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Ultrahigh discharge efficiency in multilayered polymer nanocomposites of high energy density
Author’s Accepted Manuscript Ultrahigh Discharge Efficiency in Multilayered Polymer Nanocomposites of High Energy Density Jianyong Jiang, Zhonghui Shen, Jianfeng Qian, Zhenkang Dan, Mengfan Guo, Yuanhua Lin, CeWen Nan, Longqing Chen, Yang Shen www.elsevier.com/locate/ensm
PII: DOI: Reference:
S2405-8297(18)30838-9 https://doi.org/10.1016/j.ensm.2018.09.013 ENSM506
To appear in: Energy Storage Materials Received date: 6 July 2018 Revised date: 14 September 2018 Accepted date: 14 September 2018 Cite this article as: Jianyong Jiang, Zhonghui Shen, Jianfeng Qian, Zhenkang Dan, Mengfan Guo, Yuanhua Lin, Ce-Wen Nan, Longqing Chen and Yang Shen, Ultrahigh Discharge Efficiency in Multilayered Polymer Nanocomposites of High Energy Density, Energy Storage Materials, https://doi.org/10.1016/j.ensm.2018.09.013 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Ultrahigh
Discharge
Efficiency
in
Multilayered
Polymer Nanocomposites of High Energy Density Jianyong Jianga,d, Zhonghui Shena, Jianfeng Qiana, Zhenkang Dana, Mengfan Guoa, Yuanhua Lina, Ce-Wen Nana, Longqing Chena,b, Yang Shena,c* a
State Key Lab of New Ceramics and Fine Processing, School of Materials Science and
Engineering, Tsinghua University, Beijing, 100084, China b
Department of Materials Science and Engineering, The Pennsylvania State University,
University Park, Pennsylvania 16802, United States c
Center for Flexible Electronics Technology, Tsinghua University, Beijing 100084, China
d
Laboratory of Advanced Energy Storage Materials & Devices, Research Institute of
Tsinghua University in Shenzhen, Shenzhen, 518057, China
*
Corresponding author. [email protected]
Abstract
Poly(vinylidene fluoride) (PVDF)-based dielectric polymers are in great demand for the future electronic and electrical industry because of their high dielectric constants and energy density. However, some issues that limit their practical applications remain unsolved. One of the most urgent issues is their high dielectric loss and hence low efficiency. In this contribution, we proposed and demonstrate that substantially enhanced discharge efficiency of PVDF-based polymers nanocomposites could be achieved by simultaneously optimizing 1
their
topological-structure
and
fluoride-co-hexafluoropropylene)
phase
composition.
(P(VDF-HFP))
In
the /
poly(vinylidene poly(vinylidene
fluoride-ter-trifluoroethylene-ter-chlorofluoroethylene) (P(VDF-TrFE-CFE)) multilayered nanocomposites fabricated by non-equilibrium process, an ultrahigh discharge efficiency of ~85% is achieved up to 600 MV/m, which is the highest discharge efficiency reported so far for any polar-polymer dielectric materials at such high electric field. By adjusting the quenching temperature, the phase-composition hence dielectric permittivity in the terpolymer layers could be tuned for suppressed ferroelectric loss. Results of phase-field simulations further reveal that local electric field is substantially weakened at the interfaces between the Co/Ter polymer layers, which will act as barriers to motion of charge carriers and give rise to much suppressed conduction loss and a remarkably enhanced breakdown strength. Synergy of the optimized topological-structure and phase-composition thus leads to a nanocomposite that exhibits an unprecedented high discharge efficiency of the multilayered nanocomposites that is comparable to the bench-mark biaxially oriented polypropylene (BOPP) at high electric field as well as a high discharge energy density that is over 10 times higher than that of BOPP.
KEYWORDS: PVDF; charge-discharge efficiency; energy density; multilayer structure; phase tuning; interface.
2
Graphical Abstract
Synergy of the optimized topological-structure and phase-composition is employed to suppress the conduction loss and ferroelectric loss in P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites. An unprecedented high discharge efficiency of ~85% is achieved up to ~600 MV/m, which is by far the highest discharge efficiency ever achieved in polar-polymer dielectric materials at such high electric field. More importantly, a high energy density of ~20 J/cm3 is also achieved.
1. Introduction Dielectric materials with high power density are of critical significance for pulsed power systems, smart grid and electric vehicles.[1-3] Compared to other electrochemical energy storage devices such as supercapacitors and lithium-ion batteries, the electrostatic capacitors based
on
dielectric
media
store
and
release
electrical
energy
through
the
polarization/depolarization process of dipole in dielectric materials.[4] Among all available dielectric materials, polymers are of considerable interest because of their advantages such as 3
high reliability, low cost, lightweight and ease of fabrication.[5-7] Albeit their higher power density, dielectric polymers exhibit rather low energy density and thus large volume and weight for applications.[8] For instance, in a combat vehicle, the pulse power loads, including inverters, converters, power distribution and PFN (Pulse Forming Network), etc., can occupy ~44% of power system volume.[9] In general, the discharged energy density Ue of linear dielectric materials is given by Ue=1/20rEb2, where 0 is the vacuum permittivity, r and Eb are the effective permittivity and the breakdown strength of the dielectrics, respectively.[10] Most of the dielectric polymers possess comparatively high Eb while rather low r. For instance, biaxially oriented polypropylene (BOPP), which is the bench-mark dielectric polymer of current use, only exhibits an r of 2.2 (1 kHz), hence a low discharge energy density of ~ 2 J/cm3 despite their high Eb of ~ 700 MV/m.[11] Recently, Poly(vinylidene fluoride) PVDF-based dielectric polymers with high dielectric permittivity (10~50), stemming from high polarization of the C-F bonds and spontaneous oriention of the dipoles in the crystalline domains, are drawing increasing attention for applications in high energy density
capacitors.
PVDF-based
copolymers
such
as
poly(vinylidene
fluoride-co-hexafluoropropylene), P(VDF-HFP), deliver discharge energy density of 13.5 J/cm3 and efficiency of 62% with an Eb of 600 MV/m.[12] Another class of PVDF-based polymers
is
the
terpolymers
including
poly(vinylidene
fluoride-ter-trifluoroethylene-ter-chlorofluoroethylene), P(VDF-TrFE-CFE), poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene), poly(vinylidene
P(VDF-TrFE-CTFE),
fluoride-ter-trifluoroethylene-ter-hexafluoropropylene), 4
P(VDF-TrFE-HFP),[13] which enjoy the highest r (~50 at 1 kHz) among all the dielectric polymers explored so far. Relaxor ferroelectric behaviors are observed for the terpolymers, which endow the terpolymers with a high Ue at low electric field. For example, a high Ue ~10 J/cm3 is achieved at an electric field of ~350 MV/m in P(VDF-TrFE-CFE).[14] To further increase r and Ue, nanosized ceramic fillers with high r such as BaTiO3, Ba1-xSrxTiO3, have been incorporated into PVDF-based polymers to form polymer nanocomposites.[15, 16] We recently showed that by incorporating TiO2 nanofibers embedded with BaTiO3 nanoparticles (BTO@TO_nfs) into P(VDF-HFP) matrix, hierarchical interfaces are introduced and give rise to concomitant enhancement of r and Eb, leading to a giant discharge energy density of ~31.2 J/cm3 and efficiency of 78% at 797.7 MV/m.[17] Albeit high energy density realized in PVDF-based dielectric polymers and their nanocomposites, the discharge efficiency of these dielectric materials is still low ( of ~ 50%-70%) at high electric fields.[1, 18] It means that more than 1/5 of stored electrical energy is not recovered and dissipated in the form of Joule heat that accumulates within the polymer and may cause early failure of these dielectric materials at an electric field much lower than the intrinsic Eb of the polymer matrix. The unreleased energy is also called the dissipated energy or dielectric loss. There are mainly two types of dielectric loss: ferroelectric loss and conduction loss.[19-22] The ferroelectric loss is induced during the ferroelectric switching when the dipolar switching does not synchronize with the applied alternating electric field. As a semi-crystalline ferroelectric polymer, the ferroelectric loss of PVDF-based polymers is closely related to their phase structures and crystalline size. 5
Specifically, compared to polar phase of PVDF with the all-trans configuration hence large ferroelectric loss, the non-polar phase with zig-zag configuration exhibits smaller ferroelectric loss.[23] Terpolymers, such as P(VDF-TrFE-CFE), with small crystalline size due to the nano-confinement effect of large CFE group, exhibit relaxor ferroelectric behavior where the dipoles rotate and respond rapidly to the external electrical field leading to a low ferroelectric loss.[24] Different from the prominent ferroelectric loss that is intrinsic to the ferroelectric polymers, conduction loss is more common to all dielectric polymers. Biaxially-oriented Polypropylene (BOPP), a paraelectric polymer without permanent dipolar moments along the chains exhibits almost no ferroelectric loss at low electric fields but a rather high conduction loss of ~15% at ~ 600 MV/m due to much increased leakage current at high electric fields.[20, 25] Hence, in order to achieve a high discharge efficiency, both ferroelectric and conduction losses should be suppressed concurrently in PVDF-based dielectric polymers. In this contribution, we propose and demonstrate that concomitant suppression of ferroelectric and conduction loss could be achieved by synergistically tuning the topological structure and phase composition of the multilayered P(VDF-HFP)/P(VDF-TrFE-CFE) (Co/Ter polymer) nanocomposites. The multilayered Co/Ter polymer nanocomposites are fabricated by a nonequilibrium process that combines electrospinning and fast thermal treatments. Simply by adjusting the quenching temperature, we are able to tune the phase composition of P(VDF-TrFE-CFE) terpolymer to achieve both a low ferroelectric loss and a high electric displacement. Results of the phase-field simulations show that the local electric fields are 6
substantially weakened at the interfaces between the alternating Co/Ter polymer layers, giving rise to suppressed leakage current hence low conduction loss and enhanced breakdown strength.
2. Results 2.1. Preparation of multilayered nanocomposites The P(VDF-HFP)/P(VDF-TrFE-CFE) (Co/Ter for short) multilayered nanocomposites were fabricated by a non-equilibrium processing which has been described in detail in our previous contribution.[26] Briefly, by integrating electrospinning, hot-pressing and thermal-quenching process, the non-equilibrium method is capable of yielding polymer films of very high quality. More importantly, this method provides a simple approach to the fine tuning of the microstructures, such as such as composition gradient or nanofiller orientation, of polymer nanocomposites.[26, 27] In this study, three Co/Ter multilayered nanocomposites with different layers, i.e., 4 layers, 8 layers and 16 layers, (4L, 8L, 16L for short, respectively) were fabricated, where every layer possesses the same thickness and the two same middle layers is designed to ensure same volume fractions of different multilayered nanocomposites, as shown in Figure 1e. Meanwhile, pure P(VDF-HFP) (pure copolymer for short) and P(VDF-TrFE-CFE) (pure terpolymer for short) films were also prepared as controls. Different from fabrication of pure polymer films with simple electrospinning process, alternating electrospinning was employed to prepare the multilayered nanocomposites. During the alternating electrospinning process, two different polymer solutions, i.e., 7
P(VDF-HFP) and P(VDF-TrFE-CFE) solutions, alternately electrospun. The thickness of each layer is controlled by adjusting the electrospinning time of the individual layers. For tuning of their phase-compositions, the multilayered nanocomposites is quenched from 200oC to water of different temperature after hot-pressing process. As shown in Figure 1a, b, both the as-spun copolymer and terpolymer nanofibers are bead-free and smooth, which is crucial for the subsequent hot-pressing process that finally transformed the fibrous mat into dense multilayered nanocomposites of high structural integrity. After hot-pressing and quenching process, smooth and dense multilayered nanocomposites were obtained, as evidenced by the surface and cross-sectional SEM images shown in Figure 1c, d. For a clearer observation of multilayer structure, 5 wt.% of BaTiO3 nanoparticles was incorporated into the terpolymer layer as an internal mark with minor effect, as shown in Figure S1. As seen, well-defined layered structure with same thickness of each constituent layer has been achieved.
2.2. Breakdown strength of multilayered nanocomposites Weibull statistics is employed for the analysis of the dielectric breakdown behavior of the multilayered nanocomposites. The breakdown strength is determined by fitting the experimental breakdown electric field to the Weibull equation as P(E) = 1-exp[-(E/Eb)], where P is the cumulative probability of electric failure, E is experimental breakdown strength, Eb is a scale parameter refers to the breakdown strength at the cumulative failure probability of 63.2% and is also regarded as the characteristic breakdown strength, and the shape parameter is the Weibull modulus that shows the dispersion of E.[20] The results of 8
Weibull statistical analysis for Co/Ter multilayered nanocomposites are presented in Figure S2. For a better comparison, the values of Eb parameter are summarized and plotted in Figure 2. Notably, two features could be distinguished in the dielectric breakdown behaviors of Co/Ter multilayered nanocomposites. First, substantial enhancement of Eb with the increasing number of inner layers is observed for the nanocomposites that are quenched at the same temperature. For instance, for the nanocomposites quenched at 0 ℃, the value of Eb increases from 465.7 MV/m for 4L to 545.6 MV/m for 8L and then reaches up to a maximum of 637.5 MV/m for 16L. It is of interest to note that albeit the relatively low Eb of terpolymer layers (~300 MV/m), the Co/Ter multilayered nanocomposites with increasing amount of terpolymer can still deliver a Eb that is comparable to or even higher than Eb of the copolymer layer. Second, with the same number of layers, the Co/Ter multilayered nanocomposites quenched at 45 ℃ exhibit much higher Eb compared with those quenched at either 0 oC or 60 ℃. This difference is predominant at samples with 4L and then becomes milder with increasing number of layers. For example, the Eb of 4L-45 ℃ nanocomposite (the succinct symbol of xL-y ℃ represents the sample with x number of layers quenched at y ℃) is 11.4% higher than that of 4L-0 ℃ sample, while the Eb of 16L-45 ℃ and 16L-0 ℃ samples are comparable. Plus, the values of all the Co/Ter multilayered nanocomposites shown in Figure S2 remain at a considerably high level (~20), suggesting a narrow distribution of Eb and excellent electrical reliability.
2.3. Discharge efficiency and energy density
9
The efficiencies and discharged energy densities of the Co/Ter multilayered nanocomposites and the pure copolymers and terpolymers are characterized by electric displacement–electric field (D-E) loops obtained by modified Sawyer-Tower circuit. (Figure S3 and Figure 5e) The discharge efficiency is defined as = Ue/U, where Ue and U are the discharged and stored energy densities, respectively, both of which can be derived from D-E loops by integration of the area between the discharged/charged curve and the ordinate. Ferroelectric PVDF-based polymers usually suffer from a low of ~50-70%, which limits their application in electrical insulation and energy storage devices despite of their high dielectric permittivity or high discharge energy density. As shown in Figure 3a, our Co/Ter multilayered nanocomposites exhibit an ultrahigh discharge efficiency up to 85% at a rather high electric field of ~ 612 MV/m. To the best of our knowledge, this is the highest discharge efficiency ever reported for PVDF-based dielectric polymers at such high electric field. This discharge efficiency is even comparable to that of some linear dielectrics, such as the bench-mark BOPP, at high electric field ( ~ 83% at ~ 646MV/m).[25] Also could be observed in Figure 3a is the monotonic increase of the discharge efficiency with increasing quenching temperature from 0 ℃ to 60 ℃ for both the multilayered nanocomposites and the pure polymer controls. Take the nanocomposites with 4L as an example, the discharge efficiency increases from 73.6% to 80.5% as the quenching temperature increases from 0 oC to 60 ℃. It is also of interest to note that higher discharge efficiency could be achieved by increasing the number of inner layers in the multilayered nanocomposites. As shown in Figure 3a, the Co/Ter multilayered nanocomposite with 16L quenched at 60 ℃exhibits an 10
ultrahigh discharge efficiency of 85.1%. To highlight the significance of our synergistic approach, we then plot the discharge efficiency as a function of the corresponding discharge energy density obtained at the breakdown strength (Eb) in Figure 3b, for the Co/Ter multilayered nanocomposites in this work and those PVDF-based dielectric polymers reported in previous literatures, including polymer blend, multilayered polymers and polymer nanocomposites, etc.. As shown in Figure 3b, for most of the PVDF-based dielectric polymers or nanocomposites,[12, 18, 19, 21, 28-36] the enhanced discharge energy density (~ 5-15 J/cm3) is substantially compromised by their low discharge efficiency of ~ 50-70%. In stark contrast, Co/Ter multilayered nanocomposites delivers an ultrahigh discharge efficiency of ~85% with high energy density of ~ 20 J/cm3 at a rather high electric field of ~ 600 MV/m. The higher discharge efficiency means that more energy density is released to external circuit rather than being dissipated in the polymer dielectrics, hence larger discharged energy density is obtained. As shown in Figure S4, Ue can reach up to 22.6 J/cm3 from the Co/Ter multilayered nanocomposites with 16L quenched at 0 ℃, which is an enhancement of ~130% over the pure copolymer or ~ 278% over the terpolymer quenched at the same temperature. The Co/Ter multilayered nanocomposites with 16L quenched at 45 oC or 60 oC also exhibit much enhanced discharge energy density compared with their co/terpolymer constituents (Figure S4).
3. Discussion In an effort to understand the enhanced breakdown strength, we characterize the Co/Ter multilayered nanocomposites in terms of mechanical and electrical properties and study the 11
dielectric breakdown behavior by phase-field simulations. It is commonly accepted that electromechanical breakdown is the major breakdown mechanism for polymer dielectrics. Theoretically, the electromechanical breakdown strength Eem can be expressed by the Stalk-Garton model as Eem = k(Y/(0r))1/2, where k is a constant and Y is the Young’s modulus.[37] When an electric field is applied in the out-of-plane direction of the nanocomposite films, the out-of-plane Young’s modulus of the films are evaluated by the nanoindentation method. As shown in Figure 4a, the Young’s modulus of the pure terpolymer is only 0.7 GPa, which is less than half of that of the pure copolymer (~1.5 GPa). Incorporated with the hard pure copolymer, the Co/Ter multilayered nanocomposites achieve a higher Y than pure terpolymer and hence an enhanced Eb. The results also show that the Y of multilayered nanocomposites increases with increasing number of layers, which agrees well with the tendency of experimental Eb in Figure 2 and observations reported in other studies.[38, 39] Therefore, it suggests that multilayered polymers may provide an approach to mitigate deleterious electromechanical effects in low modulus, high dielectric materials. Note that the pure copolymer has a higher modulus than multilayered nanocomposites with 16L, while a lower Eb. This could be attributed to the limitations of the Stalk-Garton model, where ideal insulators hence no electro-thermal coupling are considered. In fact, in addition to the electromechanical breakdown mechanism, other critical factors have to be taken into account in the consideration of electrical breakdown behavior of real dielectric polymers. The first factor that has to be considered is the inhomogeneous distribution of the local electric field induced by the mismatch between the dielectric permittivities of the composite 12
constituents. In the multilayered nanocomposites, the dielectric “soft” layers of high dielectric permittivity will “squeeze” the local electric field into the dielectric “hard” layers of low dielectric permittivity, resulting in substantial concentration of the local electric field hence early electrical breakdown in the layers of low dielectric permittivity. By quenching the multilayered nanocomposites at different temperatures, we are able to tune the dielectric permittivity of the constituent Co/Ter polymer layers for a balance. As seen in Figure 4b, the dielectric permittivity of the pure terpolymer shows a strong dependence on quenching temperature, which is due to the dependence of their Curie temperature (Figure S5). Yet the same heat treatment induces only minor changes in the dielectric permittivity of the pure copolymer. For instance, when the quenching temperature rises from 0 ℃ to 60 ℃, the dielectric permittivity (@ 1 kHz) increases from 27.0 to 50.9 for the terpolymer layers, while decreases marginally from 10.0 to 8.0 for the copolymer layers. It is thus speculated that the higher Eb of the the multilayered nanocomposites quenched at 45 ℃ may be attributed to the better match between the dielectric permittivities in the Co/Ter polymer layers hence the favorable local electric field distribution. In order to further verify this explanation, local electric field of multilayered nanocomposites is simulated by a phase-field model. As shown in Figure S6, the local electric field of the 4L-45 ℃ film is more uniform compared to the 4L-0 ℃ and 4L-60 ℃ films. The multilayer structure tends to suppress the leakage current leading to enhanced Eb. Presented in Figure 4c is the electric field-dependences of the leakage current for Co/Ter multilayered nanocomposites with increasing number of layers. In general, the leakage 13
current density decreases as more layers, hence more interfaces, are present especially at higher electric field. It indicates that the interface regions between the alternating layers may be at play in restricting the motion of charge carriers hence suppressing the leakage current. To elucidate the effects of the mesoscopic interfaces on electric breakdown process, a phase-field model is employed to simulate the distribution of local electric field and the time-domain evolution of electrical treeing process under an electric field along the out-of-plane direction of the multilayered nanocomposites. In this model, the interfaces between the alternating Co/Ter polymer layers are considered as a third interfacial phase with a thickness of ~ 100 nm and a dielectric permittivity of ~ 500.[22, 40-43] The top of Figure 4e exhibits the distribution of local electric field along the out-of-plane direction (z-direction) in the Co/Ter multilayered nanocomposites with 8L. The ratio between the local electric field (Ezlocal) in the constituent layers and the apparent electric field (Ezapp) is used as an indicator for the inhomogeneous distribution of local electric field. Due to the dielectric mismatch between the copolymer (r ~8.6) and terpolymer (r ~35.0) layers, a large Ezlocal/ Ezapp ratio of ~ 1.5 indicates substantially concentrated local electric field in the copolymer layers (as highlighted in red), whereas the local electric fields in the terpolymer layers are decreased (Ezlocal/ Ezapp ratio of ~ 0.5). The most striking feature in Figure 4e is that even weaker local electric field is observed within the interface regions between the Co/Ter polymer layers, as evidenced by a small Ezlocal/ Ezapp ratio of ~ 0.1 and highlighted in blue. We note that the much decreased local electric field at interfaces originates from high dielectric permittivity of interfacial phase induced by Maxwell-Wagner-Sillars (MWS) interface polarizations (see 14
Figure S7 & Section 6 in supporting information).[44] Combining with the effect of deep traps in the interface, the charge carriers first decelerate (loss their electric energy) due to weak local electric field and then are captured within the deep traps in the interface region.[45, 46] Hence, it is reasonable to infer that the charge carriers will suffer from more impediment when they encounter more interfaces on their way to breakdown. In the bottom of Figure 4e, the real-time evolutions of breakdown paths are presented in a three-dimensional view, where two features could be well distinguished. First, with the increasing number of layers, the breakdown paths become more tortuous and widen in the in-plane directions. In principle, more tortuous breakdown paths indicate a smaller probability of failure and hence a higher Eb. Second, the breakdown paths propagate faster in the nanocomposites with less layers than those with more layers. For instance, in the third column, a complete breakdown pathway is already formed in films with 4L, while partial breakdown paths are observed in the nanocomposites with 16L. The simulation results show that layered structures hinder the breakdown process by providing multiple barriers to the propagation of the breakdown pathway at interfaces between the alternating layers.[45] The effects of interface are further evidenced by the simulation results of the breakdown strength (Eb) by comparing Eb of the multilayered nanocomposites with and without interface. As shown in Figure 4d, the simulated Eb agrees reasonably well with the experimental ones with the contribution of interface considered. Yet, a much lower Eb is obtained from the simulations without interface. In addition, it is worth noting that when comparing the Ebs of multilayered nanocomposites with 16L and pure copolymer films, the former is higher than 15
the latter in the simulation with interface, which is in line with experimental results. Opposite results are obtained from the simulation without considering the interfaces. It means that besides electromechanical effect, the blocking effect of interface seems more important to enhance Eb in our systems. We then turn to the discussions on the effects of phase composition on the suppression of ferroelectric and conduction loss. First, tuning phase composition through quenching process is employed to lower ferroelectric loss. As shown in Figure 5a, the XRD pattern of the pure P(VDF-TrFE-CFE) terpolymer varies with the quenching temperature. There are two main peaks presented in the XRD pattern of the terpolymers, where the peak at 2 = 18.0°corresponds to the (110/200) reflection of non-polar phase while the other peak at 2 = 19.6° corresponds to the (110/200) reflection of polar phase.[47] The pure terpolymer quenched at 0 ℃ exhibits a mixture of non-polar and polar phase. With increasing quenching temperature, the content of the polar phase decreases from 26.2 wt.% at 0 ℃ to only 6.7 wt.% at 60 ℃ while that of the non-polar phase gradually increases from 22.9 wt.% at 0 ℃ to 42.2 wt.% at 60 ℃ (Table S1), where the non-polar phase is predominant as evidenced by the single peak at 2 = 18.0° in XRD pattern. Reducing the content of the polar phase will suppress ferroelectric loss and lead to a slimmer D-E loop (Figure 5e) and lowered dielectric loss 1- (Figure 5c). Similar to the case of pure terpolymer, the pure copolymer quenched at higher temperature possesses a smaller content of the polar phase evidenced by the XRD pattern in Figure 5b (refer to Section 6 in supporting information for detailed discussions) and thus lower dielectric loss (Figure 5d). Finally, the Co/Ter multilayered nanocomposites 16
consisting of phases of both constituents show a similar tendency, i.e., ferroelectric loss is substantially suppressed at higher quenching temperature (see in Figure S8). Second, we construct a multilayer structure to hinder leakage current and hence a suppressed conduction loss. The interfaces between copolymer and terpolymer act as barriers to the motion of charge carriers. As shown in Figure 4c, the leakage current decreases with increasing the number of layers, indicating an enhanced blocking effect. The suppressed leakage current then will lower the conduction loss. For instance, Figure S9 presents the dielectric loss (1-) of the Co/Ter multilayered nanocomposites quenched at 45 ℃ as a function of electric field. One can see that at the same electric field, the dielectric loss of the films quenched at same temperature decreases with the increasing number of layers, which should be attributed to the depressed conduction loss because of their similar ferroelectric loss.
It is of particular significance to note that, in previous contributions,[18, 19, 25, 33] the discharge efficiency decreases gradually with increasing electric field and then drops sharply at high electric field owing to the exponentially increased leakage current. In stark contrast, with increasing of electric field, the discharge efficiency of our multilayered nanocomposites firstly decreases mainly due to the dipole switching and field-induced phase transformation processes[21, 48, 49] and then increases gradually until electric breakdown, as shown in Figure S9. Hence, a higher discharge efficiency could be obtained with a larger Eb, as shown in the relationship of discharge efficiency and Eb of the Co/Ter multilayered nanocomposites (Figure 6a). To understand the enhanced discharge efficiency with higher Eb, the remnant 17
polarization Pr of 16L films is plotted as function of electric field in Figure 6b. Note that the Pr remains almost unchanged or even a little decreased at electric field >300 MV/m, while the maximum of polarization Pmax rises steadily with increasing electric field (Figure S10). It means that as electric field increased, the ratio of loss energy density to charged energy density decreases, i.e., the ratio of discharged energy density to charged energy density increases. As for the unchanged Pr, it results from the steady leakage current (Figure 4c) and saturated unswitched ferroelectric dipole at high electric field.[26]
4. Conclusions We propose a simple approach to design and fabricate PVDF-based dielectric polymers to exhibit unprecedentedly high discharge efficiency at high electric fields. Specifically, we fabricated P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites with extremely high charge-discharge efficiency (80~85%) and high discharged energy density (~20 J/cm3) through a non-equilibrium processing. By optimizing the topological structure of the multilayered nanocomposites, we simultaneously achieve suppressed leakage current, enhanced breakdown strength and reduced conduction loss. It is attributed to the multiple blocking effects of the interfaces weakening electric fields. Through tuning the phase composition, we were able to manipulate the distribution of local electric field to obtain a higher breakdown strength, and suppress the ferroelectric loss. It is also found that besides the concomitant suppression of ferroelectric and conduction losses, the enhanced breakdown strength is of particular significance for achieving an ultrahigh charge-discharge efficiency in 18
our multilayered nanocomposites, where the remnant polarization remains unchanged at high electric fields.
5. Experimental Section 5.1. Materials P(VDF-HFP) and P(VDF-TrFE-CFE) were purchased from Arkema and used as received. All solvents were obtained from China National Chemicals Corporation Ltd. and used without further purification. 5.2. Fabrication of P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites The non-equilibrium process was applied to prepare the multilayered nanocomposites. Briefly, P(VDF-HFP) and P(VDF-TrFE-CFE) powders were thoroughly dissolved in a mixed solvent of N,N-dimethylformamide (DMF) and acetone, respectively. The obtained solutions as electrospinning solutions were transferred into two syringes, respectively. A modified electrospinning process was performed to fabricate the multilayered fibrous mats. Take fabrication of a multilayered mat with four layers for example, the first layer, P(VDF-HFP) fiber, was electrospun at an applied field of 1.0 kV/cm with flow rate of 1.0 mL/h, followed by 2nd to 4th layers, i.e., P(VDF-TrFE-CFE) fiber layer, P(VDF-TrFE-CFE) fiber layer and P(VDF-HFP) fiber layer with same electrospinning conditions. The similar process was applied to fabricate the multilayered mats with eight and sixteen layers, respectively. However, the electrospinning time of each layer varied with the number of layer in different multilayered mats, which ensured the same thickness of final films. The fibrous mats were 19
then transformed into dense body by hot-pressing at 200 ℃ for 1 h under a pressure of 10 MPa. 5.3. The quenching process The as-pressed films were annealed at 200 ℃ for 15 min followed by quenching in water with different temperature, i.e., 0 ℃, 45 ℃ and 60 ℃, respectively. All of the samples had equal volume amount of P(VDF-HFP) and P(VDF-TrFE-CFE) and an overall film thickness of 15 μm. Characterization: The morphology of P(VDF-HFP) nanofibers, P(VDF-TrFE-CFE) nanofibers and multilayered nanocomposites were characterized with scanning electron microscopy (SEM, ZEISS MWRLIN compact). X-ray diffraction (XRD, Rigaku SmartLab) was performed for all samples. Dielectric constant were measured by employing a HP 4294A precision impedance analyzer (Agilent Technologies) at room temperature within the frequency range of 102 to 107 Hz at 1 Vrms. The Curie temperature of pure P(VDF-TrFE-CFE) films were tested by using the same impedance analyzer at 1 Hz and 1 Vrms within the temperature range of 30 ℃ to 100 ℃. Electric breakdown tests were performed with dielectric withstand voltage test system (Beijing Electro-mechanical Research Institute Supesvoltage 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.). The leakage current as a function of electric field were measured by the same ferroelectric test system. 5.4. Phase-field Simulations 20
In the phase-field model, two sets of order parameters were used in this model: a nonevolving variable r , with the value of 1 in copolymer matrix and 0 in terpolymer matrix, and a continuous phase field variable r,t which is both position- and time-dependent, where
r, t =1 represents the broken phase and r, t =0 represents the unbroken phase. It has to be noted that the electric inhomogeneity of the dielectric composite in this work is defined according to the different dielectric constants of the different matrix phase, the interface phase, and the electrically broken down phase. The
total
free
energy
F
of
the
dielectric
composite
can
be
written
as
1 2 F f sep g f elec 2 V (1) where the first term is phase separation energy density, the second term signifies the gradient energy density with the gradient energy coefficient g, and the last term represents the electrostatic energy density. The separation energy is expressed as
fsep =a 2 1
2
(2) where a is a coefficient related with the energy barrier of the phase transition. The electrostatic energy density f elec can be calculated as
21
1 1 f elec = 0 ij r Ei r E j r Ei r Pi S r 2 2 (3) where ij r is the position-dependent dielectric constant tensor, Ei r is the electric field component, Pi S r is the spontaneous polarization component. Therein the electric field distributions E r can be obtained by solving electrostatic Poisson equation 0 E P 0 using a Fourier spectral iterative perturbation method, which has been used
in previous work.[50, 51] The evolution of dielectric breakdown was simulated by solving a modified Allen-Cahn equation,
f r, t f elec L0 H f elec f critical sep g2 r, t t r, t r, t (4) where L0 is the kinetic coefficient related to the interface mobility, H f elec f critical is the Heaviside unit step function
H f
elec
f critical 0 and H f elec f critical =1 , f critical is
position-dependent material constant related to the maximal energy storage density of each constitute in the composite. In our phase field simulation, the characteristic length scale l0 g / a and similarly a characteristic time scale t 1 / L0a can be defined in terms of the material parameters g, a and L0. A grid size of 256x 256y 256z with grid spacing for x 1l0 and a time interval t 0.05t0 was used to yield a stable solution. Here, the initial parameters used in 22
this model, such as the breakdown strength and the dielectric constant of different pure polymers, are obtained from the experimental results. In order to reflect the interface effect, a 100 nm-thick interface region between adjacent layers with the dielectric permittivity of ~ 500 is introduced into our model. The coefficient in the separation energy term is given a value of 108 J/m3 to represent the barrier of the breakdown phase and the non-breakdown phase. A value of 10-10 J/m is used for the gradient energy coefficient to specify d0 1nm in the modeling.
Acknowledgment This work was supported by Basic Science Center Program of NSFC (Grant No. 51788104), the NSF of China (Grant No. 51625202, and 51572141), the National Key Research & Development Program (Grant No. 2017YFB0701603), the National Basic Research Program of China (Grant No. 2015CB654603) and Research Fund of Science and Technology in Shenzhen (JSGG20150331155519130).
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27
Fig 1. SEM images of (a) P(VDF-HFP) nanofibers, (b) P(VDF-TrFE-CFE) nanofibers, and (c) surface and (d) cross sections of Co/Ter multilayered nanocomposites with eight layers, respectively.(e) schematic illustration of the structures of the Co/Ter multilayered nanocomposites and the controls (pure copolymer and pure terpolymer), respectively.
28
Fig 2. Comparison of breakdown strength for Co/Ter multilayered nanocomposites and pure terpolymer and copolymer films.
29
Fig 3. (a) Efficiency of Co/Ter multilayered nanocomposites and pure terpolymer and copolymer films. (b) Comparison of efficiency and discharged energy density (at Eb) in this work and the related PVDF-based dielectric polymers reported in the previous literatures. 30
Fig 4. Breakdown mechanisms for Co/Ter multilayered nanocomposites. (a) Young’s modulus of Co/Ter multilayered nanocomposites quenched at 45 ℃. (b) Dielectric permittivity of pure copolymer and terpolymer quenched at different temperatures as a function of frequency. (c) Leakage current of Co/Ter multilayered nanocomposites quenched at 45 ℃ at different electric field. (d) Comparison of experimental and simulated breakdown strength of Co/Ter multilayered nanocomposites quenched at 45 ℃. (e) The distribution of local electric field along the direction of the applied field in the Co/Ter multilayered nanocomposites with 8L (top). Time-dependent 3D views of the breakdown evolution procedures for Co/Ter multilayered nanocomposites with 4L, 8L and 16L, respectively, by phase field simulation (bottom).
31
Fig 5. XRD patterns for (a) pure terpolymer and (b) pure copolymer quenched at different temperatures. Energy loss (1-) for (c) pure terpolymer and (d) pure copolymer quenched at different temperatures as a function of electric field. (e) D-E loops for pure terpolymer quenched at 0 ℃, 45 ℃ and 60 ℃, respectively, at varied electric field.
32
Fig 6. (a) Relationship of breakdown strength with efficiency in Co/Ter multilayered nanocomposites. (b) The remnant polarization of Co/Ter multilayered nanocomposites with 16L quenched at different temperatures as a function of electric field
33
PII: DOI: Reference:
S2405-8297(18)30838-9 https://doi.org/10.1016/j.ensm.2018.09.013 ENSM506
To appear in: Energy Storage Materials Received date: 6 July 2018 Revised date: 14 September 2018 Accepted date: 14 September 2018 Cite this article as: Jianyong Jiang, Zhonghui Shen, Jianfeng Qian, Zhenkang Dan, Mengfan Guo, Yuanhua Lin, Ce-Wen Nan, Longqing Chen and Yang Shen, Ultrahigh Discharge Efficiency in Multilayered Polymer Nanocomposites of High Energy Density, Energy Storage Materials, https://doi.org/10.1016/j.ensm.2018.09.013 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Ultrahigh
Discharge
Efficiency
in
Multilayered
Polymer Nanocomposites of High Energy Density Jianyong Jianga,d, Zhonghui Shena, Jianfeng Qiana, Zhenkang Dana, Mengfan Guoa, Yuanhua Lina, Ce-Wen Nana, Longqing Chena,b, Yang Shena,c* a
State Key Lab of New Ceramics and Fine Processing, School of Materials Science and
Engineering, Tsinghua University, Beijing, 100084, China b
Department of Materials Science and Engineering, The Pennsylvania State University,
University Park, Pennsylvania 16802, United States c
Center for Flexible Electronics Technology, Tsinghua University, Beijing 100084, China
d
Laboratory of Advanced Energy Storage Materials & Devices, Research Institute of
Tsinghua University in Shenzhen, Shenzhen, 518057, China
*
Corresponding author. [email protected]
Abstract
Poly(vinylidene fluoride) (PVDF)-based dielectric polymers are in great demand for the future electronic and electrical industry because of their high dielectric constants and energy density. However, some issues that limit their practical applications remain unsolved. One of the most urgent issues is their high dielectric loss and hence low efficiency. In this contribution, we proposed and demonstrate that substantially enhanced discharge efficiency of PVDF-based polymers nanocomposites could be achieved by simultaneously optimizing 1
their
topological-structure
and
fluoride-co-hexafluoropropylene)
phase
composition.
(P(VDF-HFP))
In
the /
poly(vinylidene poly(vinylidene
fluoride-ter-trifluoroethylene-ter-chlorofluoroethylene) (P(VDF-TrFE-CFE)) multilayered nanocomposites fabricated by non-equilibrium process, an ultrahigh discharge efficiency of ~85% is achieved up to 600 MV/m, which is the highest discharge efficiency reported so far for any polar-polymer dielectric materials at such high electric field. By adjusting the quenching temperature, the phase-composition hence dielectric permittivity in the terpolymer layers could be tuned for suppressed ferroelectric loss. Results of phase-field simulations further reveal that local electric field is substantially weakened at the interfaces between the Co/Ter polymer layers, which will act as barriers to motion of charge carriers and give rise to much suppressed conduction loss and a remarkably enhanced breakdown strength. Synergy of the optimized topological-structure and phase-composition thus leads to a nanocomposite that exhibits an unprecedented high discharge efficiency of the multilayered nanocomposites that is comparable to the bench-mark biaxially oriented polypropylene (BOPP) at high electric field as well as a high discharge energy density that is over 10 times higher than that of BOPP.
KEYWORDS: PVDF; charge-discharge efficiency; energy density; multilayer structure; phase tuning; interface.
2
Graphical Abstract
Synergy of the optimized topological-structure and phase-composition is employed to suppress the conduction loss and ferroelectric loss in P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites. An unprecedented high discharge efficiency of ~85% is achieved up to ~600 MV/m, which is by far the highest discharge efficiency ever achieved in polar-polymer dielectric materials at such high electric field. More importantly, a high energy density of ~20 J/cm3 is also achieved.
1. Introduction Dielectric materials with high power density are of critical significance for pulsed power systems, smart grid and electric vehicles.[1-3] Compared to other electrochemical energy storage devices such as supercapacitors and lithium-ion batteries, the electrostatic capacitors based
on
dielectric
media
store
and
release
electrical
energy
through
the
polarization/depolarization process of dipole in dielectric materials.[4] Among all available dielectric materials, polymers are of considerable interest because of their advantages such as 3
high reliability, low cost, lightweight and ease of fabrication.[5-7] Albeit their higher power density, dielectric polymers exhibit rather low energy density and thus large volume and weight for applications.[8] For instance, in a combat vehicle, the pulse power loads, including inverters, converters, power distribution and PFN (Pulse Forming Network), etc., can occupy ~44% of power system volume.[9] In general, the discharged energy density Ue of linear dielectric materials is given by Ue=1/20rEb2, where 0 is the vacuum permittivity, r and Eb are the effective permittivity and the breakdown strength of the dielectrics, respectively.[10] Most of the dielectric polymers possess comparatively high Eb while rather low r. For instance, biaxially oriented polypropylene (BOPP), which is the bench-mark dielectric polymer of current use, only exhibits an r of 2.2 (1 kHz), hence a low discharge energy density of ~ 2 J/cm3 despite their high Eb of ~ 700 MV/m.[11] Recently, Poly(vinylidene fluoride) PVDF-based dielectric polymers with high dielectric permittivity (10~50), stemming from high polarization of the C-F bonds and spontaneous oriention of the dipoles in the crystalline domains, are drawing increasing attention for applications in high energy density
capacitors.
PVDF-based
copolymers
such
as
poly(vinylidene
fluoride-co-hexafluoropropylene), P(VDF-HFP), deliver discharge energy density of 13.5 J/cm3 and efficiency of 62% with an Eb of 600 MV/m.[12] Another class of PVDF-based polymers
is
the
terpolymers
including
poly(vinylidene
fluoride-ter-trifluoroethylene-ter-chlorofluoroethylene), P(VDF-TrFE-CFE), poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene), poly(vinylidene
P(VDF-TrFE-CTFE),
fluoride-ter-trifluoroethylene-ter-hexafluoropropylene), 4
P(VDF-TrFE-HFP),[13] which enjoy the highest r (~50 at 1 kHz) among all the dielectric polymers explored so far. Relaxor ferroelectric behaviors are observed for the terpolymers, which endow the terpolymers with a high Ue at low electric field. For example, a high Ue ~10 J/cm3 is achieved at an electric field of ~350 MV/m in P(VDF-TrFE-CFE).[14] To further increase r and Ue, nanosized ceramic fillers with high r such as BaTiO3, Ba1-xSrxTiO3, have been incorporated into PVDF-based polymers to form polymer nanocomposites.[15, 16] We recently showed that by incorporating TiO2 nanofibers embedded with BaTiO3 nanoparticles (BTO@TO_nfs) into P(VDF-HFP) matrix, hierarchical interfaces are introduced and give rise to concomitant enhancement of r and Eb, leading to a giant discharge energy density of ~31.2 J/cm3 and efficiency of 78% at 797.7 MV/m.[17] Albeit high energy density realized in PVDF-based dielectric polymers and their nanocomposites, the discharge efficiency of these dielectric materials is still low ( of ~ 50%-70%) at high electric fields.[1, 18] It means that more than 1/5 of stored electrical energy is not recovered and dissipated in the form of Joule heat that accumulates within the polymer and may cause early failure of these dielectric materials at an electric field much lower than the intrinsic Eb of the polymer matrix. The unreleased energy is also called the dissipated energy or dielectric loss. There are mainly two types of dielectric loss: ferroelectric loss and conduction loss.[19-22] The ferroelectric loss is induced during the ferroelectric switching when the dipolar switching does not synchronize with the applied alternating electric field. As a semi-crystalline ferroelectric polymer, the ferroelectric loss of PVDF-based polymers is closely related to their phase structures and crystalline size. 5
Specifically, compared to polar phase of PVDF with the all-trans configuration hence large ferroelectric loss, the non-polar phase with zig-zag configuration exhibits smaller ferroelectric loss.[23] Terpolymers, such as P(VDF-TrFE-CFE), with small crystalline size due to the nano-confinement effect of large CFE group, exhibit relaxor ferroelectric behavior where the dipoles rotate and respond rapidly to the external electrical field leading to a low ferroelectric loss.[24] Different from the prominent ferroelectric loss that is intrinsic to the ferroelectric polymers, conduction loss is more common to all dielectric polymers. Biaxially-oriented Polypropylene (BOPP), a paraelectric polymer without permanent dipolar moments along the chains exhibits almost no ferroelectric loss at low electric fields but a rather high conduction loss of ~15% at ~ 600 MV/m due to much increased leakage current at high electric fields.[20, 25] Hence, in order to achieve a high discharge efficiency, both ferroelectric and conduction losses should be suppressed concurrently in PVDF-based dielectric polymers. In this contribution, we propose and demonstrate that concomitant suppression of ferroelectric and conduction loss could be achieved by synergistically tuning the topological structure and phase composition of the multilayered P(VDF-HFP)/P(VDF-TrFE-CFE) (Co/Ter polymer) nanocomposites. The multilayered Co/Ter polymer nanocomposites are fabricated by a nonequilibrium process that combines electrospinning and fast thermal treatments. Simply by adjusting the quenching temperature, we are able to tune the phase composition of P(VDF-TrFE-CFE) terpolymer to achieve both a low ferroelectric loss and a high electric displacement. Results of the phase-field simulations show that the local electric fields are 6
substantially weakened at the interfaces between the alternating Co/Ter polymer layers, giving rise to suppressed leakage current hence low conduction loss and enhanced breakdown strength.
2. Results 2.1. Preparation of multilayered nanocomposites The P(VDF-HFP)/P(VDF-TrFE-CFE) (Co/Ter for short) multilayered nanocomposites were fabricated by a non-equilibrium processing which has been described in detail in our previous contribution.[26] Briefly, by integrating electrospinning, hot-pressing and thermal-quenching process, the non-equilibrium method is capable of yielding polymer films of very high quality. More importantly, this method provides a simple approach to the fine tuning of the microstructures, such as such as composition gradient or nanofiller orientation, of polymer nanocomposites.[26, 27] In this study, three Co/Ter multilayered nanocomposites with different layers, i.e., 4 layers, 8 layers and 16 layers, (4L, 8L, 16L for short, respectively) were fabricated, where every layer possesses the same thickness and the two same middle layers is designed to ensure same volume fractions of different multilayered nanocomposites, as shown in Figure 1e. Meanwhile, pure P(VDF-HFP) (pure copolymer for short) and P(VDF-TrFE-CFE) (pure terpolymer for short) films were also prepared as controls. Different from fabrication of pure polymer films with simple electrospinning process, alternating electrospinning was employed to prepare the multilayered nanocomposites. During the alternating electrospinning process, two different polymer solutions, i.e., 7
P(VDF-HFP) and P(VDF-TrFE-CFE) solutions, alternately electrospun. The thickness of each layer is controlled by adjusting the electrospinning time of the individual layers. For tuning of their phase-compositions, the multilayered nanocomposites is quenched from 200oC to water of different temperature after hot-pressing process. As shown in Figure 1a, b, both the as-spun copolymer and terpolymer nanofibers are bead-free and smooth, which is crucial for the subsequent hot-pressing process that finally transformed the fibrous mat into dense multilayered nanocomposites of high structural integrity. After hot-pressing and quenching process, smooth and dense multilayered nanocomposites were obtained, as evidenced by the surface and cross-sectional SEM images shown in Figure 1c, d. For a clearer observation of multilayer structure, 5 wt.% of BaTiO3 nanoparticles was incorporated into the terpolymer layer as an internal mark with minor effect, as shown in Figure S1. As seen, well-defined layered structure with same thickness of each constituent layer has been achieved.
2.2. Breakdown strength of multilayered nanocomposites Weibull statistics is employed for the analysis of the dielectric breakdown behavior of the multilayered nanocomposites. The breakdown strength is determined by fitting the experimental breakdown electric field to the Weibull equation as P(E) = 1-exp[-(E/Eb)], where P is the cumulative probability of electric failure, E is experimental breakdown strength, Eb is a scale parameter refers to the breakdown strength at the cumulative failure probability of 63.2% and is also regarded as the characteristic breakdown strength, and the shape parameter is the Weibull modulus that shows the dispersion of E.[20] The results of 8
Weibull statistical analysis for Co/Ter multilayered nanocomposites are presented in Figure S2. For a better comparison, the values of Eb parameter are summarized and plotted in Figure 2. Notably, two features could be distinguished in the dielectric breakdown behaviors of Co/Ter multilayered nanocomposites. First, substantial enhancement of Eb with the increasing number of inner layers is observed for the nanocomposites that are quenched at the same temperature. For instance, for the nanocomposites quenched at 0 ℃, the value of Eb increases from 465.7 MV/m for 4L to 545.6 MV/m for 8L and then reaches up to a maximum of 637.5 MV/m for 16L. It is of interest to note that albeit the relatively low Eb of terpolymer layers (~300 MV/m), the Co/Ter multilayered nanocomposites with increasing amount of terpolymer can still deliver a Eb that is comparable to or even higher than Eb of the copolymer layer. Second, with the same number of layers, the Co/Ter multilayered nanocomposites quenched at 45 ℃ exhibit much higher Eb compared with those quenched at either 0 oC or 60 ℃. This difference is predominant at samples with 4L and then becomes milder with increasing number of layers. For example, the Eb of 4L-45 ℃ nanocomposite (the succinct symbol of xL-y ℃ represents the sample with x number of layers quenched at y ℃) is 11.4% higher than that of 4L-0 ℃ sample, while the Eb of 16L-45 ℃ and 16L-0 ℃ samples are comparable. Plus, the values of all the Co/Ter multilayered nanocomposites shown in Figure S2 remain at a considerably high level (~20), suggesting a narrow distribution of Eb and excellent electrical reliability.
2.3. Discharge efficiency and energy density
9
The efficiencies and discharged energy densities of the Co/Ter multilayered nanocomposites and the pure copolymers and terpolymers are characterized by electric displacement–electric field (D-E) loops obtained by modified Sawyer-Tower circuit. (Figure S3 and Figure 5e) The discharge efficiency is defined as = Ue/U, where Ue and U are the discharged and stored energy densities, respectively, both of which can be derived from D-E loops by integration of the area between the discharged/charged curve and the ordinate. Ferroelectric PVDF-based polymers usually suffer from a low of ~50-70%, which limits their application in electrical insulation and energy storage devices despite of their high dielectric permittivity or high discharge energy density. As shown in Figure 3a, our Co/Ter multilayered nanocomposites exhibit an ultrahigh discharge efficiency up to 85% at a rather high electric field of ~ 612 MV/m. To the best of our knowledge, this is the highest discharge efficiency ever reported for PVDF-based dielectric polymers at such high electric field. This discharge efficiency is even comparable to that of some linear dielectrics, such as the bench-mark BOPP, at high electric field ( ~ 83% at ~ 646MV/m).[25] Also could be observed in Figure 3a is the monotonic increase of the discharge efficiency with increasing quenching temperature from 0 ℃ to 60 ℃ for both the multilayered nanocomposites and the pure polymer controls. Take the nanocomposites with 4L as an example, the discharge efficiency increases from 73.6% to 80.5% as the quenching temperature increases from 0 oC to 60 ℃. It is also of interest to note that higher discharge efficiency could be achieved by increasing the number of inner layers in the multilayered nanocomposites. As shown in Figure 3a, the Co/Ter multilayered nanocomposite with 16L quenched at 60 ℃exhibits an 10
ultrahigh discharge efficiency of 85.1%. To highlight the significance of our synergistic approach, we then plot the discharge efficiency as a function of the corresponding discharge energy density obtained at the breakdown strength (Eb) in Figure 3b, for the Co/Ter multilayered nanocomposites in this work and those PVDF-based dielectric polymers reported in previous literatures, including polymer blend, multilayered polymers and polymer nanocomposites, etc.. As shown in Figure 3b, for most of the PVDF-based dielectric polymers or nanocomposites,[12, 18, 19, 21, 28-36] the enhanced discharge energy density (~ 5-15 J/cm3) is substantially compromised by their low discharge efficiency of ~ 50-70%. In stark contrast, Co/Ter multilayered nanocomposites delivers an ultrahigh discharge efficiency of ~85% with high energy density of ~ 20 J/cm3 at a rather high electric field of ~ 600 MV/m. The higher discharge efficiency means that more energy density is released to external circuit rather than being dissipated in the polymer dielectrics, hence larger discharged energy density is obtained. As shown in Figure S4, Ue can reach up to 22.6 J/cm3 from the Co/Ter multilayered nanocomposites with 16L quenched at 0 ℃, which is an enhancement of ~130% over the pure copolymer or ~ 278% over the terpolymer quenched at the same temperature. The Co/Ter multilayered nanocomposites with 16L quenched at 45 oC or 60 oC also exhibit much enhanced discharge energy density compared with their co/terpolymer constituents (Figure S4).
3. Discussion In an effort to understand the enhanced breakdown strength, we characterize the Co/Ter multilayered nanocomposites in terms of mechanical and electrical properties and study the 11
dielectric breakdown behavior by phase-field simulations. It is commonly accepted that electromechanical breakdown is the major breakdown mechanism for polymer dielectrics. Theoretically, the electromechanical breakdown strength Eem can be expressed by the Stalk-Garton model as Eem = k(Y/(0r))1/2, where k is a constant and Y is the Young’s modulus.[37] When an electric field is applied in the out-of-plane direction of the nanocomposite films, the out-of-plane Young’s modulus of the films are evaluated by the nanoindentation method. As shown in Figure 4a, the Young’s modulus of the pure terpolymer is only 0.7 GPa, which is less than half of that of the pure copolymer (~1.5 GPa). Incorporated with the hard pure copolymer, the Co/Ter multilayered nanocomposites achieve a higher Y than pure terpolymer and hence an enhanced Eb. The results also show that the Y of multilayered nanocomposites increases with increasing number of layers, which agrees well with the tendency of experimental Eb in Figure 2 and observations reported in other studies.[38, 39] Therefore, it suggests that multilayered polymers may provide an approach to mitigate deleterious electromechanical effects in low modulus, high dielectric materials. Note that the pure copolymer has a higher modulus than multilayered nanocomposites with 16L, while a lower Eb. This could be attributed to the limitations of the Stalk-Garton model, where ideal insulators hence no electro-thermal coupling are considered. In fact, in addition to the electromechanical breakdown mechanism, other critical factors have to be taken into account in the consideration of electrical breakdown behavior of real dielectric polymers. The first factor that has to be considered is the inhomogeneous distribution of the local electric field induced by the mismatch between the dielectric permittivities of the composite 12
constituents. In the multilayered nanocomposites, the dielectric “soft” layers of high dielectric permittivity will “squeeze” the local electric field into the dielectric “hard” layers of low dielectric permittivity, resulting in substantial concentration of the local electric field hence early electrical breakdown in the layers of low dielectric permittivity. By quenching the multilayered nanocomposites at different temperatures, we are able to tune the dielectric permittivity of the constituent Co/Ter polymer layers for a balance. As seen in Figure 4b, the dielectric permittivity of the pure terpolymer shows a strong dependence on quenching temperature, which is due to the dependence of their Curie temperature (Figure S5). Yet the same heat treatment induces only minor changes in the dielectric permittivity of the pure copolymer. For instance, when the quenching temperature rises from 0 ℃ to 60 ℃, the dielectric permittivity (@ 1 kHz) increases from 27.0 to 50.9 for the terpolymer layers, while decreases marginally from 10.0 to 8.0 for the copolymer layers. It is thus speculated that the higher Eb of the the multilayered nanocomposites quenched at 45 ℃ may be attributed to the better match between the dielectric permittivities in the Co/Ter polymer layers hence the favorable local electric field distribution. In order to further verify this explanation, local electric field of multilayered nanocomposites is simulated by a phase-field model. As shown in Figure S6, the local electric field of the 4L-45 ℃ film is more uniform compared to the 4L-0 ℃ and 4L-60 ℃ films. The multilayer structure tends to suppress the leakage current leading to enhanced Eb. Presented in Figure 4c is the electric field-dependences of the leakage current for Co/Ter multilayered nanocomposites with increasing number of layers. In general, the leakage 13
current density decreases as more layers, hence more interfaces, are present especially at higher electric field. It indicates that the interface regions between the alternating layers may be at play in restricting the motion of charge carriers hence suppressing the leakage current. To elucidate the effects of the mesoscopic interfaces on electric breakdown process, a phase-field model is employed to simulate the distribution of local electric field and the time-domain evolution of electrical treeing process under an electric field along the out-of-plane direction of the multilayered nanocomposites. In this model, the interfaces between the alternating Co/Ter polymer layers are considered as a third interfacial phase with a thickness of ~ 100 nm and a dielectric permittivity of ~ 500.[22, 40-43] The top of Figure 4e exhibits the distribution of local electric field along the out-of-plane direction (z-direction) in the Co/Ter multilayered nanocomposites with 8L. The ratio between the local electric field (Ezlocal) in the constituent layers and the apparent electric field (Ezapp) is used as an indicator for the inhomogeneous distribution of local electric field. Due to the dielectric mismatch between the copolymer (r ~8.6) and terpolymer (r ~35.0) layers, a large Ezlocal/ Ezapp ratio of ~ 1.5 indicates substantially concentrated local electric field in the copolymer layers (as highlighted in red), whereas the local electric fields in the terpolymer layers are decreased (Ezlocal/ Ezapp ratio of ~ 0.5). The most striking feature in Figure 4e is that even weaker local electric field is observed within the interface regions between the Co/Ter polymer layers, as evidenced by a small Ezlocal/ Ezapp ratio of ~ 0.1 and highlighted in blue. We note that the much decreased local electric field at interfaces originates from high dielectric permittivity of interfacial phase induced by Maxwell-Wagner-Sillars (MWS) interface polarizations (see 14
Figure S7 & Section 6 in supporting information).[44] Combining with the effect of deep traps in the interface, the charge carriers first decelerate (loss their electric energy) due to weak local electric field and then are captured within the deep traps in the interface region.[45, 46] Hence, it is reasonable to infer that the charge carriers will suffer from more impediment when they encounter more interfaces on their way to breakdown. In the bottom of Figure 4e, the real-time evolutions of breakdown paths are presented in a three-dimensional view, where two features could be well distinguished. First, with the increasing number of layers, the breakdown paths become more tortuous and widen in the in-plane directions. In principle, more tortuous breakdown paths indicate a smaller probability of failure and hence a higher Eb. Second, the breakdown paths propagate faster in the nanocomposites with less layers than those with more layers. For instance, in the third column, a complete breakdown pathway is already formed in films with 4L, while partial breakdown paths are observed in the nanocomposites with 16L. The simulation results show that layered structures hinder the breakdown process by providing multiple barriers to the propagation of the breakdown pathway at interfaces between the alternating layers.[45] The effects of interface are further evidenced by the simulation results of the breakdown strength (Eb) by comparing Eb of the multilayered nanocomposites with and without interface. As shown in Figure 4d, the simulated Eb agrees reasonably well with the experimental ones with the contribution of interface considered. Yet, a much lower Eb is obtained from the simulations without interface. In addition, it is worth noting that when comparing the Ebs of multilayered nanocomposites with 16L and pure copolymer films, the former is higher than 15
the latter in the simulation with interface, which is in line with experimental results. Opposite results are obtained from the simulation without considering the interfaces. It means that besides electromechanical effect, the blocking effect of interface seems more important to enhance Eb in our systems. We then turn to the discussions on the effects of phase composition on the suppression of ferroelectric and conduction loss. First, tuning phase composition through quenching process is employed to lower ferroelectric loss. As shown in Figure 5a, the XRD pattern of the pure P(VDF-TrFE-CFE) terpolymer varies with the quenching temperature. There are two main peaks presented in the XRD pattern of the terpolymers, where the peak at 2 = 18.0°corresponds to the (110/200) reflection of non-polar phase while the other peak at 2 = 19.6° corresponds to the (110/200) reflection of polar phase.[47] The pure terpolymer quenched at 0 ℃ exhibits a mixture of non-polar and polar phase. With increasing quenching temperature, the content of the polar phase decreases from 26.2 wt.% at 0 ℃ to only 6.7 wt.% at 60 ℃ while that of the non-polar phase gradually increases from 22.9 wt.% at 0 ℃ to 42.2 wt.% at 60 ℃ (Table S1), where the non-polar phase is predominant as evidenced by the single peak at 2 = 18.0° in XRD pattern. Reducing the content of the polar phase will suppress ferroelectric loss and lead to a slimmer D-E loop (Figure 5e) and lowered dielectric loss 1- (Figure 5c). Similar to the case of pure terpolymer, the pure copolymer quenched at higher temperature possesses a smaller content of the polar phase evidenced by the XRD pattern in Figure 5b (refer to Section 6 in supporting information for detailed discussions) and thus lower dielectric loss (Figure 5d). Finally, the Co/Ter multilayered nanocomposites 16
consisting of phases of both constituents show a similar tendency, i.e., ferroelectric loss is substantially suppressed at higher quenching temperature (see in Figure S8). Second, we construct a multilayer structure to hinder leakage current and hence a suppressed conduction loss. The interfaces between copolymer and terpolymer act as barriers to the motion of charge carriers. As shown in Figure 4c, the leakage current decreases with increasing the number of layers, indicating an enhanced blocking effect. The suppressed leakage current then will lower the conduction loss. For instance, Figure S9 presents the dielectric loss (1-) of the Co/Ter multilayered nanocomposites quenched at 45 ℃ as a function of electric field. One can see that at the same electric field, the dielectric loss of the films quenched at same temperature decreases with the increasing number of layers, which should be attributed to the depressed conduction loss because of their similar ferroelectric loss.
It is of particular significance to note that, in previous contributions,[18, 19, 25, 33] the discharge efficiency decreases gradually with increasing electric field and then drops sharply at high electric field owing to the exponentially increased leakage current. In stark contrast, with increasing of electric field, the discharge efficiency of our multilayered nanocomposites firstly decreases mainly due to the dipole switching and field-induced phase transformation processes[21, 48, 49] and then increases gradually until electric breakdown, as shown in Figure S9. Hence, a higher discharge efficiency could be obtained with a larger Eb, as shown in the relationship of discharge efficiency and Eb of the Co/Ter multilayered nanocomposites (Figure 6a). To understand the enhanced discharge efficiency with higher Eb, the remnant 17
polarization Pr of 16L films is plotted as function of electric field in Figure 6b. Note that the Pr remains almost unchanged or even a little decreased at electric field >300 MV/m, while the maximum of polarization Pmax rises steadily with increasing electric field (Figure S10). It means that as electric field increased, the ratio of loss energy density to charged energy density decreases, i.e., the ratio of discharged energy density to charged energy density increases. As for the unchanged Pr, it results from the steady leakage current (Figure 4c) and saturated unswitched ferroelectric dipole at high electric field.[26]
4. Conclusions We propose a simple approach to design and fabricate PVDF-based dielectric polymers to exhibit unprecedentedly high discharge efficiency at high electric fields. Specifically, we fabricated P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites with extremely high charge-discharge efficiency (80~85%) and high discharged energy density (~20 J/cm3) through a non-equilibrium processing. By optimizing the topological structure of the multilayered nanocomposites, we simultaneously achieve suppressed leakage current, enhanced breakdown strength and reduced conduction loss. It is attributed to the multiple blocking effects of the interfaces weakening electric fields. Through tuning the phase composition, we were able to manipulate the distribution of local electric field to obtain a higher breakdown strength, and suppress the ferroelectric loss. It is also found that besides the concomitant suppression of ferroelectric and conduction losses, the enhanced breakdown strength is of particular significance for achieving an ultrahigh charge-discharge efficiency in 18
our multilayered nanocomposites, where the remnant polarization remains unchanged at high electric fields.
5. Experimental Section 5.1. Materials P(VDF-HFP) and P(VDF-TrFE-CFE) were purchased from Arkema and used as received. All solvents were obtained from China National Chemicals Corporation Ltd. and used without further purification. 5.2. Fabrication of P(VDF-HFP)/P(VDF-TrFE-CFE) multilayered nanocomposites The non-equilibrium process was applied to prepare the multilayered nanocomposites. Briefly, P(VDF-HFP) and P(VDF-TrFE-CFE) powders were thoroughly dissolved in a mixed solvent of N,N-dimethylformamide (DMF) and acetone, respectively. The obtained solutions as electrospinning solutions were transferred into two syringes, respectively. A modified electrospinning process was performed to fabricate the multilayered fibrous mats. Take fabrication of a multilayered mat with four layers for example, the first layer, P(VDF-HFP) fiber, was electrospun at an applied field of 1.0 kV/cm with flow rate of 1.0 mL/h, followed by 2nd to 4th layers, i.e., P(VDF-TrFE-CFE) fiber layer, P(VDF-TrFE-CFE) fiber layer and P(VDF-HFP) fiber layer with same electrospinning conditions. The similar process was applied to fabricate the multilayered mats with eight and sixteen layers, respectively. However, the electrospinning time of each layer varied with the number of layer in different multilayered mats, which ensured the same thickness of final films. The fibrous mats were 19
then transformed into dense body by hot-pressing at 200 ℃ for 1 h under a pressure of 10 MPa. 5.3. The quenching process The as-pressed films were annealed at 200 ℃ for 15 min followed by quenching in water with different temperature, i.e., 0 ℃, 45 ℃ and 60 ℃, respectively. All of the samples had equal volume amount of P(VDF-HFP) and P(VDF-TrFE-CFE) and an overall film thickness of 15 μm. Characterization: The morphology of P(VDF-HFP) nanofibers, P(VDF-TrFE-CFE) nanofibers and multilayered nanocomposites were characterized with scanning electron microscopy (SEM, ZEISS MWRLIN compact). X-ray diffraction (XRD, Rigaku SmartLab) was performed for all samples. Dielectric constant were measured by employing a HP 4294A precision impedance analyzer (Agilent Technologies) at room temperature within the frequency range of 102 to 107 Hz at 1 Vrms. The Curie temperature of pure P(VDF-TrFE-CFE) films were tested by using the same impedance analyzer at 1 Hz and 1 Vrms within the temperature range of 30 ℃ to 100 ℃. Electric breakdown tests were performed with dielectric withstand voltage test system (Beijing Electro-mechanical Research Institute Supesvoltage 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.). The leakage current as a function of electric field were measured by the same ferroelectric test system. 5.4. Phase-field Simulations 20
In the phase-field model, two sets of order parameters were used in this model: a nonevolving variable r , with the value of 1 in copolymer matrix and 0 in terpolymer matrix, and a continuous phase field variable r,t which is both position- and time-dependent, where
r, t =1 represents the broken phase and r, t =0 represents the unbroken phase. It has to be noted that the electric inhomogeneity of the dielectric composite in this work is defined according to the different dielectric constants of the different matrix phase, the interface phase, and the electrically broken down phase. The
total
free
energy
F
of
the
dielectric
composite
can
be
written
as
1 2 F f sep g f elec 2 V (1) where the first term is phase separation energy density, the second term signifies the gradient energy density with the gradient energy coefficient g, and the last term represents the electrostatic energy density. The separation energy is expressed as
fsep =a 2 1
2
(2) where a is a coefficient related with the energy barrier of the phase transition. The electrostatic energy density f elec can be calculated as
21
1 1 f elec = 0 ij r Ei r E j r Ei r Pi S r 2 2 (3) where ij r is the position-dependent dielectric constant tensor, Ei r is the electric field component, Pi S r is the spontaneous polarization component. Therein the electric field distributions E r can be obtained by solving electrostatic Poisson equation 0 E P 0 using a Fourier spectral iterative perturbation method, which has been used
in previous work.[50, 51] The evolution of dielectric breakdown was simulated by solving a modified Allen-Cahn equation,
f r, t f elec L0 H f elec f critical sep g2 r, t t r, t r, t (4) where L0 is the kinetic coefficient related to the interface mobility, H f elec f critical is the Heaviside unit step function
H f
elec
f critical 0 and H f elec f critical =1 , f critical is
position-dependent material constant related to the maximal energy storage density of each constitute in the composite. In our phase field simulation, the characteristic length scale l0 g / a and similarly a characteristic time scale t 1 / L0a can be defined in terms of the material parameters g, a and L0. A grid size of 256x 256y 256z with grid spacing for x 1l0 and a time interval t 0.05t0 was used to yield a stable solution. Here, the initial parameters used in 22
this model, such as the breakdown strength and the dielectric constant of different pure polymers, are obtained from the experimental results. In order to reflect the interface effect, a 100 nm-thick interface region between adjacent layers with the dielectric permittivity of ~ 500 is introduced into our model. The coefficient in the separation energy term is given a value of 108 J/m3 to represent the barrier of the breakdown phase and the non-breakdown phase. A value of 10-10 J/m is used for the gradient energy coefficient to specify d0 1nm in the modeling.
Acknowledgment This work was supported by Basic Science Center Program of NSFC (Grant No. 51788104), the NSF of China (Grant No. 51625202, and 51572141), the National Key Research & Development Program (Grant No. 2017YFB0701603), the National Basic Research Program of China (Grant No. 2015CB654603) and Research Fund of Science and Technology in Shenzhen (JSGG20150331155519130).
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Fig 1. SEM images of (a) P(VDF-HFP) nanofibers, (b) P(VDF-TrFE-CFE) nanofibers, and (c) surface and (d) cross sections of Co/Ter multilayered nanocomposites with eight layers, respectively.(e) schematic illustration of the structures of the Co/Ter multilayered nanocomposites and the controls (pure copolymer and pure terpolymer), respectively.
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Fig 2. Comparison of breakdown strength for Co/Ter multilayered nanocomposites and pure terpolymer and copolymer films.
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Fig 3. (a) Efficiency of Co/Ter multilayered nanocomposites and pure terpolymer and copolymer films. (b) Comparison of efficiency and discharged energy density (at Eb) in this work and the related PVDF-based dielectric polymers reported in the previous literatures. 30
Fig 4. Breakdown mechanisms for Co/Ter multilayered nanocomposites. (a) Young’s modulus of Co/Ter multilayered nanocomposites quenched at 45 ℃. (b) Dielectric permittivity of pure copolymer and terpolymer quenched at different temperatures as a function of frequency. (c) Leakage current of Co/Ter multilayered nanocomposites quenched at 45 ℃ at different electric field. (d) Comparison of experimental and simulated breakdown strength of Co/Ter multilayered nanocomposites quenched at 45 ℃. (e) The distribution of local electric field along the direction of the applied field in the Co/Ter multilayered nanocomposites with 8L (top). Time-dependent 3D views of the breakdown evolution procedures for Co/Ter multilayered nanocomposites with 4L, 8L and 16L, respectively, by phase field simulation (bottom).
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Fig 5. XRD patterns for (a) pure terpolymer and (b) pure copolymer quenched at different temperatures. Energy loss (1-) for (c) pure terpolymer and (d) pure copolymer quenched at different temperatures as a function of electric field. (e) D-E loops for pure terpolymer quenched at 0 ℃, 45 ℃ and 60 ℃, respectively, at varied electric field.
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Fig 6. (a) Relationship of breakdown strength with efficiency in Co/Ter multilayered nanocomposites. (b) The remnant polarization of Co/Ter multilayered nanocomposites with 16L quenched at different temperatures as a function of electric field
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