Enhancing thermal stability of P(VDF-HFP) based nanocomposites with core-shell fillers for energy storage applications

Journal Pre-proof Enhancing thermal stability of P(VDF-HFP) based nanocomposites with core-shell fillers for energy storage applications Ling Zhou, Yufei Zhou, Yuchen Shi, Tianwei Chen, Tenghao Zou, Dongxiang Zhou, Qiuyun Fu PII:

S0266-3538(19)32359-0

DOI:

https://doi.org/10.1016/j.compscitech.2019.107934

Reference:

CSTE 107934

To appear in:

Composites Science and Technology

Received Date: 22 August 2019 Revised Date:

13 November 2019

Accepted Date: 21 November 2019

Please cite this article as: Zhou L, Zhou Y, Shi Y, Chen T, Zou T, Zhou D, Fu Q, Enhancing thermal stability of P(VDF-HFP) based nanocomposites with core-shell fillers for energy storage applications, Composites Science and Technology (2019), doi: https://doi.org/10.1016/j.compscitech.2019.107934. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier Ltd.

Enhancing thermal stability of P(VDF-HFP) based nanocomposites with core-shell fillers for energy storage applications Ling Zhoua,b, Yufei Zhoub, Yuchen Shib, Tianwei Chenb, Tenghao Zoub, Dongxiang Zhoub, Qiuyun Fub* a

Hubei Key Laboratory of Theory and Application of Advanced Materials Mechanics, School of Science, Wuhan University of Technology, Wuhan, 430070, China b

School of Optical and Electronic Information, Engineering Research Centre for

Functional Ceramics of the Ministry of Education, Huazhong University of Science and Technology, Wuhan 430074, P. R. China. Abstract Improving thermal stability of the ceramic-polymer based nano-composited electrostatic capacitors is the key element to their practical applications in harsh environment. In this paper, Fe3O4 @BaTiO3 particles with thermal conductive core and high-k shell were prepared and used as fillers to improve the thermal stability and high

temperature

energy

storage

density

of

poly-(vinylidene

fluoride-co-hexafluoropropylene) (P(VDF-HFP)) based composited films. Compared to the pure polymer, the films with 10 wt.% fillers showed enhanced temperature stability at the range of 40~120°C. The high-temperature (120°C) average breakdown strength reached at 1615 kV/cm with ~170% higher than the pure P(VDF-HFP) (970 kV/cm). The discharged energy storage density (120°C) was 1.68 J/cm3 (1750 kV/cm), which was enhanced up to 1150% comparing to that of pure P(VDF-HFP) (0.146 J/cm3 at 1300 kV/cm). The increased thermal conductivity and the Internal Barrier Layer Capacitor (IBLC) effects at the conductive-insulating interfaces contribute to the enhanced features. The results in the present work explored an effective way to prepare high temperature dielectric nanocomposites for harsh applications in energy storage field. Key words

* Corresponding author. E-mail addresses: E-mail: [email protected] (Fu Q.Y.)

Fe3O4@BaTiO3 core-shell fillers, Energy storage density, Multiple interfaces, Dielectric capacitors, Thermal stability. 1. Introduction Dielectric capacitors, which can convert a low-power, long-time input into a high-power, short-time output, are served as ideal energy storage devices and power pulse devices in electronic and electrical fields, such as hybrid electric vehicles, wind turbine generators, pulsed power sources and grid-tied photovoltaics

[1-4]

. The key

issue to prepare high energy storage density dielectric capacitors is to improve the dielectric constant (

) and breakdown field strength (Eb), while suppress the

dielectric loss (Tan δ). Ceramic - polymer based composites, which combine the high of ceramics and high Eb of polymers, are considered to be the most promising candidates and become a hot topic in the energy storage field [5-25]. However, due to the poor thermal conductivity (TC) and poor thermal stability of polymers, their practical applications are limited, especially in harsh environments [26-28]

. For example, the biaxially oriented polypropylene (BOPP), which is the best

commercial polymer dielectric, cannot match the requirements in power inverters of electric vehicles. The near-engine-temperature is above 120 °C

[29]

, but the operating

temperature of BOPP is below 105 °C [30]. In addition, in the high voltage applications, because of the heat accumulation in the dielectric materials, thermal breakdown is the main failure mode of these composites

[31]

. Therefore, improving thermal stability or

thermal conductivity, and maintaining the high

and low Tan δ of these

nano-composites are the key elements to their practical applications. To achieve both high TC and high

, many efforts have been made focusing on

fillers. Fillers can be classified to three categories, 1) high

fillers, such as BaTiO3,

0.5Ba(Zr0.2Ti0.8)O3-0.5(Ba0.7Ca0.3)TiO3 nanofibers (BZCT NFs) or with functional shells (BaTiO3@C, BZCT@SiO2)

[32-34]

, 2) high TC fillers, such as SiC, BN, BNNS

(BN nano sheets) or there composites [35-37] and 3) conductive particles or fibers such as Ni, Al, Ag, Fe3O4, graphene, carbon nanotubes and et al [38-43]. Surface modification, core-shell structures, fillers orientation with extra field, and multifunctional composited fillers were also reported to improve the thermal and electrical properties.

However, it is still a great challenge to satisfy the above requirements simultaneously. On the one hand, due to the low thermal conductivity and phonon scattering at the interface, embedding high-k fillers is difficult to improve the thermal conductivity of these composites [32]. On the other hand, due to the low polarization of high TC fillers (SiC, BN, BNNS), the dielectric constant of the composite system is limited

[36, 37]

.

Furthermore, dispersing nearly percolation threshold of conductive particles or fibers into polymers can significantly improve the

and TC. However, because of the

aggregation and the percolation effect, the electrical conductivity of the material becomes too high and the breakdown strength deteriorates rapidly, which finally limits the application under high voltage [38-43]. In view of the above contradiction, we attempted to use thermal conductive core and high-k shell micro/nanoparticles as fillers to solve the above problems. Semi-conductive Fe3O4 (FO) spheres (TC = 7 W/m.K [44]), which are usually used as fillers to enhance the TC of nanofluids [45-47], are selected as thermal conductive cores. High-k ceramic BaTiO3 nanoparticle shells are coating on the cores, and high-dielectric P(VDF-HFP) (

≈ 10) is selected as the polymer matrix to improve

the dielectric features of the core-shell-matrix composites. With this strategy, the TC of the three-phase composites with 10 wt.% fillers increases to 0.41 W/m.K (vs. 0.18 W/m.K

for

P(VDF-HFP)).

Furthermore,

the

composites

show

excellent

high-temperature stability. The discharged energy storage density Udis of the composites with 10 wt.% fillers remains stable at the range of 40~120°C. At 120°C, the Eb reaches at 1615 kV/cm, which is ~170% higher than that of neat P(VDF-HFP), and the Udis is 1.68 J/cm3 (1750 kV/cm), which is ~1150% higher than neat P(VDF-HFP) (0.146 J/cm3 at 1300 kV/cm), showing a great prospect in high-temperature energy storage applications. 2. Experimental section Materials and methods The preparation procedure is shown in figure 1. Firstly, a typical solvothermal method was used to prepare Fe3O4 submicron spheres [48]. The precursor solution was

containing 0.015 mol iron chloride (Sinopharm, China, the same below), 0.05 mol potassium acetate and 150 mL ethylene glycol. The solvothermal temperature and reaction time was 200 °C and 10 h, respectively. Then the precipitate was gathered and washed by ethanol for three times. After dried, the Fe3O4 submicron spheres were prepared. Then a self-assembling hydrolysis-hydrothermal method was used to coat the FO spheres with BTO shells [49]. A suspension containing 0.003 mol FO powders, 200 ml n-butanol, 10 ml deionized water and 0.5g cetyltrimethyl ammonium bromide (CTAB) was mixed by ultrasonic dispersion for 30 min. As the FO spheres were modified by cationic surfactant, numerous micro reactors containing FO spheres and hydrophilic group head were formed at the water/n-butanol interfaces. Then the TiO2 precursor (a solution of 0.007 mol tetra-n-butyl titanate (TBOT) and 50 ml n-butanol) was slowly dropped into the FO suspension. With mechanical stirring, TBOT hydrolyzed and uniform Ti(OH)4 shells self-assembly deposited on the FO spheres because of the electrostatic force. Mechanical stirring was maintained for another 12 h to complete the hydrolysis reaction. Then the precipitation was washed by absolute ethanol for three times and the dried powder was the core-shell FO@TiO2 particles. After a standard hydrothermal reaction with Ba(OH)2.8H2O (Ba/Ti = 1.3 in mole ratio) and 120 mL deionized water at 200 °C for 4h, the resultant core-shell FO@BTO composites were synthesized. To remove the BaCO3 impurities generated during the hydrothermal process, the powders were immersed in acetic acid (3 mol/L), washed by deionized water and then dried. Finally, the composited films were prepared by solution blending and tape casting method [50]. The mass fractions of core-shell powders were designed to be 0%, 2.5%, 5%, 10% according to our previous work. Certain amounts of Fe3O4@BaTiO3 core-shell powders were dispersed in 10 ml N, N-dimethylformamide (DMF) to form homogeneous suspensions by using ultrasound bath for 4 h. Then calculated amounts of P(VDF-HFP) (Solvay Plastics, Shanghai, China) were dissolved into the suspensions. After an overnight, the P(VDF-HFP) powder was completely dissolved

and the suspensions were in an ultrasonic bath for several hours. Then the composited films were prepared by tape casting, the thicknesses were controlled to about 15 µm. After dried at 80 °C, a thermal treatment (200 °C, 7 min and quenched in ice water) was performed. For electric measurement, the dried films were coated with Au electrodes (Φ2.5 mm) using a thermal evaporation deposition system. For thermal conductivity measurements, the films were cut into pieces and hot-pressed to pellets (Φ12.7 mm) at 200 °C under a pressure of 10 MPa for about 30 min. Then the pellets were spray-coated with a thin layer of graphite powder at both sides. Characterization The powder’s X-ray diffraction (XRD) pattern was identified by Bruker D8 ADVANCE X-ray diffractometer with Cu Kα= 1.5406 Å. The interior morphology and element mapping were detected by JEOL JEM-2100F transmission electron microscope (TEM) equipped with Oxford X-max 80TLE energy-dispersive X-ray spectroscopy (EDS). The surface morphology of the films was checked by a JSM-5610LV scanning electron microscope (SEM). Tensile tests were performed by a universal testing machine (INSTRON 5848 Micro Tester) at room temperature (RT). The films were stamped into the standard dog bone shape. The tested region was 50 mm (Length) × 10 mm (Width). The films were pulled under constant speed of 10 mm/min. The dielectric properties were measured by a WK6500B impedance analyzer at RT. The energy storage densities were calculated from electric displacement–electric field (D–E) loops. The D-E loops were measured by the Radiant Precision Ferroelectric Tester. The testing temperature was over the range from RT to 120°C. The thermal diffusion coefficient was detected by a NETZSCH LFA 467 laser-flash diffusivity instrument.

Figure 1. Schematic diagram of the preparation process.

3. Results and discussion

Figure 2. TEM images of (a) a coated FO sphere, (b) the HRTEM image of BTO nanocrystals shells, (c-g) EDS mapping images of a core–shell particle, (h) XRD patterns of FO@BTO core-shell powders, (i) surface SEM image, (j) the cross section SEM image and (h) photograph of the 10 wt.% composited film.

Figure 2(a) shows the TEM image of a coated FO particle. The core-shell particle was covered with a rough shell consisted of numerous nanocrystals. Figure 2(b) shows the higher resolution TEM (HRTEM) image of these BTO nanocrystals. The inter-planar spacing of (1 0 0) and (1 1 1) were 0.285 nm and 0.232 nm, respectively. The EDS element mapping showed in figure 2c-2g also proved the core-shell structure in further. Figure 2(h) shows the XRD patterns of the core-shell powders. The diffraction peaks and relative intensities confirmed that the composited powders were structured in typical perovskite (JCPDS #00-075-0213, space group: Pm-3m) and cubic spinel

(JCPDS #01-088-0315, space group: Fd-3m). No impurity such as BaCO3 was observed, indicating that the acetic acid wash procedure was efficient. Figure 2(i) shows the surface morphology SEM figure of the films with 10 wt.% fillers in micro scale. There are no micro defects and pinholes observed. Fillers were dispersed uniformly in the polymer matrix with few aggregates. To detect the interior morphology of the fillers and matrix, the film was initially frozen in liquid nitrogen and cut with a blade quickly. Then the cross section was sputter coated with a homogeneous platinum layer and detected by a SEM (Gemini300, ZEISS), as shown in figure 2(j). It can be seen that the particles were in capsuled with polymer matrix without obvious organic-inorganic interface separation. The hydroxylated BTO shells, which are simply washed by acetic acid, promote the tighter adhesion between fillers and polymers, thus fewer defects, better mechanical and dielectric properties of the nanocomposited films are expected

[51]

. The acetic acid washing procedure is also

proved to be a simpler and cheaper interface modification method as compared to using the expensive surfactant such as polydopamine and so on. The photograph in figure 2(k) shows that the corresponding composited film is flexible and uniform.

Figure 3. Tensile stress-strain curves of composited films.

The mechanical property, which is of great importance for device processing and manufacturing, is showed in figure 3 and presented in the tensile stress-strain curves. The tensile strain of all the films reaches up to 100%. The tensile strength of pure polymer is 17.4 MPa, which is in region of reference value (Solvay 21216 / 1001, the

tensile stress at yield is 15 - 25 MPa at 23°C). The composited films show obvious enhancements. The tensile strengths are 43.8, 33.9, and 38.8 MPa. The elasticity modulus are 274, 221, 226, 212 MPa for the pure P(VDF-HFP), 2.5 wt.%, 5 wt.%, and 10 wt.%, respectively. The fine tensile properties of composited films are originated from the uniform dispersion of fillers and excellent intermolecular force between polymers and fillers.

Figure 4. Frequency dependent behavior of the pure P(VDF-HFP) and nanocomposited films (a) εr, (b) tan δ and (c) electrical conductivity, the inset is the equivalent circuit for the composited films.

Then the frequency dependent dielectric properties of the films are showed in figure 4. As the filler contents increase, the εr increases significantly, from 9.9 to 16.0 at 1 kHz. Not only the high-k shells, but also the huge interfacial polarization

[52]

,

including FO-FO interfaces in the cores, FO-BTO core-shell interfaces, BTO-BTO interfaces in shells and BTO-P(VDF-HFP) inorganic-organic interfaces, all contribute to the enhancement of εr. At lower frequency region, the tan δ remains nearly stable between pure polymer and composited films, all the values are lower than 0.04 at 1 kHz. The suppressed loss tangent is attributed to two reasons: 1) Coating insulating BTO shells has effectively prevented the agglomeration and percolation of the conductive FO spheres; 2) hydroxylation modification of fillers has enhanced interfacial compatibility between the inorganic fillers and organic matrix. At higher frequency region (>1 MHz), the tan δ gentally increases with the filler contents. As conductivity indicates just a slight increase with filler concentration

shown in figure 4(c), the composited films remain insulating, thus we introduce a simpified capaciter model to interprete this phenomenon [53-54]. The typical equivalent circuit consisting of one R element (Re) and one parallel RC elements connected in represent surface electrode and

series is showed in figure 4(c). Here

( )

represent the resistances (capacitances) of insulating dielectrics, respectively. Thus the total complex impedance ( ∗

where (

∗

=

+

∗

) can be written as, (1)

1+

=2

is the angular frequency (

= ∗) is derived as, ∗

=

1+ + +

). Then the complex admittance

(2)

The conductance G, which is the real part of =

1+ (1 +

)

∗

, can be written as,

(3)

In the measured frequency range of 100 Hz ~ 30 MHz, the ≈

,

≪ 1, and

is the measured capacitance of films. Thus the tan δ can be

approximately as equation (4), !=

≈

1

+

(4)

Using this equation, the complex behavior of tangent loss can be interpreted. At lower frequency, the first term dominates. Though the the slight increase in conductivity in figure 4c) and

decreased (concluded from increased (seen from figure 4a)

with filler contents, the slight change in the denominator ( obvious change in

) could not cause an

!, thus the value remained stable between the different films.

At higher frequency, the second term dominates. Considered that the stable,

increased as the filler contents, thus the

fillers concentrations. Also, the

remained

! slightly increased with

! reached their maximum values at

=#

$ %&' &(

.

Figure 5. D-E loops measured at the range of 40 °C to 120 °C, (a) pure P(VDF-HFP), (b) 2.5 wt.%, (c) 5 wt.%, (d) 10 wt.%. The high temperature (120 °C) D-E loops measured from 800 kV/cm to the breakdown strength, (e) pure P(VDF-HFP), (f) 2.5 wt.%, (g) 5 wt.%, (h) 10 wt.%.

Figure 5a-5d shows the temperature dependent D–E loops measured at 1000 kV/cm. The testing frequency was 100 Hz and the temperature range was from 40 °C to 120 °C. For the pure P(VDF-HFP), as the temperature increases, the loops become fatter and fatter, showing a typical leakage current feature. It is suggested that the large conduction loss is related to the temperature-dependent conduction mechanisms. As the filler contents increases, the loops become thinner, especially for the 10 wt.% sample. There is only a slight increment in remnant polarization (Pr) as temperature increases, which means much lower conduction loss than the pure P(VDF-HFP). The potential mechanism will be discussed later. Then the high temperature (120 °C) D-E loop groups measured from 800 kV/cm to the break down strength (1300 kV/cm, 1550 kV/cm, 1650 kV/cm and 1750 kV/cm, respectively) are showed in figure 5e-5h. For the clearer presentation, the (Dmax - Pr) value as the function of testing temperature and testing electric field are showed in figure 6. From the figure 6(a), the values of all the samples remain stable at the range of 40 °C to 80 °C. However, at the range of 80 °C to 120 °C, the values sharply decrease from 1.48 µC/cm2 to 0.32 µC/cm2 for pure P(VDF-HFP) (78.4% drop) and from 1.55 µC/cm2 to 0.62 µC/cm2 for 2.5 wt.% (60.0% drop). For the sample of 5 wt.% and 10 wt.% filler, the values decrease gently, from 1.68 µC/cm2 to 1.28 µC/cm2 (23.8% drop) and 1.62 µC/cm2 to

1.43 µC/cm2 (11.7% drop), respectively. Figure 6(b) shows the (Dmax-Pr) value as the function of electric field at 120 °C. As the electric field increases, the values of pure P(VDF-HFP) decrease from 0.70 µC/cm2 (800 kV/cm) to 0.05µC/cm2 (1300 kV/cm), while the core-shell particles filled samples exhibit opposite tendency. The values of 10 wt.% film increase from 1.18µC/cm2 (800 kV/cm) to 2.27 µC/cm2 (1750 kV/cm). Thus the sample with 10 wt.% filler is expected to have larger energy storage density at 120 °C.

Figure 6 (a) The (Dmax - Pr) values as the function of testing temperature, the electric field is 1000 kV/cm. (b) The (Dmax - Pr) values, (c) discharged storage energy density (Udis) and (d) the efficiency (η) as the functions of electric field, the testing temperature is 120 °C.

From the D-E loops measured at 120°C, the Udis were calculated and showed in figure 6(c). The composited films exhibit obviously higher energy density than the neat polymer. The highest energy density is observed in 10 wt.% film and reaches at 1.68 J/cm3 (1750 kV/cm). The value is 1150% enhancement of the pure P(VDF-HFP) (0.146 J/cm3 at 1300 kV/cm). Figure 6(d) shows the charge−discharge efficiency (η) as the fuction of electric field at 120°C. For the samples with 5 wt.% and 10 wt.% fillers, though the η drops as the electric field increasing, they still exhibit enhanced high temperature stability than the pure P(VDF-HFP). For the pure P(VDF-HFP) and the 2.5 wt.%, the η are less than 25% during the whole testing electric field. For the samples with 10 wt.% fillers, the η

is neerly 80% at 800 kV/cm and maintains above 60% until 1200 kV/cm. Figure 7(a) and 7(b) indicate the Weibull distribution of composited films Eb at room temperature and high temperature. The distribution is described by the two parameters according to the equations as follow [7]. ) = ln(, )

(5)

= ln (− ln(1 − /( + 1))) (6) where Ei is the experimental Eb value of each specimen and the values should be sorted in ascending order. The subscripts i and n are the serial number and the total number of specimens, respectively. Then Xi and Yi are ploted and fitted using a linear renationship. The slope of the fitting line is called the shape parameter β, which is considered to be one of important paramater in the Weibull distribution. It is related to the scatter of the experimental data and higher value presents higher level of reliability. Then the abscissa of intersection of the fitting line and the line of Yi = 0 is the logarithms of the average Eb. Figure 7(a) and 7(b) present the average Eb values and the shape parameters β, too. All the β values are >1, indicating a reliability of the experimental data and the availability of the Weibull distribution. Compared to the pure P(VDF-HFP) film (β is 8.496 at RT and 6.0 at 120°C), the values of composited films are higher, indicating better mechanical and eletrical features in the composited films. A clearer comparison of RT and high temperature average Eb versus the filler content is shown in figure 7(c). From the column diagram, the RT average Eb (Eb_RT) remains nearly stable as reported previously [50]. For the high temperature Eb (Eb_HT), the values increase as the filler contents. At 120°C, the average Eb of pure P(VDF-HFP) is 970 kV/cm, while that of 10 wt.% increases to 1615 kV/cm. Compared the RT average Eb with high temperature average Eb within a group of films, the average Eb drops ((Eb_RT - Eb_HT)*100/ Eb_RT) were 57.4%, 38.2%, 32.4%, and 28.7% for the pure P(VDF-HFP), 2.5 wt.%, 5 wt.% and 10 wt.%, respectively. Both the energy storage density comparison and the breakdown strenth comparison indicate that the composited films are more stable at high temparature and have great potential in high temperature applications.

Figure 7. Weibull plots and Weibull breakdown strength (a) measured at RT and (b) measured at 120 °C, (c) column diagram of breakdown strength.

The enhancements in temperature stability are mainly originated from the higher thermal conductivity of Fe3O4@BaTiO3 core-shell fillers, which improves the total thermal diffusivity of the composited films and suppresses the heat accumulation. Therefore, the thermal conductivity is discussed theoretically and experimentally. The theoretical thermal conductivity of composite (/012 ) is calculated by the equation 7 [55]. /012 = (1 − ∅)/2 + ∅ /4 (7) Where ∅ is the filler volume, /2 and /4 are the thermal conductivity of polymer matrix and fillers, respectively. Concerning the core-shell fillers, equation 7 can be written as the equation 8: 6

/012 = 5 ∅ / (8) 7

Where i presents the core, shell and matrix. The thermal conductivity of Fe3O4 core, BaTiO3 shell and P(VDF-HFP) matrix are 7 W/m.K, 6.2 W/m.K and 0.18 W/m.K, respectively. The core-shell filler volume is 9.66% (10 wt.%), the diameter of Fe3O4 core is ~500 nm and the thickness of BaTiO3 shell is ~50 nm. Then the calculated thermal conductivity of the 10 wt.% composites is 0.81 W/m.K, which is 4 times higher than the neat P(VDF-HFP). For a clearer presentation, the behavior of heat conduction of pure P(VDF-HFP) and 10 wt.% composites were simulated. The simulation is based on finite element analysis method. The sample size is set to 10 mm×10 mm. The initial condition is room temperature (23°C), boundary condition is 200 °C with infinite constant heat

source. Figure 8 shows the temperature gradient diagram after the heat source is applied for 30 µs. It can be definitely seen that the thermal diffusion rate of 10 wt.% composited film is much higher than that of pure P(VDF-HFP).

Figure 8. The comparison of simulation analysis of heat conduction of (a) pure P(VDF-HFP) and (b) 10 wt.% composited film.

Then the thermal conductivity of the composites was measured further. However, only the sample with 10 wt.% fillers was successfully measured. The density (ρ) of the pellet was 1.883 g/cm3 (RT). The tested thermal diffusion coefficient (α) of 10 wt.% composites was 0.186 mm2/s (23°C). According to the simple mixing role and the reference specific heat values at 23°C (Fe3O4, ~0.62 J/g.K, BaTiO3, ~0.46 J/g.K

[56]

,

P(VDF-HFP), 1.2 J/g.K), the specific heats (c) of the composite films was calculated to 1.138 J/g.K. Using the equation λ(T) = α(T) ρ(T) c(T), the thermal conductivity (λ) of 10 wt.% composite pellets was 0.41 W/m.K. The value is only a half of the theoretical value (0.81 W/m.K), which mainly due to the huge interface thermal resistance between the pieces of films in the pellet. Combined with the high-temperature energy storage measurements, the in-plane thermal conductivity of single layer film is expected much higher than the pressed pellets. Besides, considering the lower conductive loss of composited films at the same temperature and the same electric field (see figure 5(a-d)), we supposed the conductive mechanism is different between pure polymer and composites. In the high-temperature region, Schottky emission and Poole-Frenkel emission

are the main conduction mechanisms in polymers

[27]

. Schottky emission is an

electrode-limited conduction mechanism, which is strongly depending on energy barrier at the metal electrode and dielectric interface. Poole-Frenkel emission is related to the trap energy level in the interior of polymers and is considered a bulk-limited type of conduction mechanisms. As the electric field is applied, charge carriers can overcome the metal-dielectric interface energy barrier and inject into the dielectrics when they get enough energy from thermal activation. In the dielectrics, the charge carriers will be trapped in trapping centers and detach from the trapping centers when they obtain enough energy to overcome the potential barrier, resulting to the leakage current and conductive loss. In our work, there are the same metal electrode-dielectric interfaces and measure conditions between the four group samples, thus the deference in Schottky emission might be ruled out. The increasing FO@BTO fillers contents contributes to the suppressed conductive loss. As the charge carriers inject to the dielectrics, the large conductive - insulating (FO-BTO) interfaces act as trapping centers and supply extra internal barriers in polymers. The extra internal barriers between conductive cores and insulating shells refer to the Internal Barrier Layer Capacitor (IBLC) Effects, which is mentioned in our previous work

[50]

. At room temperature region, the phenomenon

that Eb declined first and raised again as filler contents increasing is mainly because of the competition between polymers’ intrinsic high Eb features and the fillers’ IBLC effects. However, at high temperature region, the fillers’ IBLC effects dominated. As the filler contents increased, carrier entrapment was increasing while carrier detachment was decreasing, which finally resulted to the suppressed leakage loss and enhanced breakdown field strength. A brief comparison in electric and thermal features of PVDF serials based composites is shown in table 1. Compared to the bare high-K BTO fillers, conductive graphene or CNT and high TC BN or BNNS, the FO@BTO fillers filled composited films exhibited balanced and integrated features. Moreover, the FO@BTO fillers have significantly improved the high temperature energy storage density and breakdown

strength of P(VDF-HFP), which broaden the operating temperature range and may satisfy demanding applications.

Table 1. A comparison in thermal and electric features of PVDF serials based composited films fillers

BTO

matrix

εr & tanδ

Eb

Udis

TC(RT)

(1 kHz)

(kV/cm)

(J/cm3)

(W/m.K)

10 wt%

~15,<0.05

~1000

--

~0.27

50 wt%

~43,<0.05

~100

--

0.701

contents

P(VDF-HFP)

Ref

29

BT@C/SiC

P(VDF-HFP)

BT@C 50 wt.%,SiC 7.8%

1394, 0.9

--

--

0.78

30

SiC-CNT*

PVDF

CNT 1.5%

714, 0.49 @0.1Hz

--

--

~0.4

32

mBN**

PVDF

10 wt%

~1.7, <0.05

--

--

3.09

33

40 wt%

~2.6, <0.05

--

--

6.5

SMG***

PVDF

(16 wt%)/

83.8, 0.34

--

--

0.679

41

ZnO

PVDF

40 wt%

123.9, --

463

--

~0.6

57

BNNS****

P(VDF-HFP)

0.5 wt%

~34, 0.05

3250

5.6

~0.5

58

FO@BTO

P(VDF-HFP)

10 wt.%

16, 0.04

2350 (RT)

7.018

0.41

This work

1750 (120°C)

1.68

*SiC-CNT CNT-carbon nanotubes ***SMG- surface-modified graphene

** mBN - polydopamine modified hexagonal boron nitride ****BNNS - hexagonal boron nitride nanosheets

4. Conclusions In this paper, Fe3O4 @BaTiO3 particles with thermal conductive core and high-k shell were prepared and used as fillers to improve the thermal conductivity and dielectric constant of P(VDF-HFP) based composited films. The composited films, especially the sample with 10 wt.% fillers show great enhancements in εr, TC, Udis and high temperature stability to the pure P(VDF-HFP) films. The εr increases from 9.9 to 16.0 at 1 kHz, the TC increases from 0.18 to 0.41 W/m.K and the tan δ remains less than 0.04 (1 kHz). The Udis of 10 wt.% films maintained stable at the range of 40~120°C. The average high-temperature (120°C) Eb reached at 1615 kV/cm, which is 170% higher than the pure P(VDF-HFP) (970 kV/cm). The obtained maximum Udis is

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Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: