Enhanced ferroelectricity and energy storage in poly(vinylidene fluoride)–clay nanocomposite films via nanofiller surface charge modulation

Abstract

Ferroelectric polymer poly(vinylidene-fluoride) (PVDF) and its copolymers films have long been considered as the most promising candidates for non-volatile organic electronic devices. Fabrication of high-quality PVDF film with a high breakdown electric field and excellent ferroelectricity is required. Designing nanocomposites with a combination of the advantages of both PVDF and nanofillers provides a feasible route to obtain high-performance ferroelectric polymer films. Instead of the usually selected high-k dielectric ceramic nanoparticles, we chose nanolaminate shaped clay as the nanofiller. We used surfactant modification to tune the surface charge and the dispersion of the filler simultaneously so that enhanced ferroelectricity and energy storage in PVDF/clay nanocomposite film at rather low clay loading were realized. Compared to the pristine PVDF film, the energy density was increased from 5.34 J cm⁻³ to 5.91 J cm⁻³ at 1 wt% MMT content and the “maximum” energy density could reach 10.2 J cm⁻³. Our results demonstrate a low-cost and facile method to tune the electrical properties of PVDF film so that it could be easily integrated into all-organic electronic devices.

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Introduction

Polymer based dielectric materials are presently the choice for energy storage applications because of their superiorities of high energy density, high electric breakdown field (E_BD), low dielectric loss, fast speed, low cost, good processability and high reliability.^1–4 The discharged energy density of a dielectric material is defined as the integral , where E is the electric field and D is the electric displacement. According to the equation, the energy density of a dielectric material is determined by both the E_BD and saturated electric displacement. For the linear dielectrics, the maximum energy density U_e = 1/2ε₀ε_rE_BD², where ε₀ is the vacuum permittivity (ε₀ = 8.85 × 10⁻¹² F m⁻¹) and ε_r is the permittivity of the dielectric material. However, the commercial dielectric polymers (e.g. biaxially oriented polypropylene shows the highest energy density of ∼1–2 J cm⁻³) used for high energy density capacitors have generally low ε_r (i.e., less than 3), so the high energy density originates from the high E_BD (>500 MV m⁻¹). Ferroelectric poly(vinylidene-fluoride) (PVDF) and its copolymer with trifluoroethylene (TrFE) have been extensively investigated as ideal candidates since they exhibit large ε_r (∼10–50), bistable remanent polarization that can repeatedly be switched by an electric field and facile solution-processing thin films fabrication. Zhang et al. have shown that, by combining the reversible nonpolar and polar molecular structural changes to realize high D with matched dielectric constant values to avoid the early D-saturation, a very high energy density (>17 J cm⁻³) with fast discharge speed (<1 μs) and low dielectric loss can be obtained in defect-modified P(VDF-chlorotrifluoroethylene) polymers.¹

Pure PVDF is a semi-crystalline polymer with at least four different polymorphs referred to as the α, β, γ and δ-phase. The α-phase is non-polar and paraelectric because of the centro-symmetric symmetry of the unit cell. The β-PVDF is ferroelectric due to macroscopic dipole moment which is the most desired phase. The polymer chains in γ-PVDF have a conformation in between that of α and β-phases making it ferroelectric but hard to acquire experimentally. The δ-phase is a polar version of the α-phase and has recently been proved to show good ferroelectricity comparable with the copolymer P(VDF-TrFE).⁵ Although commercial PVDF films are available, fabrication of high-quality ferroelectric PVDF thin films as required for microelectronic applications is a long-standing problem.

An alternative way to obtain high energy density of ferroelectric polymer based materials is designing nanocomposites. The potential combination of increased E_BD and ε_r, along with gradual failure modes, material tunability and easy processability, provides an attractive approach. The electroactive β-phase PVDF has been obtained in a wide variety of nanocomposites, including the use of BaTiO₃,⁶ TiO₂,⁷ (BaSr) TiO₃,^8,9 ferrite nanoparticles,¹⁰ clay,^11–13 PMMA¹⁴ and ionic-liquid.¹⁵ Zhang X. et al. have recently reported an ultrahigh energy density of 20 J cm⁻³ of PVDF/BaTiO₃@TiO₂ nanofibers nanocomposites.¹⁶ The presence of various nanostructured fillers in the polymer matrix leads to significant modifications of the breakdown field, charge transport, and charge distribution of the dielectric materials brought by the multi-scale interfacial effects.^17,18 The literature has already concluded that nanofillers with high aspect ratio and their spatial alignment in the polymer matrix are beneficial for large ε_r and E_BD.^9,11,19 Moreover, as pointed out by Vaia R. A. et al.,²⁰ for the dielectric polymer nanocomposite design, it is insufficient to control just the nanoparticle morphology and surface, the ability to tune the relative polarity of the nanoparticle surface with regard to the matrix while maintaining dispersion should be paid attention to.

From this point of view, we used clay montmorillonite (MMT) as the filler, which could be exfoliated easily into nanolaminates acting as barriers capable of retarding breakdown and carrier mobility instead of the usually chosen high-k dielectrics (e.g., BaTiO₃). Meanwhile, cetrimonium bromide (CTAB) was chosen as the surfactant for two distinct aspects: (i) the long alkyl chain of CTAB could be intercalated into the interlayers of MMT so as to improve the dispersity of the nanolaminates. (ii) CTAB provides positive surface charge which can tune the polarization of negatively charged MMT nanolaminates. The dielectric and ferroelectric behaviors of PVDF/MMT nanocomposite films were investigated. Enhanced ferroelectric behavior and energy density were realized in PVDF/MMT nanocomposite film with rather small loading 1 wt% of MMT.

Experimental

PVDF/MMT nanocomposite film preparation

The clay used in this study is Na⁺-type montmorillonite supplied by Zhejiang Fenghong Clay Ltd. Co. The raw MMT was first dispersed in water and stirred for one week, then centrifuged at 3000 rpm for 10 min to remove the MMT that was not exfoliated. After centrifugation, the supernatant became transparent, and the clay nanolaminates were kept suspended without precipitation for one week. Next, exfoliated MMT nanolaminates were modified by CTAB in the following procedures. CTAB was added to the supernatant and stirred in water bath at 80 °C for 8 h. After filtration and washed with the mixture of ethanol and deionized water (mole ratio 1 : 1) several times to remove the unreacted CTAB, then the product was dried in vacuum oven at 80 °C for 10 h, and modified MMT was obtained.

Composite films were fabricated via a solution casting method. First, modified MMT nanolaminates were dispersed in N,N-dimethylformamide (DMF, pure grade) ultrasonicated and stirred for one week to form a stable suspension. Then PVDF powders (Shanghai 3F Co., China) were dissolved in the suspension and stirred vigorously for 2 h to make a homogeneous solution. Films were fabricated on an Electric coating machine (Eleometer4340). The solution was casted onto glass sheets which were laid on a heating panel at 80 °C. After the solvent was thoroughly evaporated, the panel was cooled down gradually to room temperature in the air, and the dried films could be easily peeled off the glass. The final thickness of the films varied from 10 to 15 μm. Scanning electron microscopy (SEM, Quanta 250 FEG) measurements were performed. Fourier transformer infrared (FTIR) spectra were collected from Nicolet In 10MX infrared spectrophotometer. Zeta potential measurements were carried out to determine the surface charge of MMT nanolaminates before and after CTAB modification on Zetasizer (NanoZS90). Measurements were performed in pure water and an average value for each sample was obtained from 3 measurements. Differential scanning calorimetry (DSC, Netzsch STA 449C) was used to determine the crystalline percentage of the composite films with the measurement temperature range from room temperature to 200 °C under a nitrogen atmosphere at a heating rate of 10 °C min⁻¹.

Dielectric and electrical characterization

For electrical measurements, Au electrodes (4 mm in diameter) were sputtered on both sides of the films using a mask. Dielectric constant and loss were measured on an Agilent 4294A impedance analyzer in a frequency range 10² to 10⁷ Hz at an ac driving voltage of 1 V. The breakdown strength measurements were carried out on a CS2671A dielectric strength tester (Nanjing Changsheng Instrument Ltd. Co.). The D–E loops were measured at 10 Hz on a Premier II ferroelectric test system (Radiant Technologies Inc.). For each electrical characterization, several measurements were collected for each sample to ensure the repeatability of the data.

Results and discussion

To clarify the effects of interface modification on electrical properties of the composites, how the CTAB modified the MMT nanolaminates was first examined. FTIR spectra clearly indicate that the alkyl groups from CTAB have been added onto MMT nanolaminates (not shown here). Furthermore, the surface charge of MMT nanolaminates before and after CTAB modification was changed from negative −37 mV to positive 7.92 mV.

The literature already suggested that the key factor of the PVDF β-phase nucleation is the electric interaction between the bonds of the PVDF chains and the electric charged surface of nanofillers.²¹ Previous studies showed that highly polar solvents can induce β or γ phases in PVDF films, and the fraction of the polar phase strongly depends on the solvent evaporation rate.²² High quality PVDF films in polar phase are still hard to prepare as discussed in recent review article.²³ Due to the similarity in polymer chain conformation of the β and γ phases, it usually needs combination of different characterization methods to quantitatively identify the fraction of each phase.

XRD patterns of pure PVDF film, CTAB modified MMT and PVDF/MMT composite films were shown in Fig. 1(a). All the diffractions show low intensity and broadened peaks, indicating generally low crystalline percentage of the films, probably due to the low crystallization temperature (i.e. 80 °C). As calculated from the DSC curves (shown in Fig. 1(b)), degree of crystallinity of all the composite films is less than 10%. Nevertheless, the degree of crystallinity increased from 6.88% of pristine PVDF film to 8.79% by addition of 1 wt% MMT nanolaminates, 7.45% for 3 wt% MMT and 8.12% for 10 wt% MMT. The melting temperature range obtained from DSC curves for all the films is around 168–170 °C, corresponding to the melting temperature for both α and β phases. Seen from XRD patterns, the characteristic diffractions from α and β phases were observed in pure PVDF film. As MMT fraction increased, the fraction of β phase gradually increased, while diffractions from α phase almost disappeared as MMT content exceeds 3 wt%. These results indicate that MMT nanofillers could promote the crystallization of PVDF films especially the nucleation of β-PVDF.

Fig. 1 (a) X-ray diffraction patterns and (b) DSC curves of PVDF/MMT composite films with various MMT fractions.

As shown in FTIR spectra (Fig. 2), the distinguishable absorption band of γ phase at 1234 cm⁻¹ also appeared in pristine PVDF film, which suggests that the film is composed of α, β and γ phases. The FTIR spectra showed that all the PVDF/MMT composite films contain β and γ phases. The characteristic peak of α-PVDF at 765 cm⁻¹ gradually shifted to higher frequency and the intensity became weakened as MMT loading increased, which further verify that MMT nanolaminates suppress the crystallization of α phase and facilitate the formation of β phase.

Fig. 2 FTIR spectra of PVDF/MMT composite films with MMT nanolaminates loadings 0, 1, 3, 5, and 10 wt%.

Seen from the surface topography images of the PVDF/MMT composite films shown in Fig. 3, the size of spherulites formed in pristine PVDF film varies from 2 to 3 μm. As MMT loading increases from 1 to 10 wt%, the size of the spherulite decreases and surface fluctuation becomes even. The SEM images again demonstrate that MMT nanolaminates act as nucleating agents, since more nucleation sites greatly reduced the grain size. Meanwhile, due to the decrease in spherulite size, the pinholes appeared in pristine PVDF film [see Fig. 3(a)] vanish and the surface becomes smooth as MMT nanolaminates were loaded.

Fig. 3 SEM topography images of MMT/PVDF composite films with various MMT fractions: (a) pure PVDF film, (b) 1 wt% MMT, (c) 3 wt% MMT, (d) 5 wt% MMT, and (e) 10 wt% MMT.

Fig. 4(a) presents the frequency dependence of the dielectric properties in the range 10² to 10⁷ Hz at room temperature. The composites generally follow the trend of pristine PVDF with a slight increase in ε at low frequency, i.e., below 2 kHz, which comes from the contribution of the surface charges of MMT nanolaminates. At high frequency range, the frequency dependence of dielectric constants of the composites mainly exhibit the characteristic of PVDF matrix, except for the 10 wt% MMT film, in which the high content of MMT nanolaminates begins to impede the polarization of PVDF matrix. With the increase of MMT nanolaminates loading, the ε increased and reached maximum value at 3 wt% MMT loading as shown in Fig. 4(b). Continuing increase of the MMT loading would not increase the ε but cause the increase of the dielectric loss as shown in the inset of Fig. 4(b), which is due to the partial agglomeration of the MMT nanolaminates at high loading as shown in Fig. 3(e) the surface image of 10 wt% MMT film. Though the loss of the composite film is increased by MMT nanolaminates, it is generally kept at a low level, i.e., less than 0.05 within 5 wt% MMT loading, indicating good dispersion of MMT nanofillers and combination between nanofillers and PVDF matrix.

Fig. 4 Dielectric behaviors of PVDF/MMT nanocomposite films with various MMT mass fractions. (a) Dependence of dielectric constants on the frequency at room temperature with the inset showing the dependence of dielectric loss on the frequency. (b) Dielectric constant of PVDF/MMT nanocomposite film as a function of mass fraction of MMT, measured at 1 kHz. Inset shows the dielectric loss of PVDF/MMT nanocomposite film as a function of mass fraction of MMT.

Fig. 5 summarizes the analysis of PVDF/MMT composites failure utilizing a two parameter Weibull cumulative probability function: P(E) = 1 − exp[−(E/E_BD)^β], where P(E) is the cumulative probability for electric failure and E is experimentally recorded breakdown strength. The fitting parameters are E_BD, the breakdown field where there is 63.2% probability for failure, and β, the shape parameter associated with the least-squares linear fit of the distribution. Seen from the inset of Fig. 5, the change of E_BD as a function of MMT mass fraction, the E_BD showed great increase from 380.2 kV mm⁻¹ of the pristine PVDF film to 463.4 kV mm⁻¹ by adding only 1 wt% MMT nanolaminates and reached maximum value of 480.8 kV mm⁻¹ at 3 wt% MMT nanolaminates. The shape parameter β associated with the least-squares linear fit of the distribution describes the scattering in the experimental data, and higher β represents the homogenous microstructure of the sample. The remarkable increase in E_BD (by 26.3% compared to pristine PVDF film) of MMT 3 wt% film comes from the homogenous distribution of MMT nanolaminates microstructure and the improvement of the film quality, which is in consistent with the low loss of the composite film. And moreover, the increase in E_BD indicates that MMT nanolaminates are effective in retarding the propagation of breakdown.

Fig. 5 Experimentally observed dielectric breakdown strengths, plotted as a probability of failure (P) in the form of lg(−ln(1 − P)) vs. electric field (Weibull statistics). The inset shows breakdown strength as a function of MMT mass fraction, determined at 63.2% failure probability from a linear regressive fit of the two-parameter Weibull failure analysis.

Since the MMT nanolaminates facilitate the crystallization of β phase PVDF, the ferroelectric behavior of PVDF/MMT composite films was examined as shown in Fig. 6. For the composite films with low nanofiller loadings, i.e., 1 and 3 wt% [see Fig. 6(b) and (c)], enhanced remnant polarizations (P_r) (P_r = 6.5 μC cm⁻² for MMT 1 wt% and 6.3 μC cm⁻² for MMT 3 wt%) and well-saturated D–E loops were obtained compared to pristine PVDF film with P_r = 5.0 μC cm⁻². Composites with MMT loading higher than 5 wt% cannot have saturated hysteresis D–E loops due to the increased leakage current and early breakdown before reaching saturation as presented in Fig. 6(d) and (e). The discharged energy density was calculated from the measured D–E loop according to the integral and the change of energy density with MMT mass fraction is shown in Fig. 6(f). Compared to the pristine PVDF film, incorporation of a small amount of MMT nanolaminates enhances the energy density from 5.34 J cm⁻³ of PVDF film to 5.91 J cm⁻³ for MMT 1 wt% and 5.5 J cm⁻³ for MMT 3 wt% film, respectively.

Fig. 6 Displaced charge density of PVDF/MMT composite films presented as a function of applied electric field measured at 10 Hz. (a) Pure PVDF film, (b) 1 wt% MMT, (c) 3 wt% MMT, (d) 5 wt% MMT, and (e) 10 wt% MMT. (f) The discharged energy density of PVDF/MMT composite film as a function of MMT mass fraction.

We further analyze the change of energy density as a function of electric field. Since the key factors in achieving a high energy density include a high E_BD, a high D value and a proper ε to avoid the electric displacement saturation at fields well below E_BD.¹ Fig. 7 showed the calculated discharged energy densities versus electric field for the pristine PVDF film and composite films with various MMT loadings. The fitted curves in dash lines were used to estimate the “maximum” energy density (ED_max) at the electric breakdown strength of each film. As seen, the ED_max of the PVDF/MMT with only 1 wt% loading could reach as large as 10.2 J cm⁻³.

Fig. 7 Calculated discharged energy densities according to the D–E loops versus electric field for the pristine PVDF film and composite films with MMT loadings 1, 3, 5 and 10 wt%. The vertical arrows denote the dc electric breakdown strength (E_BD) of each material and the horizontal arrows denote the corresponding “maximum” energy density (ED_max).

Therefore, the promotion of polar β-phase crystallization in PVDF matrix by MMT nanolaminates, the modification of the surface polarization of MMT nanolaminates with regard to the PVDF matrix, the good dispersion of the nanolaminates in the PVDF matrix, as well as the compatibility between the MMT nanolaminates and PVDF polymer come to the best effect in enhancing the ferroelectricity and energy density of the PVDF/MMT composite film with addition of 1 wt% CTAB-modified MMT nanolaminates. The results demonstrate a facile and low-cost route to obtain organic ferroelectric materials with high-performance in energy storage and ferroelectricity demanded by all-organic electronic devices.

Conclusions

In summary, PVDF/clay nanocomposite film was designed based on the thought to tune the relative polarity of the nanoparticle surface yet maintaining its dispersion with regard to the matrix. Surfactant CTAB was utilized to tune the surface charge as well as to improve the compatibility between polymer matrix and inorganic MMT nanolaminate fillers. MMT nanolaminates promote the nucleation of β phase PVDF and the homogeneous dispersion of MMT nanolaminates retards the breakdown propagation. Therefore, enhanced ferroelectricity and energy storage in PVDF/MMT nanocomposite film at rather low clay loading were realized. Our results demonstrate a low-cost and facile method to tune the electrical properties of PVDF film so that it could be easily integrated into all-organic electronic devices.

Acknowledgements

The work was supported by the State Key Development Program for Basic Research of China (Grant No. 2012CB933200), National Natural Science Foundation of China (No. 51172008 and 51002006), National Natural Science Fund Innovation Group (No. 51221163), and the Fundamental Research Funds for the Central Universities.

References

1. B. J. Chu, X. Zhou, K. L. Ren, B. Neese, M. R. Lin, Q. Wang, F. Bauer and Q. M. Zhang, Science, 2006, 313, 334.
2. Q. Wang and L. Zhu, J. Polym. Sci., Polym. Phys. Ed., 2011, 49, 1421.
3. Z. M. Dang, J. K. Yuan, S. H. Yao and R. J. Liao, Adv. Mater., 2013, 25, 6334.
4. S. J. Kang, I. Bae, J. H. Choi, Y. J. Park, P. S. Jo, Y. Kim, K. J. Kim, J. M. Myoung, E. Kim and C. Park, J. Mater. Chem., 2011, 21, 3619.
5. M. Y. Li, H. J. Wondergem, M. J. Spijkman, K. Asadi, I. Katsouras, P. W. M. Blom and D. M. de Leeuw, Nat. Mater., 2013, 12, 433–438.
6. Z. M. Dang, S. H. Yao, J. K. Yuan and J. B. Bai, J. Phys. Chem. C, 2010, 114, 13204.
7. J. J. Li, S. I. Seok, B. J. Chu, F. Dogan, Q. M. Zhang and Q. Wang, Adv. Mater., 2009, 21, 217.
8. S. H. Liu, J. W. Zhai, J. W. Wang, S. X. Xue and W. Q. Zhang, ACS Appl. Mater. Interfaces, 2014, 6, 1533.
9. H. X. Tang and H. A. Sodano, Nano Lett., 2013, 13, 1373.
10. P. Martins, C. Caparros, R. Goncalves, P. M. Martins, M. Benelmekki, G. Botelho and S. Lanceros-Mendez, J. Phys. Chem. C, 2012, 116, 15790.
11. S. P. Fillery, H. Koerner, L. Drummy, E. Dunkerley, M. F. Durstock, D. F. Schmidt and R. A. Vaia, ACS Appl. Mater. Interfaces, 2012, 4, 1388.
12. V. Tomer, E. Manias and C. A. Randall, J. Appl. Phys., 2011, 110(4), 044107.
13. V. Tomer, G. Polizos, C. A. Randall and E. Manias, J. Appl. Phys., 2011, 109(7), 074113.
14. M. Y. Li, N. Stingelin, J. J. Michels, M. J. Spijkman, K. Asadi, K. Feldman, P. W. M. Blom and D. M. de Leeuw, Macromolecules, 2012, 45, 7477.
15. F. P. Wang, A. Lack, Z. L. Xie, P. Frubing, A. Taubert and R. Gerhard, Appl. Phys. Lett., 2012, 100(6), 062903.
16. X. Zhang, Y. Shen, Q. Zhang, L. Gu, Y. Hu, J. Du, Y. Lin and C.-W. Nan, Adv. Mater., 2015, 27, 819–824.
17. L. Y. Xie, X. Y. Huang, Y. H. Huang, K. Yang and P. K. Jiang, J. Phys. Chem. C, 2013, 117, 22525.
18. K. Yu, Y. J. Niu, Y. Y. Bai, Y. C. Zhou and H. Wang, Appl. Phys. Lett., 2013, 102(10), 102903.
19. C. W. Nan, Y. Shen and J. Ma, Annu. Rev. Mater. Res., 2010, 40, 131.
20. C. A. Grabowski, S. P. Fillery, N. M. Westing, C. Z. Chi, J. S. Meth, M. F. Durstock and R. A. Vaia, ACS Appl. Mater. Interfaces, 2013, 5, 5486.
21. G. J. Zhong, L. F. Zhang, R. Su, K. Wang, H. Fong and L. Zhu, Polymer, 2011, 52, 2228.
22. D. L. Chinaglia, R. Gregorio, J. C. Stefanello, R. A. P. Altafim, W. Wirges, F. P. Wang and R. Gerhard, J. Appl. Polym. Sci., 2010, 116, 785.
23. P. Martinsa, A. C. Lopesa and S. Lanceros-Mendez, Prog. Polym. Sci., 2014, 39, 683.

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