High energy density with ultrahigh discharging efficiency obtained in ceramic-polymer nanocomposites using a non-ferroelectric polar polymer as matrix

Journal Pre-proof High energy density with ultrahigh discharging efficiency obtained in ceramic-polymer nanocomposites using a non-ferroelectric polar polymer as matrix Xu Lu, Xiaowan Zou, Jialiang Shen, Lin Zhang, Li Jin, Z.-Y. Cheng PII:

S2211-2855(20)30108-7

DOI:

https://doi.org/10.1016/j.nanoen.2020.104551

Reference:

NANOEN 104551

To appear in:

Nano Energy

Received Date: 16 September 2019 Accepted Date: 29 January 2020

Please cite this article as: X. Lu, X. Zou, J. Shen, L. Zhang, L. Jin, Z.-Y. Cheng, High energy density with ultrahigh discharging efficiency obtained in ceramic-polymer nanocomposites using a non-ferroelectric polar polymer as matrix, Nano Energy, https://doi.org/10.1016/j.nanoen.2020.104551. 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. © 2020 Elsevier Ltd. All rights reserved.

High energy density with ultrahigh discharging efficiency obtained in ceramicpolymer nanocomposites using a non-ferroelectric polar polymer as matrix Xu Lua,*, Xiaowan Zoua, Jialiang Shenb, Lin Zhangc, Li Jinc, Z.-Y. Chengb,* a

Laboratory of Functional Films, School of Materials Science and Engineering, Xi’an

University of Technology, Xi'an 710048, China b c

Materials Research and Education Center, Auburn University, Auburn, AL 36849, USA

Electronic Materials Research Laboratory, Key Laboratory of the Ministry of Education &

International Center for Dielectric Research, School of Electronic Science and Engineering, Faculty of Electronic and Information Engineering, Xi'an Jiaotong University, Xi'an 710049, China *[email protected] (Xu Lu); [email protected] (Z.-Y. Cheng) Abstract To overcome the low charging efficiency of ceramic-polymer composites using ferroelectric polymers as matrix and to take the advantage in fabrication offered by the polar polymer, a polar but non-ferroelectric polymer – poly(methyl methacrylate) (PMMA) – was selected as the matrix in the development of high performance composites for energy storage. Ba0.5Sr0.5TiO3 (BST) nanoparticles was selected as the filler. Freestanding and flexible BST-PMMA ceramic-polymer nanocomposites with BST content up to 30 vol.% were fabricated in thickness of about 5 µm using spin-coating process and were systemically studied. Due to the strong interaction between the polar groups of PMMA and the hydroxyl groups on the surface of BST nanoparticles, the suspension of BST nanoparticles in PMMA solution exhibits excellent stability and, hence, the nanocomposite films have an excellent microstructure uniformity and compatibility between the BST nanoparticles and PMMA. All the BST-PMMA films exhibit an excellent frequency (100 Hz to 1 MHz) and temperature (-90 to 120 oC) stability in their dielectric properties with a high energy storage density of more than 11 J/cm3. Most importantly, an ultra-high discharging efficiency of almost 100% is obtained in all the nanocomposites. Keywords: Nanocomposite; PMMA; BST; Energy storage; Discharging efficiency

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1. Introduction With the rapid development of green energy, mobile devices/systems, and wearable electronics, devices for storing electric energy become more and more important in our daily life. Basically, there are two types of electric energy-storage devices: battery and capacitor. Batteries store energy into a chemical form and the output is the electric energy, while capacitors store the energy in an electrostatic field and release the electric energy as output. Normally batteries have a high energy density (102~103 J/cm3) but a low power density, while capacitors can provide a high power density but a low energy density (100~101 J/ cm3). Batteries have a fixed and low output voltage (~1 V), while the output voltage of capacitors can be easily adjusted over a very broad range. Additionally, some batteries are not rechargeable and some are rechargeable but with a very limited number of charging-discharging cycles, while all capacitors are rechargeable with an almost infinite number of charging-discharging cycles. That is, as the electric energy-storage device, capacitors are better than batteries in many aspects except the energy density. Therefore, it is highly desirable to increase the energy density of capacitors, which is determined by the energy density of the dielectrics used to fabricate the capacitors. The energy-storage density (U) of a dielectric can be expressed as: D

U = ∫ EdD

(1)

0

where E is the applied electric field, D is the electric displacement under the corresponding E, respectively. Since the D increases with increasing E monotonically, the maximum U that can be achieved in a dielectric is expressed as U =

2

∫

D( Eb )

0

EdD , where Eb (i.e., dielectric

strength or the electric breakdown field) is the maximum E that can be applied on the dielectric, D(Eb) is the D under an electric field of Eb [1.2]. For the dielectrics under the same E, a higher D can be achieved in a dielectric with a higher relative permittivity, εr, (i.e., dielectric constant). The discharging efficiency (η), defined as the ratio of discharged energy to the stored energy, is also a critical parameter for the capacitors/dielectrics used as an energy-storage device. A lower η means a larger amount of heat generated. Therefore, it is highly desirable for the energy-storage capacitors to use the dielectrics with a higher Eb,

εr and η. Both inorganic (i.e., ceramics) and organic (i.e., polymers) dielectrics have been extensively studied as energy-storage dielectrics [1-4]. In general, inorganics exhibit a high

εr (101~104) but a low Eb (101~102 kV/cm), while the organics exhibit a high Eb (103~104 kV/cm) but a low εr (100~101). Dielectrics can be also classified into polar and nonpolar dielectrics. Usually, the polar dielectrics exhibit a higher εr than the nonpolar dielectrics [2]. Among the polar dielectrics, ferroelectrics have been widely studied as energy-storage dielectrics due to their extremely high εr. For example, the εr of the nonpolar ceramics ranges from 100 up to 102, while the εr of ferroelectric ceramics can be up to 104; the εr of nonpolar polymers is usually smaller than 3, while the εr of ferroelectric polymers can be up to a few tens. It has to be mentioned that the nonpolar dielectrics usually exhibit a linear response: D=ε 0ε r E [2], where ε0 (=8.85×10−12 F/m) is the permittivity of free space. That is, when they are used to fabricate the capacitor, the capacitor will have a discharging efficiency of about 100% (η≈100%). The polar dielectrics, especially the ferroelectrics, exhibit a nonlinear response with a polarization-electric field (P-E) hysteresis loop, which 3

unfortunately results in a low η. To improve the U of the dielectrics, the efforts have been focused on the improvement of Eb of the inorganic dielectrics and the enhancement of the εr of polymers. For example, for inorganic dielectrics, ceramic-glass systems have been studied since the glass exhibits a high Eb [5-8]; fine-grained ceramics have been used since the Eb of a ceramic increases with decreasing grain size [9-11]; ceramic films have been investigated since the Eb of a ceramic increases with decreasing the thickness of the sample [12-15]. For polymers, polar polymers especially ferroelectric polymers have been studied since they exhibit a high εr [16-20]. To overcome the low η observed in the ferroelectrics, different approaches have been studied to reduce the P-E hysteresis, such as fine-grained ferroelectric ceramics [10,11,21], relaxor ferroelectric ceramics [22-24], antiferroelectric ceramics [25-27], quenched ferroelectric polymers [28-30], relaxor ferroelectric polymers [31-33], and polymer blends [34-37]. To fully utilize the advantages offered by ceramics (i.e., high εr) and polymers (i.e., high Eb, good flexibility and low processing temperature), ceramic-polymer composites, especially ceramic-polymer 0-3 composites, in which particles of ceramics with a high εr are embedded in a polymer matrix with a high Eb, have been extensively studied in last three decades for the development of high-performance dielectrics for energy-storage applications [1,20,38-44]. The properties of a composite are determined by the properties of its constituents, its connectivity, microstructure and microstructure uniformity [41]. Additionally, the compatibility between the ceramic particles and polymer matrix also plays a critical role on the composite’s performance, especially the Eb. From a device fabrication 4

point of view, it is highly desirable to have the composites with a uniform microstructure. Both the compatibility between the ceramic particle and polymer matrix and microstructure uniformity of the composite are strongly dependent on the process used to fabricate the composites. Ceramic-polymer 0-3 composites can be fabricated using either the dry or wet processes. For the dry processes, such as extrusion, the ceramic particles and polymer matrix are mixed mechanically [45-47]. Therefore, for the composites fabricated using a dry process, they usually do not have a good microstructure uniform and there is a poor compatibility between the ceramic particles and polymer matrix. In the wet process [1,20,38-44], the polymer matrix is first dissolved in a solvent to prepare the polymer solution, and then the ceramic particles are dispersed in the polymer solution to prepare a suspension. The suspension is used to prepare the composites through evaporating the solvent. So, in the wet process, the polymer chains could be completely released in the solvent and fully interact with the ceramic particles, leading a molecular level mixing of the polymer matrix with the ceramic particles. In these wet processes, a relatively stable suspension, which indicates there is a strong interaction between the ceramic particles and polymer solution, is the key to achieving a uniform microstructure in the fabricated composites. Regarding the ceramic particles used in the composites, there are two critical parameters: the size and εr. The smaller is the size of the ceramic particles; the better is the microstructure uniformity in the composites, especially along the thickness direction of the composite films [48]. Ferroelectric ceramics are generally selected as filler due to their high

εr. As mentioned above, ferroelectrics exhibit a strong polarization hysteresis, which can be

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weakened by reducing the size of the ferroelectric particles. Therefore, nanoparticles of various ferroelectric ceramics are favorable and, actually, widely used for the development of ceramic-polymer composites. It has to be mentioned that the ferroelectrics can exhibit a very weak or even no polarization hysteresis when it is used at temperatures higher than its Curie temperature (Tc). Regarding the polymer matrix used in the ceramic-polymer composites, it is predictable that nonpolar polymers, such as polypropylene (PP), are favorable since they exhibit an extremely high Eb and are linear dielectrics meaning a high η. Unfortunately, these polymers are extremely difficult to be processed using the wet processes due to the facts: 1) it is very difficult to prepare the solutions of these polymers; 2) even when the polymer solution can be prepared, it is extremely difficult to obtain a relative stable suspension of the ceramic particles in the polymer solution since there is not a strong interaction between the ceramic particles and polymer matrix. Due to these facts, the most popular polymers used in the investigation of ceramic-polymer composites are ferroelectric polymers, such as PVDF-based polymers [49-54]. The suspensions of ferroelectric ceramic particles in the solutions of PVDF-based polymers are relatively stable due to the strong interaction between the particles and the polar groups of the polymers. For example, it is experimentally found that the stability of the suspensions of ceramic nanoparticles in the solutions of PVDF-based polymers is strongly related to the density of hydroxyl on the surface of the ceramic nanoparticles and believed that the interaction between the hydroxyl surface of the ceramic nanoparticles and the polar groups in the PVDF-based polymers results in a great stability of the suspensions and hence an excellent compatibility between

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the polymer matrix and ceramic nanoparticles [55-59]. Unfortunately, these PVDF-based polymers exhibit a quite high polarization hysteresis (i.e., a low η), which would cause disastrous heating problem, especially under high E since the dielectrics usually have a low thermal conductivity. To reduce the polarization hysteresis, quenching has been widely used to reduce the ferroelectric phase in PVDF-based polymers [43,44,50,52,54]. It is predicable that the quenched polymers do not have a good stability. Additionally, the dielectric properties of these ferroelectric polymers are strong dependent on the temperature [16]. Again, all these are undesirable for the real applications. In a short summary, the recent research on the ceramic-polymer composites seems to be an undesirable direction due to the usage of ferroelectric constitutes, especially ferroelectric polymers as the matrix. To avoid the usage of ferroelectric polymers but, at the same time, to take the advantage of polar polymers offered, there is interest in using other polar but non-ferroelectric polymers as matrix for the development of ceramic-polymer composites. The non-ferroelectric polar polymers would exhibit a very weak or even zero polarization hysteresis so that the η should be almost 100%. Regarding the temperature dependence of the dielectric response, there are two critical temperatures for a polymer: glass transition temperature (Tg) for the amorphous and the melting temperature (Tm>Tg) for the polymer crystals. For a dielectric polymer, its properties are weakly dependent on temperature in the temperature range below its Tg. That is, polymers with a high Tg seems favorable for the development of high-performance energy-storage dielectrics. Based on the above discussions, one would expect that thermoplastic and polar polymer – poly(methyl methacrylate) (PMMA), also known as organic glass for its

7

commercial product, should be a great candidate for the polymer matrix used in the ceramic-polymer composites. PMMA has a Tm of about 160 °C, which is the same as the Tm of PP, and a glass transition temperature of about 120 °C, which is far higher than room temperature so that the dielectric properties of PMMA at temperatures around room temperature have a very weak temperature dependence [60]. Comparing to PP (εr=2.2), PMMA exhibits a higher εr (about 4) and a similar or even higher Eb (>7000 kV/cm). Comparing to PVDF-based polymers, PMMA not only demonstrates a much weaker temperature dependence in its properties but also a much higher Eb with no polarization hysteresis. At the same time, PMMA can be well dissolved in different solvents such as DMF, at room temperature. Additionally, PMMA has a good degree of compatibility with human tissue, and it has been widely used in surgeries and artificial organs. Therefore, one would expect that the ceramic-PMMA composites should be a great candidate for the wearable electronics and electronics used in implant devices. In this work, flexible and freestanding ceramic-polymer 0-3 nanocomposites are developed using PMMA as the polymer matrix. The nanocomposite films were fabricated using a wet process – spin coating. Ba0.5Sr0.5TiO3 (BST) nanoparticles with an average size of about 100 nm were selected as the ceramic filler. The BST has a Curie temperature of about -40 °C. That is, at temperatures above -40 oC, the BST would exhibit a very weak or even zero polarization hysteresis. It is well known that the properties of PMMA are strongly related to the molecular weight. PMMA with a molecular weight of up to 2×106 is of high transparency, good mechanical properties, good weather and chemical resistances, and excellent electrical insulation. Therefore, PMMA with a molecular weight of 2×106

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was selected for this work. The microstructures, thermal behaviors, dielectric properties and energy-storage performances of these BST-PMMA nanocomposite films were systemically investigated. Due to the utilization of BST and PMMA selected, it is experimentally found that all the BST-PMMA nanocomposite films show an extremely high discharging efficiency of almost 100%. 2. Experiments 2.1 Materials and chemicals All the materials and chemicals used in this work are commercially available, and they were used as received without further treatments. Poly(methyl methacrylate) (PMMA) pellets with an nominal molecular weight about 2×106 were obtained from Shuguang Co., Ltd. (Xi’an China). Ba0.5Sr0.5TiO3 (BST) nanoparticles with an average size of about 100 nm were purchased from nGimat Co. (Kentucky, USA). N, N-dimethyl formamide (DMF) with a chemical purity grade was purchased from Sinopharm Group Co., Ltd. (Shaanxi, China). Pre-cleaned glass slides (50 mm×50 mm×0.8 mm) were purchased from Guluoglass Co., Ltd. (Luoyang, China). 2.2 Fabrication of BST-PMMA nanocomposite films A spin-coating process has been successfully developed for the fabrication and investigation of ceramic-polymer nanocomposite films [58,59], and was employed here to fabricate the BST-PMMA nanocomposite films. Seven groups of nanocomposite films, with BST content as 0, 5, 10, 15, 20, 25, 30 vol.%, respectively, were fabricated. Those samples are abbreviated as xBST-PMMA, where x represents the BST content in volume percentage.

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For the fabrication of each group of films, the PMMA used was 1 g and the DMF solvent was 10 mL. The BST nanoparticles used were calculated from the compositions accordingly. In the calculations, the densities of the PMMA and BST were used as 1.18 g/cm3 and 4.91 g/cm3, respectively. The materials and chemicals used in the fabrication are shown in Table 1. Table 1 Materials and chemicals used in the fabrication of the BST-PMMA nanocomposite films xBST-PMMA DMF (mL) PMMA (g) BST (g)

x=0 10 1 0

x=5 10 1 0.22

x=10 10 1 0.46

x=15 10 1 0.74

x=20 10 1 1.04

x=25 10 1 1.39

x=30 10 1 1.78

Step 1 Preparation of PMMA solution: PMMA pellets of 1 g were dissolved in 10 mL DMF under magnetic stirring and ultrasonication. It is experimentally found that the dissolution of PMMA in DMF is very slow process. The complete and fully dissolution usually takes about 2-3 days. Therefore, two steps, swelling and dissolution, were clearly observed. It has to be mentioned that the dissolution of PMMA used here in DMF is quite slow. Interestingly, although the swollen PMMA in DMF looked quite sticky, the viscosity of the final solution was not very high, which indicates that the loosen coils of the swollen PMMA could be finally separated and fully dispersed into the solution. After the solution became clear and uniform, it was further stirred and untrasonicated for about 5 hours to make sure that the PMMA was completely and homogeneously dissolved in the DMF. Step 2 Preparation of BST suspensions in PMMA solution: Based on the data in Table 1, specific amount of BST nanoparticles was added in each PMMA solution and then fully 10

dispersed under magnetic stirring and ultrasonication for about 12 hours until homogeneous suspensions of BST nanoparticles in PMMA solution were obtained. It was experimentally found that the suspensions show great stability at room temperature. For example, without stirring and untrasonication, there was no visible phase separation or sediment after being stored for 24 hours. As discussed in the “Introduction” part, the stabile suspension is the key to fabricate the composites with a uniform microstructure. The excellent stability of the suspensions could be associated with the interaction between the polar groups in PMMA and the hydroxyl on the surface of the BST nanoparticles, indicating an excellent compatibility between the polymer matrix and the ceramic nanoparticles in the fabricated composites. Step 3 Fabrication of BST-PMMA nanocomposite films: The BST-PMMA nanocomposite films were fabricated using a spin-coating process, and the details are given below. The program of the spin-coater was set as 300 rpm for 30 seconds. About 1.5 mL of the suspension was dropped on a pre-cleaned glass slide and the spin-coater ran subsequently and immediately. After this, the glass substrate with the spin-coated film was placed in an oven at 80 °C for one hour to further dry the nanocomposite film. Then, all the solidified films were carefully stripped off from the glass slides in DI water. Finally, the free-standing BST-PMMA nanocomposite films were annealed at 140 °C for about 24 hour. 2.3 Characterizations To demonstrate the microstructure and dispersion condition of the BST nanoparticles in the PMMA matrix, the fracture morphologies of the BST-PMMA nanocomposite films were examined by a field emission scanning electron microscope (FE-SEM) (JSM-7000F JEOL, Tokyo, Japan). Before that, the nanocomposite films were fractured in liquid 11

nitrogen, and then the fracture surfaces were coated with a thin layer of gold to improve the conductivity. The X-ray diffraction (XRD) patterns of all the BST-PMMA nanocomposite films were recorded by an X-ray diffractometer (PANalytical, Cambridge, UK) in an 2θ range from 10° to 60°, during which the step size and scanning speed were set as 0.02° and 10°/min, respectively. Both the SEM observation and XRD analysis were done at room temperature. The thermal analysis was measured by a differential scanning calorimetry analyzer (DSC 250, TA Instruments, DE, US) in a temperature range from 0 °C and 200 °C with a heating/cooling rate of 10 °C/min in a nitrogen atmosphere. For the electrical measurements, the BST-PMMA nanocomposite films were coated with gold electrodes (3 mm in diameter) on both sides by a gold coater (JFC-1600, JEOL, Tokyo, Japan). The weak-field (500 mV) dielectric properties were measured by a precision impedance analyzer (4294A, Agilent, CA, USA) over a frequency range from 100 Hz to 1 MHz at selected temperatures in the range from -90 °C to 180 °C with a step of 10 °C. The temperature dependences of dielectric properties were measured by a precision impedance analyzer (4980, Agilent, CA, USA) in a temperature range from -90 °C to 180 °C under five selected frequencies (100 Hz, 1 kHz, 10 kHz, 100 kHz, and 1 MHz), during which the samples were placed in a heating & freezing stage (TP94, Linkam, Surrey, UK), and the heating/cooling rate was set as 10 °C/min. The polarization-electric field (P-E) loops were measured at 100 Hz and at room temperature using a ferroelectric testing system (TF analyzer 2000, aixACCT, Aachen, Germany). During the measurements, the electric filed was increased gradually with an increasing step of 100 kV/cm until the dielectric strength (Eb). 3. Results and discussions 12

The photos of the BST-PMMA nanocomposite films with different BST contents are presented in Fig. 1(a). The pure PMMA film is of good transparency and flexibility. The transparency of a ceramic-polymer composite film can be used as an experiential criterion for the microstructure uniformity. It can be observed that all the BST-PMMA nanocomposite films fabricated here are of uniform transparency, which indicates a good uniformity both in microstructure and film thickness.

Fig. 1(a) Photo of the BST-PMMA nanocomposite films with different BST contents; (b)(c) SEM images of the fracture surface of the 30BST-PMMA film.

It is well known that the transparency and flexibility of the ceramic-polymer composite films decreases with increasing ceramic content. However, it is experimentally found that all the BST-PMMA nanocomposite films, even with 30 vol.% of BST, are still translucent and flexible. To confirm the microstructure uniformity, the morphologies of the fracture surfaces of the BST-PMMA nanocomposite films were examined by FE-SEM. It is found 13

that all the BST-PMMA nanocomposite films are quite dense in microstructure, and the BST nanoparticles are uniformly dispersed in PMMA matrix. As an example, a fracture surface of the 30BST -PMMA film is given in Fig. 1(b)-(c). It is clearly observed that the films are of a uniform thickness of about 5 µm, and that the sphere-like BST nanoparticles are uniformly dispersed in the PMMA matrix without any observable agglomerations and gap between the BST nanoparticles and PMMA matrix. That is, all the nanocomposite films have a uniform microstructure, and there is an excellent compatibility between the BST nanoparticles and PMMA matrix. All these confirm that there is a strong interaction between the BST nanoparticles and the PMMA matrix, which is also evident by the good stability of the suspensions. It is believed that the dispersion condition and the interface compatibility are critical to obtain a high Eb and good mechanical properties for the ceramic-polymer composites. As has been well reported, some new methods, such as the hot-pressing [61-63] and surface treating [64-67], have been introduced recently to enhance the dispersion condition and the interface compatibility in ceramic-polymer composites. However, the fabrication processes of using those methods are quite complex and/or needs to be carefully controlled, which are not favorable for industrial production. The process used here for the fabrication of the BST-PMMA nanocomposite films is very simple and results in the nanocomposite films with high quality. All these also confirm that the spincoating process used here is a simple and effective wet process for the fabrication of ceramic-polymer composites. XRD patterns of all the BST-PMMA nanocomposite films are shown in Fig. 2. For the pure PMMA film, no crystalline peak is observed, but broad peak-like intensity outlines in

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the 2θ ranges of 10° to 20° and 20° to 30° can be observed, which indicate its amorphous nature in the structure. As expected, the intensity of the XRD peaks associated with the BST increases with increasing the BST content as shown in the inset plot of Fig. 2, which can be attributed to composition change, and it also indicates the formation of composites.

Fig. 2 XRD patterns of the BST-PMMA nanocomposite films with different BST contents. The inset is the plot of the intensity vs BST content for the representative peak of BST.

The DSC results obtained from all the nanocomposite films are shown in Fig. 3(a), where the samples were first heated from 0 °C to 200 °C, and then cooled down from 200 °C to 0 °C. The results from former reflects the real state of the composite films studied, while the results from the latter reflects the intrinsic behavior of the BST-PMMA system. It is well known that the glass-transition process can be reflected by a step-like change on the DSC curve. In Fig. 3(a), it can be seen that there is a clear step-like change on each cooling DSC curve of the BST-PMMA nanocomposite film, which indicates the glasstransition process of the PMMA. Clearly, the step-like change become weaker with 15

increasing the BST content, which is attributed to the composition change since BST does not have any phase transition in the measured temperature range [8]. The characteristic temperature of the step-like change on the cooling DSC curve is determined as the point of intersection of the bisector of the angle between the tangents with the step edge of DSC curve at higher temperature. Therefore, the cooling DSC curves indicate that the Tg of PMMA reported here is about 117~121 °C. The composition dependence of the characteristic temperature is shown in Fig. 3(b). One can find that the Tg of the PMMA in the nanocomposite films increases with increasing BST content. That is, the BST nanoparticles have a clear influence on the microstructure of PMMA, which indicates there is a strong interaction between the BST nanoparticles and PMMA matrix. This interaction limits the mobility of the polymer segments due to the immobility of BST particles and in turn results in the increase in the Tg of the PMMA in the composites. Interestingly, during the heating, the DSC curves show a weak peak-like behavior (i.e., weak endothermic peak). The peak-like DSC behavior during heating was also observed in other reports and was explained as the formation of PMMA crystals systems during the fabrication [68]. However, as mentioned above, there is no measurable crystallinity based on the XRD result of PMMA. Additionally, it is well known that the melting temperature (Tm) of a polymer should be much higher than its Tg, but the peak temperature obtained during the heating process is only less than 10 °C higher than the Tg obtained during the cooling process. When the peak temperature of the heating DSC curve is selected as the characteristic temperature of the heating process, it is found that the characteristic temperature (i.e., peak temperature) changes with the composition as shown in Fig. 3(c). Clearly, the peak temperature increases with increasing the BST content. One can find that the composition 16

dependence of the characteristic temperatures during the cooling and heating process shown in Fig 3(b)- 3(c), respectively, are almost identical except the temperature value. It is well known that there is a thermal hysteresis in the Tg determined using DSC during the heating and cooling process. That is, the Tg determined during the heating is higher than the Tg determined during the cooling. Therefore, based on the results shown in Fig. 3(b)-(c) and considering the thermal hysteresis, one can conclude that the above two characteristic temperatures should reflect the same nature – glass transition. That is, the peak-like DSC curve should also be the result of the glass transition process.

Fig. 3 (a) DSC curves of all the BST-PMMA nanocomposite films during the heating (bottom) and cooling (top) cycle; (b) Characteristic temperature of cooling DSC curve vs BST content; (c) Peak temperature of heating DSC curve vs BST content.

Besides the peak associated with the Tg, the DSC results during the heating also show a weak and wide peak at temperatures below the Tg. This weak and wide peak was also 17

reported and was explained due to the thermal history [69]. It should be mentioned that most of the polymers at a temperature much lower than their Tg are brittle. However, as mentioned above, the PMMA and all the BST-PMMA nanocomposite films reported here are pretty flexible at room temperature. This is another advantage of the PMMA for the development of high-performance ceramic-polymer composites.

Fig. 4 (a) εr, (b) tanδ, (c) ∆εr/εr(100 Hz), and (d) σac as a function of frequency for the BSTPMMA nanocomposite films at room temperature.

The dielectric frequency spectra of the BST-PMMA nanocomposite films at room temperature are shown in Fig. 4. It can be observed that for the pure PMMA film, the εr is about 4, which is significantly higher than the nonpolar polymer, such as PP (εr=2.2). That 18

is, besides the electronic and ionic polarizations, there are the contributions from dipole polarization to the dielectric properties of PMMA. It is also observed that both the εr and tanδ of PMMA slightly decreases with increasing frequency, which indicates that there is no relaxation process over the entire frequency range used here and can be attributed to the fact that room temperature is far below the Tg of PMMA. This is clearly different with the dielectric behavior of the widely studied PVDF-based polymers at room temperature. For example, PVDF-based polymers at room temperature exhibit a strong relaxation process with a relaxation frequency in the MHz range due to the glass-transition process, which results in a high tanδ, a strong frequency dependence of both the εr and tanδ, and an increase in the tanδ with increasing frequency [58,59,70,71]. All these behaviors observed in the PVDF-based polymers would result in a low discharging efficiency. Therefore, a much higher discharging efficiency is expected for the PMMA and the BST-PMMA composites reported here. Both the εr and tanδ of a ceramic-polymer 0-3 composite should be somewhere between the εr and tanδ of its two constituents [72]. It was reported that BST ceramics sintered from the same nanoparticles exhibit a very weak frequency dependence of the εr and a low tanδ over the frequency range from 100 Hz to 1 MHz at room temperature [8]. It can be seen in Fig. 4(a)-(b) that although the εr of the nanocomposite films increases with increasing BST content, the frequency dependence of the εr and tanδ of all the nanocomposite films are similar with that of the pure PMMA film. That is, the frequency dependences of the dielectric properties of the BST-PMMA nanocomposites are mainly determined by the PMMA. The increase in the εr of the BST-PMMA composite films with the BST content can be attributed to the high εr of the BST and the composition 19

change. Similar with the pure PMMA, the tanδ of all the BST-PMMA composite films is low and slightly decreases with increasing BST content as shown in Fig. 4(b). These can be attributed to the fact that the tanδ of BST is very small. The tanδ of the BST-PMMA nanocomposite films is exactly between that of its two constituents, which also indicates the high quality of the nanocomposite films fabricated in this work. Above results strongly indicate that the PMMA-based composites are great candidates for the development of high-performance composites for energy-storage applications. A frequency dispersion factor of the εr defined the ratios of the changes in εr at different frequencies (∆εr, i.e., ∆ε r = ε r (100 Hz ) − ε r ( f ) ) to the εr at 100 Hz [εr(100Hz)] has been used to evaluate the frequency dependence of the εr and the influences of ceramic content on the εr of the ceramic-polymer composites [70,71]. The frequency dispersion factors of all the BST-PMMA nanocomposite films are shown in Fig. 4(c). Interestingly, it is also found that this factor of the BST-PMMA nanocomposite films decreases with increasing BST factor, especially at high frequencies. For the pure PMMA and 30BST-PMMA, this factor at 1 kHz is 0.060 and 0.052, while at 1MHz is 0.159 to 0.143, respectively. This can be attributed to the weak frequency dependence of the εr of BST. In other words, the nanocomposites have a better frequency stability of the εr than the pure PMMA. It is very interesting to observe that the curvature in Fig. 4(c) is simply bending downward, while the curvature for PVDF-based composites is bending upward and relatively complex [70,71]. As well known, these phenomena are related to the relaxation processes of the polymer matrix, especially the glass-transition process. This is due to the fact that the Tg (~120 oC) of PMMA is much higher than room temperature, while the Tg (~-40 oC) of PVDF-based 20

polymers is much lower than room temperature. Again, the above results indicate that the PMMA-based composites are better than the PVDF-based composites in terms of the frequency stability of the dielectric properties at room temperature.

21

Fig. 5 (a) εr, (b) tanδ, and (c) σac as a function of the BST content for the BST-PMMA nanocomposite films at room temperature. The AC conductivity (σac) of a dielectric can be calculated using σac = 2π f ε0εr tan δ [2], where f is the frequency. The calculated σac of the BST-PMMA nanocomposite films is plotted in Fig. 4(d). It is observed that the σac of all the nanocomposite films is quite low and strongly dependent on frequency, and the relationship between the σac and f is quite linear in log-log scale, which indicates that the BST-PMMA composite films are very good electrical insulators with excellent capacitance behavior. As reported, the curvature of the σac spectra is also related to the relaxation processes of the polymer matrix [70,71]. It can be seen that the curvature of σac is quite straight as described by σac = 2π f ε0εr tan δ , which there is no relaxation process in the measured frequency range. To clearly show the influences of the BST content on the εr, tanδ, and σac of the BST-PMMA nanocomposite films, the results shown in Fig. 4 are summarized in Fig. 5. Clearly, the εr of the composite films increases with increasing BST content, while the tanδ is almost independent of the BST content. The σac increases with increasing BST content, which can be attributed to the increasing εr as expressed by σac = 2π f ε0εr tan δ . To study the influence of temperature on the dielectric frequency spectra of the PMMA, the εr and tanδ were examined in the frequency range from 100 Hz to 1 MHz and at different temperatures from -90 °C to 180 °C with a step of 10 °C and results at some selected temperatures are shown in Fig. 6 (a). Clearly, the frequency spectra of both the εr and tanδ are significantly affected by temperature due to the fact that the relaxation

22

processes associated with the molecular mobility are directly related to temperature. At temperatures above 60 °C, a clear relaxation process can be observed, and the relaxation frequency, at which the dielectric loss reaches its maximum, increases with the temperature. This relaxation process should be associated with the glass transition process. However, in the temperature range from 0 °C to 60 °C, there is not a clear relaxation process over the measured frequency range. For example, over the entire frequency range, the tanδ does not show any peak, although the εr shows some dependence on the frequency. One may assign this to the high frequency tail of the glass transition process. At low temperatures (< 0 °C), the tanδ is very low and the εr is low with a very weak frequency dispersion, which can be attributed to the fact that those temperatures are far below the Tg of PMMA.

23

Fig. 6 (a) Frequency spectra of the εr and tanδ of the pure PMMA film at different temperatures. (b) The temperature spectra of the εr and tanδ at selected frequencies of the pure PMMA film. To further study the dielectric behavior of the PMMA, the temperature spectra of both the εr and tanδ were characterized at five selected frequencies over a temperature range from -90 °C to 180 °C and the results are shown in Fig. 6(b). Clearly, there are two relaxation processes which can be recognized, especially at 100 Hz, and marked as “α” and “β”, respectively. It was reported that the α process is associated with the segmental motion in the amorphous polymer, i.e., the glass-transition process. The β process is located in the glassy 24

state, i.e., below the Tg of PMMA, which indicates that it should be a faster process than the

α-process. The β -process was reportedly induced by the partial rotation of −COOCH3 side groups [73]. It was reported that with increasing temperature, the β process of PMMA can become ‘‘locally coordinative’’, involving more molecular cooperative motions [73]. That is, the α process should involve more molecular cooperative motions than the β process, but the β process should have a higher relaxation frequency than the α process. Therefore, the dielectric dispersion observed in PMMA at temperatures from 0 °C to 60 °C observed in Fig. 6(a) can be assigned to the β process. At higher frequencies over 1 kHz or with increasing temperature above Tg of PMMA, owing to the coordinative motion between neighboring side chains, the β process merges with the β process so that only a single relaxation process (i.e., the glass transition process) is observed. In addition, above the Tg, a rapid increase in tanδ with an increasing temperature can be observed in the temperature spectrum measured at 100 Hz, which has been attributed to the relaxation process of the whole polymer chains at quite high temperatures [58,59,70,71]. As mentioned above, it is well known that a polymer at temperatures below its Tg should be brittle, but the PMMA and all the BST-PMMA composite films are flexible at room temperature. Based on the dielectric properties shown in Fig. 6, one may link the flexibility of PMMA at room temperature to the β process. The temperature spectra of the εr and tanδ of the BST-PMMA nanocomposite films are shown in Fig. 7(a)-(f). Clearly, these temperature spectra are similar with that of the pure PMMA film, which indicates that the temperature dependences of the dielectric properties of the composite films are mainly determined by the polymer matrix. It is quite interesting 25

that, besides the relaxations associated with the α process and α process of PMMA, there is another weak relaxation process that can be observed at about -40 °C for the nanocomposite films with a high BST content such as 25 and 30 vol.%, as marked with an arrow in Fig. 7(e)-(f). This new weak relaxation process was not observed in the pure PMMA. The BST used here has a ferroelectric-to-paraelectric phase transition at about 40 °C. Therefore, the observed new relaxation process in the nanocomposite films is due to the phase transition of the BST. It should be mentioned that the same BST nanoparticles had been used to fabricate the ceramic-polymer composites by using PVDF-based polymers as matrix [51,59]. In these composites, the relaxation process associated with the BST was never observed. This was caused by the polymer matrix used. PVDF-based polymers exhibit a strong dielectric dispersion in the above temperature range so that it is difficult to observe the dielectric relaxation of the ceramic filler in the dielectric properties of the composites. However, for the BST-PMMA nanocomposites reported here, the phase transition of BST and the glass transition of PMMA are located at different temperature ranges, and the dielectric properties of the polymer matrix exhibit almost no dispersion in the temperature range of the phase transition of BST, so the relaxation process of the BST can be observed in the dielectric properties of the composites. This provides a way to determine the phase transition temperature of nanoparticles.

26

Fig. 7 Temperature spectra of the εr and tanδ at selected frequencies for the BST-PMMA nanocomposite films. To clearly identify the influences of the BST content on the α process and β process of the PMMA matrix, the temperature spectra of the εr and tanδ measured at 100 Hz are plotted together, as shown in Fig. 8. It can be seen in Fig. 8(a) that although the εr of the nanocomposite films increases with the BST content, the εr exhibits a similar temperature dependence. For the tanδ peak associated with the α process, which reflects the Tα, it is found from the results shown in Fig. 8(b) that the Tα of the PMMA matrix in the nanocomposites increases with increasing BST content, which means that the BST nanoparticles could lower the local mobility so that the relaxation time of the glasstransition process becomes longer at the same temperature. In other words, the motion of PMMA segments can be constrained by the BST nanoparticles, which again indicates that there is a strong interaction between the BST nanoparticles and PMMA matrix. However, for the tanδ peak associated with the β process, as shown in Fig. 8(b), the peak position (i.e., 27

temperature) does not change with the BST content, while the peak intensity (i.e., the value of the tanδ at its maximum) gradually decreases with increasing BST content. The latter can be attributed to the composition change of the nanocomposites. The former indicates that the BST nanoparticles do not have an observable influence on the β process of the PMMA due to the fact that the β process is associated with –COOCH3 group that is very small. These results suggest that the addition of ceramic nanoparticles in polymer only have significant effects on the relaxation processes corresponding to the large-scale molecular motions, which could be observed in the dielectric property at high temperatures. Besides the α process and β process, there is another process observed at high temperatures, which is believed that this process is due to the motion of whole polymer chains. As can be seen in Fig. 8(b), at temperatures above the α process, the tanδ increases with increasing temperature. It should be mentioned that this was also observed in the PVDF-based ceramic-polymer composites [58,59,70,71]. For the PVDF-based ceramicpolymer composites, it is observed that the mobility of whole polymer chains increases initially with a small load of ceramic nanoparticles into PVDF-based polymers, but then decreases with further increasing ceramic content. However, the influence of the ceramic nanoparticles on the mobility of whole polymer chains in PMMA is different with that in the PVDF-based polymers. It can be seen in Fig. 8(b) that, at temperatures around 180 °C, the tanδ of the BST-PMMA nanocomposite films firstly decreases with 5 vol.% of BST, and then continuously increases with further increasing BST content, which is inverse to that of the PVDF-based composites. The above results indicate that the coordinative motion

28

between neighboring side chains of PMMA might be enhanced with the addition of BST nanoparticles, especially at high temperatures.

Fig. 8 Temperature spectra of the (a) εr and (b) tanδ at 100 Hz of the BST-PMMA nanocomposite films. The insect is the characteristic temperatures of the α process and the β process vs the BST content. To study the energy-storage performances of the BST-PMMA nanocomposite films, the P-E hysteresis loops were measured at 100 Hz and under various applied electric fields lower than the Eb, and the results are shown in Fig. 9(a)-(g). The Eb of the BST-PMMA nanocomposite films, representing the average of three distinct measurements, is plotted in Fig. 9(h). The discharged energy density (U) is shown in Fig. 7(i).

29

Fig. 9 (a)-(g) Polarization-electric field (P-E) hysteresis loops at 100 Hz for the BSTPMMA nanocomposite films; (h) Eb and (i) U for the BST-PMMA nanocomposite films. From the data shown in Fig. 9 (a), it is easy to conclude: 1) PMMA exhibits a very high Eb (>7100 kV/cm); 2) the P-E loop of the PMMA film is quite linear with very small hysteresis, which indicates that the PMMA film is a linear dielectric at room temperature and it has an excellent discharge efficiency (almost 100%). This is very different with the PVDF-based polymers, which exhibit a strong polarization hysteresis that results in a low discharging efficiency ~ 70-80% [44,52,54]. The P-E hysteresis loop of the ceramicpolymer composites is influenced by not only the polymer matrix but also the ceramic filler, 30

especially for the composites using linear polymers as matrix. From the results shown in Fig. 9(b)-(g), all the BST-PMMA nanocomposite films exhibit a quite small hysteresis as the pure PMMA, which can be attributed to not only the dielectric linearity of the PMMA matrix, but also the quite low hysteresis of the BST nanoparticles since the Cure temperature of the BST is much lower than the room temperature. Regarding the Eb, as shown in Fig. 9(h), the Eb of the nanocomposite films decreases with increasing BST content, which is commonly observed in the ceramic-polymer composites. Although the maximum U of the BST-PMMA nanocomposite films decreases with increasing BST content due to the decrease of Eb, the U under the same electric field increases with increasing BST content. Application Remarks: The results reported here indicate that using PMMA as matrix is a promising direction to develop high-performance ceramic-polymer composites. Therefore, if the Eb of the BST-PMMA nanocomposite films can be improved, an extremely high U could be obtained. As demonstrated in the research of ceramic-polymer composites using PVDF-based polymers as the matrix, the Eb of the ceramic-polymer composites can be significantly enhanced by different approaches [76-81]. All the results indicate that using PMMA as matrix confirms the speculation that using polar but non-ferroelectric polymers as the matrix should be a promising direction for the development of high-performance ceramic-polymer composites for energy-storage applications. 4. Conclusion In this work, BST nanoparticles and PMMA were selected as the ceramic filler and polymer

matrix,

respectively,

to

develop

high-performance

ceramic-polymer

nanocomposites for energy storage. Free-standing and flexible BST-PMMA nanocomposite 31

films with the BST content from 0 to 30 vol.% were fabricated in a thickness of about 5 µm using a spin-coating process. Comparing with nonpolar polymers (i.e., PP). PMMA has the same high Eb (>7000 kV/cm), but exhibits a significant higher εr. which results in a higher energy storage density (>11 J/cm3). Comparing with widely used ferroelectric polymers, PMMA exhibits a much higher Eb and a linear polarization response, which results in a high energy density with an ultrahigh charging efficiency (i.e., ~100%). In term of fabrication process, it is experimentally found that PMMA has all the advantages offered by the ferroelectric polymers. For example, due to the polar groups of PMMA, the suspension of BST nanoparticles in the PMMA solution is very stable, which results in an increase in the Tg of PMMA in the composites and a uniform microstructure and an excellent compatibility between the BST and PMMA in BST-PMMA nanocomposite films. The Tg of pure PMMA is about 120 oC, which results in an excellent temperature and frequency stability in its dielectric properties at temperatures below 100 oC. More interestingly, although the Tg is much higher than the room temperature, all the freestanding BST-PMMA nanocomposite films are flexible at room temperature. Besides the glass transition process (i.e., α process), another relaxation process (i.e., β process) due to the partial rotation of – COOCH3 group was also observed. It is believed that the β process is responsible for the flexibility observed in the nanocomposites at room temperature. For the nanocomposite films under the same electric field, the energy storage density increases with increasing BST content. All the nanocomposite films exhibit an ultrahigh charging efficiency (i.e. ~100%). Due to the usage of non-ferroelectric polar PMMA and the BST, all the nanocomposite films exhibit an excellent frequency and temperature stability, meaning the

32

nanocomposites can be used over a wide temperature range with a very short charging/discharging time. The results reported here demonstrate that using non-ferroelectric polar polymers as the matrix is a promising approach for the development of ceramic-polymer composites to achieve a high charging/discharging performance. This opens a new avenue for the development of high-performance ceramic-polymer composites as energy storage dielectrics. It has to be mentioned that the Eb observed in the BST-PMMA nanocomposite films reported here is significantly lower than that of non-ferroelectric polar PMMA. It is the same as the experimental results obtained in the ceramic-polymer composites using the ferroelectric polymers as the matrix. In recent year, a great deal of efforts has been given on the improvement of the Eb of the ceramic-polymer composites using ferroelectric polymers as the matrix. It is experimentally demonstrated that the Eb of these composites can be significantly improved by different approaches. These very same approaches can also used to improve the Eb of the ceramic-polymer composites using b=non-ferroelectric polar polymer as the matrix. That is, the Eb of the ceramic-polymer composites using nonferroelectric polar polymers as the matrix can be significantly improved. Therefore, an ultra high U with an ultrahigh charging/discharging efficiency would be observed in these composites.

Acknowledgement This research was financially supported by a doctoral starting fund of Xi'an University of Technology (No. 101-256211306) and an USDA NIFA Nanotechnology Grant (No. G00009848). 33

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Highlights: •

•

•

Non-ferroelectric polar polymer is introduced as the matrix for the development of flexible ceramic-polymer nanocomposite films with a high energy density and an ultrahigh discharging efficiency (~100%), which is demonstrated by using PMMA. Flexible and freestanding BST-PMMA nanocomposites with an uniform microstructure and an excellent compatibility between the PMMA matrix and BST nanoparticles were fabricated due to the selection of the polymer matrix. BST-PMMA nanocomposites exhibit an excellent temperature and frequency stability and can be used for medical applications.