Liang
Gao
a,
Jiaqi
Zhang
bc,
Linping
Wang
*a,
Dongyang
Zhang
a,
Fangzhou
Li
a,
Haoyu
Shen
a,
Ben-Lin
Hu
*a and
Run-Wei
Li
*a
aCAS Key Laboratory of Magnetic Materials and Devices, and Zhejiang Province Key Laboratory of Magnetic Materials and Application Technology, Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, 1219 West Zhongguan Road, Zhenhai District, Ningbo, P. R. China 315201. E-mail: wanglinping@nimte.ac.cn; hubenlin@nimte.ac.cn; runweili@nitme.ac.cn
bKey Laboratory of Engineering Dielectrics and Its Application, Ministry of Education, Harbin University of Science and Technology, Harbin 150080, P. R. China
cSchool of Electrical and Electronic Engineering, Harbin University of Science and Technology, Harbin 150080, P. R. China
First published on 27th September 2024
Polymer-based relaxor ferroelectrics with high dielectric constant are pivotal in cutting-edge electronic devices, power systems, and miniaturized pulsed electronics. The surge in flexible electronics technology has intensified the demand for elastic ferroelectric materials that exhibit excellent electrical properties and mechanical resilience, particularly for wearable devices and flexible displays. However, as an indispensable component, intrinsic elastomers featuring high dielectric constant and outstanding resilience specifically tailored for elastic energy storage remain undeveloped. Elastification of relaxor ferroelectric materials presents a promising strategy to obtain high-dielectric elastomers. In this study, we present a strain-insensitive, high elastic relaxor ferroelectric material prepared via peroxide crosslinking of a poly(vinylidene fluoride) (PVDF)-based copolymer at low temperature, which exhibits an intrinsic high dielectric constant (∼20 at 100 Hz) alongside remarkable thermal, chemical, and mechanical stability. This relaxor ferroelectric elastomer maintains a stable energy density (>8 J cm−3) and energy storage efficiency (>75%) under strains ranging from 0 to 80%. This strain-insensitive, high elastic relaxor ferroelectric elastomer holds significant potential for flexible electronic applications, offering superior performance in soft robotics, smart clothing, smart textiles, and electronic skin.
New conceptsThis paper presents a new concept of using peroxide cross-linking to prepare intrinsic high-elasticity relaxor ferroelectric materials for elastic energy storage. While relaxor ferroelectric materials for energy storage have been widely studied, current research primarily focuses on BTO-based composites, PMN-based materials, and PVDF-based terpolymers or quaterpolymers. These materials are typically brittle or plastic, failing to meet the elasticity demands of rapidly developing flexible electronics. Additionally, dielectric constant and breakdown field are crucial for energy storage materials. Some previous reports have attempted to achieve elasticity by blending elastomers with relaxor ferroelectrics, but this significantly increases the risk of electrical breakdown, making them unsuitable candidates. The newly developed elastic relaxor ferroelectric materials with high dielectric constants effectively address this issue, offering excellent elasticity and fatigue resistance. Our study provides a new approach to designing elastic energy storage materials, promising advancements in flexible electronics, and expanded applications in organic relaxor ferroelectric materials for elastic energy storage. |
High dielectric-constant materials are primarily used in capacitors and other energy storage devices.10–13 Among these, relaxor ferroelectric materials are emerging as highly promising candidates for next-generation high-performance capacitors. This is due to their unique electrical properties, including high dielectric permittivity, broad temperature stability, and enhanced energy storage capability, which arise from their disordered nanoscale ferroelectric domains that rapidly respond to an electric field, generating polarization for efficient energy storage and release.14–17 Energy storage in relaxor ferroelectric materials is achieved through polarization changes induced by an applied electric field. This process can be quantified using the formula
With the progression of flexible electronics technology, new requirements for the mechanical elasticity of materials have arisen.21–24 In applications such as wearable devices, flexible displays, and bioelectronic devices, materials must not only exhibit excellent electrical properties but also withstand various deformations without compromising functionality. This has driven researchers to investigate elastic dielectric materials capable of meeting these stringent demands.25,26 Our previous work introduced elasticity to relaxor ferroelectric materials through chemical crosslinking, successfully merging a high dielectric constant with exceptional mechanical elasticity. This innovation enables the material to maintain robust ferroelectric performance under complex deformation conditions, positioning elastic relaxor ferroelectric materials as strong contenders for dielectric energy storage applications.27,28 However, challenges remain for elastic energy storage materials, including low energy storage density and variations in energy storage efficiency under strain. Consequently, the development of specifically designed for elastic energy storage has yet to be realized successfully.
To address this challenge, we focused on poly(vinylidene fluoride-chlorotrifluoroethylene) (P(VDF-CTFE)).29,30 P(VDF-CTFE) is cost-effective and can be chemically modified in an alkaline environment to eliminate HCl forming P(VDF-CTFE-DB), which contains double bonds (DB) within the molecular chain.29 This transformation of the bulky CTFE monomer into –CC– bonds reduces spatial hindrance between molecular chains, facilitating the transition of P(VDF-CTFE-DB) from a paraelectric phase to a relaxor ferroelectric phase with enhanced ferroelectric characteristics. In previous work, crosslinking reactions typically required high temperatures,28,31,32 posing a significant risk for CMOS or organic electronic processes.33–35 Therefore, identifying a milder and more efficient crosslinking method to achieve elasticity at lower temperatures was crucial. We strategically chose benzoyl peroxide (BPO) for crosslinking due to its lower initiation temperature (∼80 °C) and high crosslinking efficiency.36,37 The high reactivity of the –CC– sites in the crosslinking material P(VDF-CTFE-DB) significantly reduces the crosslinking reaction temperature and enhances reaction efficiency. Additionally, triallyl isocyanurate (TAIC) was selected as a co-crosslinking agent to further boost crosslinking efficiency.38,39
In this work, we report an intrinsically elastic relaxor ferroelectric material prepared by peroxide crosslinking at a low crosslinking temperature (120 °C). This material exhibits a dielectric constant of 19.4 at 100 Hz at room temperature, alongside exceptional elasticity and thermal, chemical, and mechanical stability. Moreover, it maintains an energy storage efficiency exceeding 75% at frequencies above 1000 Hz and under strain levels ranging from 0 to 80%, demonstrating stable storage capability under different strains (>8 J cm−3 at 300 MV m−1). Compared to the widely studied BOPP which has an energy density of 2–4 J cm−3, this new material showcases outstanding energy density and strain-insensitive energy storage characteristics. This advancement positions it as a highly promising candidate for next-generation energy storage applications in flexible and wearable electronics.
The peroxide crosslinking mechanism is straightforward. During heating, the peroxide bonds in BPO undergo homolytic cleavage, yielding alkoxyl free radicals.38 These radicals subsequently attack the unsaturated CC bonds within the PVDF (polyvinylidene fluoride)-based polymer, generating polymer free radical intermediates. These intermediates then interact with the allylic bonds present in the tri-functional TAIC (triallyl isocyanurate), initiating the crosslinking reaction.40,41 This process forms a crosslinked network that enhances the thermal stability, mechanical properties, and chemical resistance of the polymer.
Crosslinked P(VDF-CTFE-DB) films exhibited notable resistance to common organic solvents such as cyclohexanone, acetone, dimethylformamide, and isophorone. After 15 days of immersion in these solvents, the crosslinked sample displayed swelling behavior with a gel content of approximately 90% (Fig. S3 and Table S1, ESI†).
Furthermore, in subsequent acid–base resistance tests, no significant volume changes were observed in the crosslinked P(VDF-CTFE-DB) films after two weeks of immersion in concentrated sulfuric acid and saturated sodium hydroxide aqueous solution, respectively (Fig. S4, ESI†). The crosslinked films also demonstrated excellent thermal stability, with a decomposition temperature (Td5%) exceeding 367 °C, as evidenced by the thermogravimetric analysis curve (Fig. S5, ESI†).
Fourier-transform infrared spectroscopy (FT-IR) and X-ray diffraction (XRD) were employed to verify the crystalline structure in both pristine and crosslinked P(VDF-CTFE-DB) films, as depicted in Fig. 1(B) and (C). Notably, the CC double bonds within the P(VDF-CTFE-DB) chain did not fully participate in the crosslinking process, as evidenced by the persistent presence of CC double bonds at 1720 cm−1 (Fig. 1(B)). The phase content analysis of P(VDF-CTFE-DB) before and after crosslinking was conducted with reference to previous studies.42,43 The identification of existing phases was based on the appearance of specific characteristic bands (listed in Table S2, ESI†), followed by calculations using specific equations depending on whether the sample contained α- and β-phases or α- and γ-phases. Further details are provided in Supplementary Note (ESI†). After crosslinking, the content of the α-phase (764 cm−1) increased from 14.7% to 47.48%, and the β-phase (840 cm−1) rose from nearly 0 to 52.52%. The X-ray diffraction (XRD) analysis (Fig. 1(C)) also clearly shows changes in the crystalline phases of the films before and after crosslinking. Although the crystallinity of the films decreased after crosslinking, leading to reduced diffraction peak intensity, the β-phase diffraction peak at 19.2° remains visible.
The elastic recovery of the crosslinked P(VDF-CTFE-DB) was evaluated through cyclic stress–strain measurements under strains ranging from 45 to 75% (5 cycles in Fig. 2(C) and over 3000 cycles under 50% strain as shown in Fig. 2(D)). Compared to the pristine P(VDF-CTFE-DB) samples (inset of Fig. 2(C)), the crosslinked films exhibited excellent elastic recovery. During the initial cycles, the recovery exceeded 95%, and after 600 cycles of stretching and releasing, the elastic recovery remained stable, exceeding 90% without significant decline. Compared to commercial fluororubber, the crosslinked P(VDF-CTFE-DB) film exhibited superior fatigue resistance and higher recovery ratios. The elasticity of crosslinked P(VDF-CTFE-DB) was attributed to entropy elasticity rather than energy elasticity, as estimated by the “force–temperature” curve of crosslinked P(VDF-CTFE-DB) under different strains (Fig. S8, ESI†). These results affirm that the intrinsic elasticity of polymer relaxor ferroelectrics can be obtained through peroxide crosslinking at a low temperature.
A capacitor-type device was fabricated to measure the ε–T curves of both pristine and crosslinked P(VDF-CTFE-DB) films. The ε–T curves of crosslinked P(VDF-CTFE-DB) are shown in Fig. 3(A). Compared to the pristine P(VDF-CTFE-DB) (Fig. S9, ESI†), the crosslinked film exhibits a broader ferroelectric-to-paraelectric (F–P) transition temperature range and a higher dielectric constant of 19.4 at 100 mHz under room temperature. Additionally, the Curie temperature (Tc) shifts to higher temperatures with increasing frequency, indicating the presence of relaxation behavior in the crosslinked P(VDF-CTFE-DB). The P–E loops of crosslinked P(VDF-CTFE-DB) were obtained using a sandwich structure device (Au/crosslinked-P(VDF-CTFE-DB)/Au/Si). As shown in Fig. 3(B), the loops exhibited low rectangularity, and appeared slender, which is a typical feature of a relaxor ferroelectric material. As the applied electric field increases, the initial hysteresis maintains its slender shape, characterized by a significantly large ratio between Pmax and Pr, resulting in lower rectangularity. At 1 kHz and 500 MV m−1, the Pmax and Pr of the crosslinked P(VDF-CTFE-DB) film are 8.2 and 1.05 μC cm−2, respectively. In comparison, the pristine P(VDF-CTFE-DB) exhibits Pmax and Pr values of 3.32 and 0.52 μC cm−2 at 280 MV m−1 and 1 kHz, respectively (Fig. S10A, ESI†). The coercive field (Ec) of crosslinked P(VDF-CTFE-DB) is approximately 47 MV m−1. The Pr of the crosslinked P(VDF-CTFE-DB) increases from 0.46 to 2.13 μC cm−2 across test frequencies ranging from 10 kHz to 100 Hz (Fig. 3(C)). Compared to pristine P(VDF-CTFE-DB) (Fig. S10B, ESI†), the crosslinked P(VDF-CTFE-DB) film exhibits more pronounced relaxor behavior due to the more condensed network introduced by peroxide crosslinking, which further breaks down the crystalline domains.
PFM was used to evaluate the piezoelectric properties of the crosslinked P(VDF-CTFE-DB) films. This technique enabled the visualization of local polarization domains and provided quantitative data on the piezoelectric coefficients. The crosslinked P(VDF-CTFE-DB) films exhibited enhanced piezoelectric responses compared to the pristine samples (Fig. S11, ESI†), confirming the improved ferroelectric relaxor behavior (Fig. 3(D)–(F)). The hysteresis and butterfly curves obtained from a single scan demonstrate the complete switching of ferroelectric domains under the influence of an electric field (Fig. 3(D)).
By applying a +10 V bias to a 5 × 5 μm2 region, followed by a −10 V bias to a 3 × 3 μm2 region, and finally a +10 V bias to a 1 × 1 μm2 region, phase and amplitude maps of a “box in box in box” pattern were obtained (Fig. 3(E) and (F)). This pattern suggests that the polarity of the ferroelectric domains in the crystalline structure can be reversibly switched by an applied field in a defined area rather than just at a single point. Additionally, PFM revealed a piezoelectric coefficient of 11.57 pm V−1 (Fig. S12, ESI†), indicating well piezoelectric performance. Relaxor ferroelectricity response and energy storage characteristics of crosslinked P(VDF-CTFE-DB) under strain.
To evaluate the relaxor ferroelectric response and energy storage properties of crosslinked-P(VDF-CTFE-DB) films under strain, we fabricated a fully elastic capacitive ferroelectric device as shown in Fig. 4(A). Flexible interdigitated electrodes were prepared using liquid metal gallium mixed with gallium oxide, with PDMS serving as the elastic substrate, and crosslinked P(VDF-CTFE-DB) films acting as the intermediate dielectric layer. The elastic device was affixed in a custom-made single-shaft tensile clamp and gradually stretched to 80% strain, as depicted in Fig. 4(B). The test results are summarized in Fig. 4(C)–(H). The P–E loops of this fully elastic device (Fig. S13, ESI† and Fig. 4(C)) without strain are comparable to those of the rigid device with Au electrodes. Additionally, without strain, as the frequency increases (from 100 Hz to 10 kHz), the Pmax and Pr of the crosslinked film slightly decrease. The energy storage density and efficiency were calculated at different frequencies, showing that the energy storage efficiency increases from 32.9% at 100 Hz to 90.7% at 10 kHz, while the energy storage density shows a minor decline, from 7.6 to 6.0 J cm−3 at 300 MV m−1. These results indicate good energy storage density and efficiency at frequencies of 1000 Hz and above (Fig. 4(D) and (E)).
The Pmax and Pr remain almost constant throughout the stretching process (Fig. 4(F), (G) and Fig. S14–S21, ESI†), which forms the basis for its strain insensitivity as an energy storage material. The energy storage density and efficiency were calculated under different strains (0–80%) at 330 MV m−1, and the energy storage efficiency remained above 75%, maintaining a high energy storage density (>8.0 J cm−3) (Fig. 4(H)). Traditional PVDF and its copolymers usually exhibit changes in crystal structure under tensile strain.44,45 However, due to the stress dispersion provided by the crosslinked network, the crystalline structure remains stable during stretching, resulting in consistent electrical performance.
To ensure the reliability of energy storage, we conducted electrical breakdown tests on the films before and after crosslinking. Aluminum electrodes with a diameter of 3 mm were deposited on the surface of the 10 μm thick films, and DC voltage was used for the electrical breakdown tests, as shown in Fig. S22 (ESI†). The breakdown strength of the crosslinked films significantly improved, with the pristine P(VDF-CTFE-DB) film exhibiting a breakdown field of 202 MV m−1 at a 63.2% breakdown probability, and the crosslinked P(VDF-CTFE-DB) film showing an increased breakdown field of 255 MV m−1.
Compared to the breakdown field observed during P–E loop testing, the electric breakdown tests on the 10 μm thick films likely include more film defects and consider factors such as the edge effects of the electric field due to the larger electrode size, which generally results in a lower measured breakdown field. Consequently, we calculated the energy storage density and efficiency under different strains at a test field of 230 MV m−1 (Fig. S23, ESI†). The energy storage efficiency consistently remained above 80%, with energy storage density always exceeding 4.5 J cm−3. This demonstrates the reliability of our testing and calculation of energy storage efficiency and density are reliable when miniaturizing the device.
Footnote |
† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4mh00998c |
This journal is © The Royal Society of Chemistry 2024 |