark contrast to the homogeneous distribution of BTO_nps in the nanocomposites prepared through non-equilibrium process. This phenomenon is more serious with increasing volume fraction of BTO_nps in the nanocomposites, as shown in Figure S3.

Characterization of crystalline morphology. The implication of non-equilibrium process on the phase composition of the BTO_nps/P(VDF-HFP) nanocomposites is first studied with Fourier transformed infrared spectroscopy (FTIR) in the wavenumber range of 700-1500 cm-1. As shown in Figure 3, pure P(VDF-HFP) films prepared by equilibrium process crystallize predominantly into the α−phase, as evidenced by the absorption bands associated with α-phase at 763, 796, 976 cm-1 ( black solid curve in Figure 3).47,48 Same FTIR bands are observed for the BTO_nps/P(VDF-HFP) nanocomposite films prepared by equilibrium process (see Figure S4a). As for the pure P(VDF-HFP) films fabricated through non-equilibrium process, in addition to the four bands from α-phase, two additional bands are observed at 840 and 1279 cm-1 (red solid curve in Figure 3) and assigned to β-phase.47,48 Non-equilibrium process also induces β-phase in the BTO_nps/P(VDF-HFP) nanocomposites, as shown in Figure S4b. The phase ACS Paragon Plus Environment

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ACS Applied Materials & Interfaces composition of the nanocomposites prepared through different process is further investigated with wide angle X-ray diffraction (WAXD) measurement. As shown in Figure S5a, the WAXD profile of the pure P(VDF-HFP) film prepared by equilibrium process shows two distinct peaks at around 2θ = 18.3°, 19.9°, which are indexed to (020) and (110) reflections of the α-phase.48 No peaks of β-phase could be observed, indicating the presence of primarily nonpolar α-phase in pure EQ film. Very similar WAXD patterns are observed for the BTO_nps/P(VDF-HFP) nanocomposite films prepared by equilibrium process, except for the additional peak at 2θ = 22.0° corresponding to the (100) reflection of BTO_nps (as shown in Figure S1b). Two striking features could be readily distinguished in the WAXD profiles for the pure P(VDF-HFP) and BTO_nps/P(VDF-HFP) nanocomposite (Figure S5b) films prepared by non-equilibrium process. First, the peak at 2θ = 19.9° shifts to 2θ = 20.2°, indicating the formation of β-phase in the non-equilibrium films.48 Second, non-equilibrium process induces smaller crystalline size in the nanocomposite films, as evidenced by larger full width at half maximum (FWHM) of the two major peaks in the non-equilibrium films compared to those in the equilibrium films. The exact crystalline size and the content of α and β phases are then determined by fitting the WAXD profiles with a sum of Gaussians and Lorentzians.9,49,50 An automatic constant baseline is subtracted according to the following formula: 1 − y = − + 2

$ − %&'() − 1+" #

(2)

where a0 is the amplitude, a1 is the center and a2 is the width.51

The results of WAXD fitting are presented in Figure 4. The contents of α and β phases and the average crystalline size determined from the fitting are summarized in Table 1. It is worth noting that both pure P(VDF-HFP) and the corresponding nanocomposite films prepared through non-equilibrium process exhibit higher content of β-phase and smaller crystalline size as compared to their equilibrium counterparts. For instance, there is almost no β-phase in equilibrium pure P(VDF-HFP) film, while a significant amount of 21.5 wt.% β-phase is present in the non-equilibrium pure P(VDF-HFP) film. The apparent crystalline size along [110]α[D(110) α] of the non-equilibrium pure P(VDF-HFP) film is ~ 15.3 nm, which is less than half of that in equilibrium pure P(VDF-HFP) film (~ 32.4 nm). We note that the stretching of the P(VDF-HFP) nanofibers during the non-equilibrium process is in favor of the formation of β-phase,48 while the rapid evaporation of solvents also prevents the growth of P(VDF-HFP) crystal grains hence leading to smaller crystalline size of α-phase in the non-equilibrium films. Figure 5 shows DSC heating and cooling thermograms for both equilibrium and non-equilibrium films. The DSC heating curves show one main melting peak at 143.0 °C (Tm) and either the addition of BTO_nps or the fabrication method seems to change Tm. However, with increasing BTO_nps contents, the degree of crystallinity decreases in both equilibrium and non-equilibrium films (inset of Figure 5), which is in line with WAXD and FTIR ACS Paragon Plus Environment


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results. In fact, the same phenomenon has also been observed in CNF/PVDF, CB/P(VDF-HFP) nanocomposites.52,53 It may be attributed to the inhibition role of nanoinclusions in reorganization of polymer chains. Here, the degree of crystallinity Xc is determined by the following equation: *+ = ,

∆./ ∅ ∆./

1 × 100%

(3)

∅ where ∆Hm is the melting enthalpy of the films and ∆./ is the melting enthalpy of 100% crystalline material, i.e.,

104.7 J/g.52-54 The overall degree of crystallinity of the equilibrium and non-equilibrium films is also listed in Table 1. The non-equilibrium films exhibit slightly higher overall degree of crystallinity (~ 44.5%) compared to their equilibrium counterparts (~ 42.6%). From the DSC cooling curves (Figure 5), it is of interest to note that the addition of BTO_nps in equilibrium films leads to a shift of crystallization temperature, Tcrystallization, towards a

lower temperature, while an opposite tendency is observed for non-equilibrium films. In general, during the recrystallization process of polymer, the fillers play a dual role depending on both their quantity and state of dispersion. In more detail, they can act not only as nucleating agents but physical barriers to the motion of polymer chains and crystal growth.49 In the non-equilibrium films, homogeneous dispersion of BTO_nps gives rise to more interfaces, hence larger interfacial area between BTO_nps and polymer, which facilitates nucleation of polymer on the surface of BTO_nps and raises the crystallization temperature of the non-equilibrium films. As for the equilibrium films, the large agglomerates of BTO_nps slow down the recrystallization process, leading to lower Tcrystallization.

Dielectric properties at low electric field. The frequency dependence of room temperature dielectric permittivity and dielectric loss for equilibrium and non-equilibrium films are shown in Figure 6. In both cases, the dielectric permittivity of the nanocomposites increases consistently with increasing BTO_nps content over the whole frequency range (Figure 6a), which is attributed to the introduction of BTO_nps of high dielectric permittivity. At a fixed frequency of 1 kHz, the dielectric permittivity of all nanocomposites as a function of the BTO_nps loading is presented in the inset of Figure 6a. As seen, compared with equilibrium films, the non-equilibrium films with the same BTO_nps loading shows a higher dielectric permittivity. We speculate that the addition of the polar β-phase in the non-equilibrium films leads to a higher dielectric permittivity of polymer matrix, which is evidenced by the higher dielectric permittivity of the pure non-equilibrium P(VDF-HFP) film (εr ~10.0, @1 kHz) than that of the pure equilibrium film (ε ~9.5, @1 kHz). More importantly, the formation of agglomeration in the equilibrium films may decreases the interfacial area between BTO_nps and P(VDF-HFP) matrix and thus suppresses the interfacial polarization. It also causes pores and voids of lower εr in the nanocomposites, which further compromises the dielectric permittivity of the nanocomposites. This is especially so at high BTO_nps loadings.


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ACS Applied Materials & Interfaces Figure 6b shows the dielectric loss of equilibrium films and non-equilibrium films with varying volume fractions of BTO_nps over the frequency range of 100 Hz-100 MHz. Both the equilibrium and non-equilibrium films show two relaxation peaks at low (~ 100 Hz) and high (~ 10 MHz) frequencies, respectively. The relaxation peak at low frequencies (~ 1 Hz) is the so-called αc-relaxation, which is attributed to the conformational changes between TG+TG- and G+TG-T of VDF-chain segments within the nonpolar crystalline α-phase.55,56 According to the aforementioned results of WAXD and FTIR, there is a reduction in the content of α-phase in the non-equilibrium films compared to the equilibrium films. Minor αc-relaxation is thus observed for the non-equilibrium films. The

αa-relaxation at high frequencies (~ 10 MHz) is originated from the segmental motions within the amorphous phase of P(VDF-HFP).55 The lower degree of crystallization in the equilibrium pure P(VDF-HFP) films (~ 42.6%, as listed in Table1), hence higher content of amorphous phase, results in stronger αa-relaxation at high frequencies. The implication of BTO_nps on the αa-relaxation lies in two folds. First, upon the incorporation of BTO_nps, the content of amorphous phase increases, which may induce stronger αa-relaxation process. Second, BTO_nps may also suppress the αa-relaxation process by pinning the segmental motions in the amorphous phase through the interactions between the surface of BTO_nps and the polymer matrix. We speculate that the second mechanism plays a dominant role in our case. Therefore, the αa-relaxation peak amplitude decreases with increasing content of BTO_nps for both the equilibrium and non-equilibrium films.

Dielectric breakdown behavior of the nanocomposites The breakdown strength is the key parameter in determining the energy storage density of nanocomposites. In this study, the breakdown strength was also measured for the equilibrium and non-equilibrium films with BTO_nps volume fractions from 0 to 10% and analyzed using two-parameter Weibull statistics. The Weibull cumulative distribution function, P, is defined as50 567 = 1 − exp −

< ;

(4)

where E is experimental breakdown strength, Eb is a scale parameter refers to the breakdown strength at the cumulative failure probability of 63.2% that is also regarded as the characteristic breakdown strength, and the shape parameter β is the Weibull modulus that shows dispersion of E and is also considered as a bench mark for the quality of the dielectrics. The results of Weibull statistical analysis for the equilibrium and non-equilibrium films are shown in Figure 7. The most striking feature of Figure 7a is the high Weibull modulus achieved for the non-equilibrium films, e.g., β ~ 35.0 is obtained for the pure P(VDF-HFP) films. To the best of our knowledge, this is the highest β value ever achieved in dielectric polymer films prepared through a solution approach, which indicates that the films are of high quality and high structural integrity. Also could be readily observed in Figure 7a is the substantially enhanced Eb of the BTO_nps/P(VDF-HFP) nanocomposites by the incorporation of BTO_nps.



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As seen, Eb increases from ~ 620.4 MV/m for the pure P(VDF-HFP) films to ~ 651.9 MV/m for the nanocomposite filled with 1 vol.% of BTO_nps, followed by gradual decrease to a minimum of 519.2 MV/m with further addition of BTO_nps up to 10 vol.%. The improvement of breakdown strength at a low concentration of nanoparticles has been observed for a number of nanocomposites.57-60 For the BTO_nps/P(VDF-HFP) nanocomposites, the polymer chains are tightly bonded to the surface of BTO_nps within the interface between P(VDF-HFP) matrix and BTO_nps. Within this interfacial region, the chain mobility is decreased such that the transfer of charge carriers through the mobile polymer chains is suppressed, giving rise to enhanced breakdown strength of the nanocomposites. However, the introduction of BTO_nps also induces concentration of local electric field close to the BTO/P(VDF-HFP) interfacial regions owing to the mismatch between the dielectric “soft” BTO_nps of high εr and the dielectric “hard” P(VDF-HFP) of low εr.5,35,38,45 With the increasing volume fraction of BTO_nps in the P(VDF-HFP) matrix, the interfacial regions of concentrated local electric field may gradually percolate and eventually form a networks at a critical volume fraction of BTO_nps. During this process, the nanocomposites exhibit monotonically decreased breakdown strength as more BTO_nps are incorporated (see Figure 7). To better distinguish the advantages of non-equilibrium process, Eb and β of the equilibrium and non-equilibrium films are summarized and compared in Figure 8. There are two striking features to the data when presented this way in Figure 8. First, the non-equilibrium nanocomposite films exhibit consistently higher Eb and β at all volume fractions of BTO_nps compared to their equilibrium counterparts. Second, monotonic decrease in Eb is observed for the equilibrium nanocomposite films with increasing volume fraction of BTO_nps while an enhanced Eb is achieved for the non-equilibrium nanocomposite loaded with 1 vol.% of BTO_nps. Although further increasing the volume fraction of BTO_nps also results in decrease in Eb of the non-equilibrium nanocomposites, a high Eb of ~ 519.2 MV/mm is still obtained at 10 vol.% of BTO_nps. Same amount of BTO_nps results in a much compromised Eb of ~ 354.0 MV/mm for the equilibrium nanocomposites, which is decreased by ~ 41.4% over Eb of the equilibrium pure P(VDF-HFP) films (~ 603.6 MV/mm). The large agglomerates of BTO_nps in the equilibrium nanocomposite films not only compromise the favorable bonding of polymer chains on the surface of BTO_nps but results in more serious concentration of local electric field even at low volume fractions. We note that the non-equilibrium process and the layer-by-layer fabrication of the nanocomposite films promote the homogenous dispersion of BTO_nps in both in-plane and out-of-plane directions of the nanocomposites. During the equilibrium process, there is a wetting phenomenon due to the difference of between solution and glass substrate in surface tension. This phenomenon will lead to the final composite films with thick middle part and thin edge section, which is detrimental to the thickness uniformity of the nanocomposite films. While the non-equilibrium process can avoid such a phenomenon. These benefits raise the nanofiller-homogeneity and thickness-uniformity of the nanocomposites films to an unprecedented high level, giving rise to the superior dielectric breakdown performance of the non-equilibrium nanocomposite films. ACS Paragon Plus Environment

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High-field polarization behavior and energy storage performances The high-field polarization behaviors of the pure polymer and nanocomposites are characterized by the electric displacement (D)-electric field (E) loops (Figure S6) obtained by a modified Sawyer-Tower circuit.31 As shown in Figure 9, the addition of BTO_nps induces enhanced D in both the non-equilibrium and equilibrium nanocomposite films, which is in line with the low-field dielectric behaviors presented in the previous sections. In general, the equilibrium nanocomposites exhibit higher D than their non-equilibrium counterparts with the same volume fraction of BTO_nps and at the same electric field. For instance, with 5 vol.% of BTO_nps, the equilibrium nanocomposites exhibit a D of ~ 8.9 µC/cm2 at 400 MV/m while its non-equilibrium counterparts exhibit D of ~ 7.5 µC/cm2. Unfortunately, the higher D of the equilibrium nanocomposites is normally the results of accumulated space charges induced by structural imperfections, such as pores or voids, at elevated electric fields. This is evidenced by the substantially increased remnant polarization (Dr) of the equilibrium nanocomposites with the increasing volume fraction of BTO_nps. The variation of Dr with electric field is shown in Figure 10a for both equilibrium and non-equilibrium nanocomposites with different loading of BTO_nps. Dr of the pure equilibrium nanocomposites increases monotonically with increasing electric field from ~ 0.01 µC/cm2 at 100 MV/m to ~ 1.1 µC/cm2 at 600 MV/m. This is in stark contrast to the scenario for the pure non-equilibrium nanocomposites, where Dr increases quickly from ~ 0.01 µC/cm2 at 100 MV/m to ~ 0.76 µC/cm2 at 300 MV/m then rather mildly to ~ 0.80 µC/cm2 at 600 MV/m. It is of interest to note that same trend is also observed in the variation of leakage current with electric field for the two types nanocomposites (see Figure 10b). The leakage current of the pure equilibrium nanocomposites increases consistently from ~ 0.2 × 10-7 A/cm2 at 50 MV/m to ~ 5.6 × 10-7 A/cm2 at 500 MV/m. As for the pure non- equilibrium nanocomposites, the leakage current rapidly increases from ~ 0.2 × 10-7 A/cm2 at 50 MV/m to ~ 1.2 × 10-7 A/cm2 at 250 MV/m then very mildly to ~ 1.9 × 10-7 A/cm2 at 500 MV/m. We are thus led to the conclusion that the variation of Dr with electric field is primarily associated with the intrinsic leakage current of polymer matrix. The incorporation of BTO_nps in the equilibrium nanocomposites, and the resultant structural imperfections, results in even more prominent increase in Dr. For instance, at a fixed electric field of ~ 300 MV/m, Dr increases drastically from ~ 0.8 µC/cm2 to ~ 2.1 µC/cm2 with the increasing amount of BTO_nps from 1 vol.% to 10 vol.%. As for the non-equilibrium nanocomposites, Dr also increases mildly with increasing amount of BTO_nps but is consistently lower than its equilibrium counterparts at all electric field. Even loaded with 5 vol.% of BTO_nps, the non-equilibrium nanocomposites only exhibit a Dr of ~ 1.0 µC/cm2 at 400 MV/m which is only 40% of the Dr of its equilibrium counterparts (~ 2.5 µC/cm2). The suppressed Dr is in favor for higher discharged energy density (Ue) of the nanocomposites as more electric energy stored during the charging process could be discharged instead of being stored in the form of remnant polarization. Ue is determined from the D-E loops and plotted as a function of electric field for both equilibrium ACS Paragon Plus Environment


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and non-equilibrium nanocomposites in Figure 11. Indeed, the non-equilibrium nanocomposites deliver a higher Ue compared to their equilibrium counterparts (Figure 11). For instance, the maximum Ue can reach up to 20.6 J/cm3 for the nanocomposites with 1 vol.% of BTO_nps at 648 MV/m, which represents an enhancement of ~13% over the pure P(VDF-HFP) (Ue ~18.3 J/cm3 at 620 MV/m) and is even 1617% greater than Ue of the bench-mark BOPP (~ 1.2 J/cm3 at 640 MV/m).33 Further increase in the content of BTO_nps may give rise to enhanced D hence higher Ue at an electric field below Eb. Yet, the maximum Ue is decreased due to the substantially compromised Eb by larger amount of the BTO_nps. As plotted in Figure 11b, despite of the much enhanced D, the maximum Ue of the equilibrium nanocomposites decreases monotonically from ~16.6 J/cm3 to ~7.4 J/cm3 as the volume fraction of BTO_nps increases from 0 to 10 vol.%, due to the much decreased Eb as well as to the substantially raised Dr (see Figure 10a). The discharge efficiency (η), which is defined as η = Ue/U×100% where U represents the total charged energy density, is also plotted as a function of electric field in Figure 11. As seen, η of the non-equilibrium nanocomposites decreases with increasing electric field to ~ 300 MV/m then increases monotonically up to Eb, which is in line with the variation of Dr with electric field (as shown in Figure 10a). Therefore, a rather high η of 74% is obtained at 648 MV/m for the non-equilibrium nanocomposites filled with 1 vol.% BTO_nps. In stark contrast, the equilibrium nanocomposites exhibit consistently lower η than their non-equilibrium counterparts at all electric field. Plus, η decreases monotonically and dramatically with increasing electric field, as shown in Figure 11b.

Effects of BTO_nps on the field-induced phase transformation Two striking features could be well distinguished in the high-field polarization behavior of the non-equilibrium P(VDF-HFP) nanocomposites. First is the suppressed leakage current beyond 300 MV/m (Figure 10b) and the second is the abrupt increase in the discharge efficiency at electric fields > 300 MV/m (Figure 11). It is thus of particular interest to explore the possible origin of the substantial change in the high-field polarization, especially beyond 300 MV/m. It has been speculated that field-induced phase transformation may play critical roles in determining the high-field polarization behavior of PVDF dielectrics. The field-dependent dielectric permittivity

εr(E), which reflects the rate of the change of dipoles with electric field and is closely correlated with phase composition and crystalline orientation of PVDF, has been employed as an indicator for the possible field-induced phase transformation.27,61-64 Although εr(E) may not be directly calculated for non-linear dielectrics due to the non-reversible variation of D with electric field with the following equation as: => =

1 ? = ?

(5)

Nevertheless, (∂D/∂E)/ε0 is qualitatively correlated with εr as a function of electric field. We therefore use (∂D/∂E)/ε0 to qualitatively correlate with the field-dependent dielectric permittivity (εr(E)) at high electric fields.


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ACS Applied Materials & Interfaces (∂D/∂E)/ε0 values are obtained by differentiating D-E loops for both of the two types of polymer nanocomposites and plotted as a function of electric field in Figure 12-13 and Figure S7-S15. Different field-induced phase transformations for the pure equilibrium and non-equilibrium P(VDF-HFP) films are revealed by Figure 12. As shown in Figure 12a for the non-equilibrium pure films, two peaks are observed in the (∂D/∂E)/ε0 curves at 200 MV/mm. Given the fact that both α and β polymorphs are present in the non-equilibrium P(VDF-HFP) films (as summarized in Table 1), the two peaks at ~ 60 MV/m and ~ 120 MV/m during charging process could be associated with the dipolar switching in β and α polymorph respectively, which agrees with their corresponding coercive fields.27 This is also evidenced by the two fast depolarization processes observed in the (∂D/∂E)/ε0 curves in discharging process. At higher poling electric field of ~300 MV/m, the two fast switching peaks are still present but shift to higher electric field, i.e., to ~ 128 MV/m and 230 MV/m respectively. The shifts of the fast switching peaks to higher electric field indicate that the dipoles are stabilized by the space charges hence need higher electric field to switch. This is in line with the substantial increase in leakage current observed for the non-equilibrium nanocomposites (as shown in Figure 10b). Further increase in poling electric field to ~ 400 MV/m leads to not only further shift of the two peaks, i.e., to ~ 188 MV/m and 295 MV/m, but the dramatic change in the relative intensity of the two peaks. For instance, the ratio between the (∂D/∂E)/ε0 values of the β /α polymorph changes from ~ 0.81 at 300 MV/m to ~ 1.12 at 400 MV/m, suggesting that more β polymorph is present in the non-equilibrium P(VDF-HFP) films as a result of the phase transformation from α to β polymorph, or to a polar αp polymorph.65 When the poling electric field increases to 500 MV/m, the two peaks merge into one single peak at ~ 209 MV/m, indicating that the field-induced phase transformation is complete beyond 500 MV/m.66 No further phase transformation could be observed even raising the poling electric field to ~ 600 MV/m (see Figure S8 of supporting information). For the equilibrium P(VDF-HFP) films, two fast switching peaks could also be observed at ~ 78 MV/m and ~ 195 MV/m in Figure 12b. Since only α polymorph is present in the equilibrium P(VDF-HFP) films (as summarized in Table 1), we speculate that these two peaks might be associated with the dipolar switching in α crystalline lamellar and the rest amorphous regions.67 It is of interest to note that the two peaks only shift to higher electric field upon increasing the poling electric field, e.g. ~ 78 MV/m to ~ 109 MV/m, or ~ 195 MV/m to ~ 274 MV/m as the electric field increases from 300 MV/m to 500 MV/m. The ratio between the (∂D/∂E)/ε0 values corresponding to the two peaks remains almost constant at ~ 0.91 with increasing electric field, indicating that no field-induced phase transformation is induced in the equilibrium P(VDF-HFP) films. Further verification of the above field-induced phase transformations is provided by the study of the FTIR patterns of equilibrium and non-equilibrium P(VDF-HFP) films after poling at ~300 MV/m. As shown in Figure 3, for equilibrium films (blue dot line), no obvious changes are detected by FTIR, suggesting that there is no obvious phase transformation after poling up to ~300 MV/m. However, this is not the scenario for non-equilibrium films (pink dot line). After poling, the FTIR peak intensity of theα-phase absorption bands at 763, ACS Paragon Plus Environment


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796, 976 cm-1 decrease obviously, while that of the β-phase absorption bands at 840 and 1279 cm-1 slightly increase. These results indicate that there is a field-induced phase transformation during electric poling process in the non-equilibrium films. The incorporation of BTO_nps in the non-equilibrium nanocomposites facilitates the field-induced phase transformation in two aspects (as shown in Figure 13). First, phase transformation is initiated and complete at lower poling electric field. For instance, at the same poling electric field of ~ 300 MV/m, the ratio between the (∂D/∂E)/ε0 values of the β and α polymorph increases from ~ 0.68 to ~ 0.94 as the volume fraction of BTO_nps increases from 1 vol.% to 5 vol.%. Further increase in the content of BTO_nps to 10 vol.% leads to complete phase transformation at 300 MV/m as evidenced by the single peak at ~ 166 MV/m in Figure 13c. This is compared with a high electric field of ~ 500 MV/m needed for complete phase transformation in the pure non-equilibrium P(VDF-HFP) films. Second, the fast switching peak of β polymorph shifts to lower electric field as more BTO_nps are introduced. At the same poling electric field of ~ 400 MV/m, the peak corresponding to the fast switching in

β polymorph shifts monotonically from ~ 211 MV/m to ~ 164 MV/m with increasing BTO_nps loading from 1 vol.% to 10 vol.%. As for the equilibrium nanocomposites, BTO_nps exhibit no effects on the field-induced phase transformation (as shown in Figure S8 of supporting information). We speculate that in the poling process, the BTO_nps are polarized and saturated at low electric field as result of their high susceptibility. The resultant strong depolarization field inside the BTO_nps distorts the homogeneous distribution of local electric field and induces substantial concentration of local electric field at the interface between BTO_nps and P(VDF-HFP).68 P(VDF-HFP) matrix is thus subjected to a much higher local electric field despite of the low overall electric field. Plus, the charged surfaces of the polarized BTO_nps are in favor for the nucleation of β polymorph,69,70 which also facilitates the field-induced phase transformation.

Conclusions Polymer nanocomposites with high reliability, high energy density has been prepared by a non-equilibrium processing combing electrospinning, hot-pressing and thermal quenching. During this process, the rapid evaporation of solvent, layer-by-layer assembly and high-ratio stretching ensure homogeneous dispersion of BTO_nps in P(VDF-HFP) matrix and high quality of the nanocomposites. An ultrahigh Weibull modulus β of ~ 30 is achieved in the non-equilibrium nanocomposites. The concomitant enhancement of both D and Eb leads to an ultrahigh energy density of 20.6 J/cm3. The non-equilibrium processing also induces mixed α and β phases and small crystalline size of P(VDF-HFP), which leads to suppressed leakage current at high electric field hence high discharge efficiency of the nanocomposites. Transformation from α to β polymorph is induced by high electric field in the non-equilibrium nanocomposites while no field-induced phase transformation is observed for the nanocomposites fabricated by an equilibrium solution-casting. The incorporation of BTO_nps facilitates the phase ACS Paragon Plus Environment

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ACS Applied Materials & Interfaces transformation at low electric field due to the concentrated local electric field at the BTO_nps/P(VDF-HFP) interfaces and the charged surface of BTO_nps that is in favor for the nucleation of β polymorph.

Supporting Information. SEM images and XRD patterns of BTO_nps; Cross-sectional SEM images, FTIR & WAXD patterns of both equilibrium and non-equilibrium nanocomposites with different volume fractions of BTO_nps; Schematic D-E loop for a dielectric materials; Derivative curves of D, (∂D/∂E)/ε0, as a function of electric field under different poling electric field for both equilibrium and non-equilibrium nanocomposites.

Acknowledgement This work was supported by the NSF of China (Grant No. 51625202, 51572141 and 51532002), the National Basic Research Program of China (Grant No. 2015CB654603) and Research Fund of Science and Technology in Shenzhen (JSGG20150331155519130).

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ACS Applied Materials & Interfaces Nanocomposite Capacitors. Adv. Energy Mater. 2013, 3, 451-456.

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Figures and Tables

Scheme 1. Schematic illustration of fabrications of BTO/P(VDF-HFP) nanocomposite films by (a) equilibrium process (or solution casting process) and (b) non-equilibrium process.



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Figure 1. SEM images of BTO_nps@P(VDF-HFP) composite fibers (a,d,g,j), top surfaces (b,e,h,k) and cross sections (c,f,i,l) of BTO_nps/P(VDF-HFP) nanocomposite films by non-equilibrium process with different BTO_nps volume fractions: 0 vol.% (a,b,c), 1 vol.% (d,e,f), 5 vol.% (g,h,i), 10 vol.% (j,k,l).


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Figure 2. SEM images of top (a) and bottom (b) surfaces and cross sections (c, d) of 5 vol.% BTO_nps/P(VDF-HFP) nanocomposites prepared by equilibrium process.



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Figure 3. FTIR patterns of pure non-equilibrium and equilibrium P(VDF-HFP) films before (solid line) and after (dot line) poling, respectively.


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Figure 4. WAXD patterns of pure equilibrium (a) and non-equilibrium (b) P(VDF-HFP) films, respectively.



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Figure 5. Melting and crystallization DSC curves of non-equilibrium and equilibrium nanocomposite films. The inset presents the corresponding crystallinity of nanocomposites with different volume fractions.

Figure 6. Frequency dependence of (a) dielectric constant and (b) dielectric loss of non-equilibrium and equilibrium nanocomposite films. The inset in (a) shows the corresponding dielectric constant of nanocomposites with different volume fractions measured at 1 kHz.


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Figure 7. Weibull distributions of electric breakdown strength at room temperature for non-equilibrium nanocomposite films (a) and equilibrium nanocomposite films (b).



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Figure 8. Breakdown strength and β-parameters for non-equilibrium and equilibrium nanocomposite films loaded with various concentrations of BTO_nps.

Figure 9. Unipolar electric displacement-electric field (D-E) loops for non-equilibrium films (a) and equilibrium films (b) with different BTO_nps volume fractions.


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Figures 10. (a) Remnant displacement of the equilibrium (top) and non-equilibrium (bottom) nanocomposite films with different volume fractions as a function of the electric field summarized from D-E loops in Figure 9a,b, respectively. (b) Variations of leakage current with electric field for pure equilibrium and non-equilibrium P(VDF-HFP) films, respectively.

Figure 11. Discharged energy density and efficiencies of the non-equilibrium films (a) and equilibrium films (b) with different volume fractions as a function of the electric field calculated from D-E loops in Figure 9a,b, respectively.



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Figure 12. Derivative curves of D, (∂D/∂E)/ε0, as a function of electric field under different poling electric field (200-500 MV/m) for pure non-equilibrium films (a) and pure equilibrium films (b), respectively.


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Figure 13. Derivative curves of D, (∂D/∂E)/ε0, as a function of electric field under different poling electric field

(300-500 MV/m) for non-equilibrium nanocomposite films with 1vol.% (a), 5 vol.% (b) and 10 vol.% (c) BTO_nps loading, respectively.



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Table 1. α and β contents and average crystalline sizes for α- and β-phase in pure P(VDF-HFP) films fabricated with two methods. Samples

α6wt. %7D

β6wt. %7D

67F 6'G7;

67F 6'G7;

6/7I 6'G7;

Crystallinity 6wt. %7+

Pure EQ film

43.4

—

30.6

32.4

—

42.6

Pure NEQ film

29.1

21.5

12.0

15.3

14.1

44.5

D Determined by wide-angle X-ray diffraction (WAXD). ; Determined by wide-angle X-ray diffraction (WAXD) and calculated from

the Scherrer equation: KLM = 6NO7/6PKLM %&(Q7, where KLM is the mean crystallite size along the [ℎTU] direction, N is the shape factor, O is the wavelength, and PKLM is the full width at half-maximum for the 6ℎTU7 direction. + Determined by differential

scanning calorimetry (DSC).


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