Open Access

Probing atomic-scale structure of dielectric ceramics with scanning transmission electron microscopy

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, P. R. China

Search for more papers by this author

Rui Wei

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, P. R. China

Search for more papers by this author

Jinquan Zeng

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, P. R. China

Search for more papers by this author

, and

Yuan-Hua Lin

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, P. R. China

E-mail Address: linyh@tsinghua.edu.cn

Corresponding author.

Search for more papers by this author

https://doi.org/10.1142/S2010135X23430014Cited by:0 (Source: Crossref)

Abstract

High performance dielectric capacitors are ubiquitous components in the modern electronics industry, owing to the highest power density, fastest charge–discharge rates, and long lifetime. However, the wide application of dielectric capacitors is limited owing to the low energy density. Over the past decades, multiscale structures of dielectric ceramics have been extensively explored and many exciting developments have been achieved. Despite the rapid development of energy storage properties, the atomic structure of dielectric materials is rarely investigated. In this paper, we present a brief overview of how scanning transmission electron microscopy (STEM) is used as a tool to elucidate the morphology, local structure heterogeneity, atomic resolution structure phase evolution and the correlation with energy storage properties, which provides a powerful tool for rational design and synergistic optimization.

Keywords:

1. Introduction

Dielectric capacitors play an important role in the currently available energy storage equipment and are a promising prospect for pulsed power systems.^1,2,3,4 They have the attractive properties of fastest charging–discharging rates (on the order of microseconds), the highest power density (on a megawatt scale) and the longest work lifetime.^5,6 However, its relatively low energy density limits its wider application.⁷ In order to find a more favorable compromise between energy density, lifetime, cost, power density and operating temperature range, multiscale structures of dielectric materials are extensively explored.^{8,9,10,11,12,13,14,15,16} In order to reveal the structural characteristics at multiscale, various characterization techniques are used, including but not limited to X-ray diffraction (XRD), neutron powder diffraction (NPD) and electron microscopy. XRD and NPD have been used to characterize the structure. The lattice parameters can be determined by XRD, and NPD has the powerful function of determining the details of light atoms, isotopes and magnetic atoms. By comparison, electron microscopy represents a significant advantage at the microscale.

Electron microscopy is an important characterization technique, which includes scanning electron microscopy (SEM), scanning tunneling microscopy (STM), transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM). SEM is mainly used to observe the surface morphology at 5–10nm using secondary electrons (SEs).¹⁷ STM is used to reveal the atomic level resolution image of the sample surface by using the change in tunneling current.¹⁸ TEM and STEM have incomparable advantages in detecting the structure and morphology of sample.¹⁹ In particular, STEM can directly observe the atomic-scale structure. For example, high angle annular dark field (HAADF) and annular bright field (ABF) STEM imaging can be used to observe the cation and anion sublattices, respectively.²⁰ This ability can offer unique insights to understand the relationship between atomic-scale structure and macroscopic properties. Therefore, STEM has been widely used in fields such as materials, biology, physics, and chemistry.

In this review, we mainly focus on STEM techniques to characterize the morphology, local structure heterogeneity, atomic resolution structure phase evolution of dielectric ceramics and their correlation with energy storage properties.

1.1. Basic knowledge of dielectric materials

Dielectric materials are mainly solid materials, which are divided into organic and inorganic materials. The structural property of the dielectric material determines its microscopic morphology and thus its characterization image. The STEM image with atomic resolution is more dependent on the properties and atomic structure of the dielectric material itself. Therefore, obtaining STEM images helps to obtain atomic-scale structures and provides guidance for the scientific and rational design of dielectric energy storage materials. From the perspective of molecular structure, inorganic dielectric materials have microcrystalline ionic structure and amorphous structure (such as ceramics, glass, mica, etc.). Organic dielectric materials are mainly composed of covalently bonded polymer structures, which can be divided into nonpolar (such as polypropylene, polystyrene, etc.) and polar (polyethylene terephthalate, etc.) according to the symmetry of the structure. The structure of both under the STEM has a great difference.

Organic dielectric materials have high breakdown field strength,²¹ However, organic dielectric materials usually have a low dielectric constant.²² So surface modification and size effect are needed to improve the material properties.^23,24,25 One strategy for solving this problem is the use of fillers with high dielectric constants, such as high dielectric ceramics.^23,26 Three commonly used dielectric polymers are linear polyethylene terephthalate (PET), bidirectional tensile polypropylene (BOPP) and polyvinylidene fluoride (PVDF).²⁷ The discharge density of PET and BOPP is about 5–6Jcm $^{- 3}$ $^{- 3}$ , the dielectric constant is 2–3, the dielectric loss is 0.0003–0.002, and the breakdown field strength is about 600–700MVm $^{- 1}$ $^{- 1}$ .²⁸ The dielectric constant of PVDF is about 10 and the breakdown strength is higher than that of PET and BOPP.²⁹ Due to the requirement of the wide operating temperature of the dielectric capacitor, inorganic dielectric materials have attracted more attention due to its high temperature endurance.

Inorganic dielectric energy storage materials can be divided into single-phase and multi-phase. Ferroelectric, relaxation ferroelectric and antiferroelectric ceramics have the advantages of high dielectric constant, good temperature stability and low leakage current.³⁰ Lead-free ceramics are the focus of dielectric ceramics research because of the biological toxicity of lead and lead oxides. The energy storage performance of inorganic oxide ceramics can be improved by designing the core-shell structure and the lamination structure.^8,31 The core-shell structure can be clearly observed by TEM. Xiang et al. used HAADF-STEM characterization techniques to reveal the evolution of the atomic structure of Au and Pd in the Au@Pd core-shell structure as a function of Au core size and Pd shell thickness.³² Liu et al.³³ showed that by coating silica evenly on the surface of Ba $_{0.3}$ $_{0.3}$ Sr $_{0.7}$ $_{0.7}$ TiO₃ nanoparticles to form a core-shell structure, and used STEM to characterize the film quality, the energy storage density can be increased to 1.52J/cm², and the energy storage efficiency can reach 82.2%. (Ba $_{0.7}$ $_{0.7}$ Ca $_{0.3}$ $_{0.3}$ TiO₃/BaZr $_{0.2}$ $_{0.2}$ Ti $_{0.8}$ $_{0.8}$ O₃) multilayer thin film structure designed by Sun and others,³⁴ the interface can effectively prevent electrical tree growth, so that the film of the energy storage density reached 52.4Jcm $^{- 3}$ $^{- 3}$ , energy storage efficiency of 72.3%. Multi-phase materials that can be used as energy storage materials include glass-ceramics materials. The dielectric constant of glass-ceramic is low, and the energy storage density is much lower than expected.³⁵ It is generally believed that the interfacial polarization between the crystalline and amorphous phases is the main reason for the reduction in energy storage density.⁴

The basic physical process of capacitor energy storage (charging) is that under the applied voltage V, the various polarization dipoles in the dielectric are transferred in a directional way and the accumulated charge Q is formed on the electrode surface to balance the applied voltage V, and finally reach the balance (Fig. 1). After charging is completed, the potential generated by the surface charge is equal to the applied voltage V, and the relationship between the capacitor C, the accumulated charge Q and the applied voltage V can be expressed by the formula: $C = Q ∕ V$ $C = Q ∕ V$ . The energy storage capacity of capacitor mainly refers to the amount of energy released under operating conditions (given electric field, temperature, etc.). The energy storage can be obtained by directly testing the discharge current or by testing the P–E relationship (Fig. 2). The hysteresis loop P–E method can easily and intuitively characterize the energy storage density and released energy density, while the discharge current test method can give the discharge speed, peak current, discharge current waveform curve and other information.^2,36

Fig. 1. Schematic of a dielectric in the presence of an electric field.

Fig. 2. Schematic P–E loop of dielectric energy storage.

1.2. Characterization of dielectric materials

With the increasing requirement of miniaturization and integration in the continuous development of modern society, how to improve the energy density of dielectric capacitors is one of the main research problems.⁷ And the energy storage density of the dielectric capacitors is largely influenced by the dielectric materials. Since the structure of the dielectric material determines the performance, it is necessary to analyze the structure and phase composition of dielectric materials accurately when studying dielectric materials. Dielectric materials are commonly characterized by XRD, SEM, TEM, STEM and so on.

Through XRD of the material and analysis of its diffraction pattern, XRD can obtain the phase composition of the sample and its structure. Besides, whether the sample is amorphous or crystalline can also be determined by XRD pattern. In addition, due to the difference in the radius of doped ions and matrix ions, cell expansion or contraction can be qualitatively analyzed by measuring the sample and standard figure 2 $θ$ $𝜃$ values.³⁷ The SEM scans the sample surface with a focused electron beam in a grating scanning mode. Because of the interaction between the electron beam and atoms at different depths in the sample, signals containing surface morphology and sample composition are generated. The SE is the most important imaging signal. The sample morphology can be obtained based on the number of SEs and the signal strength. The SEM image is usually combined with other analytical instruments to illustrate the microscopic morphology and composition. For example, an energy dispersive spectrometer (EDS) mounted on a SEM can be used to analyze the type and content of elements in the micro-components of materials.³⁸ Besides, TEM transmits accelerated and concentrated electron beams into a sample, and images after focusing and amplifying. Therefore, TEM observes the internal fine structure of the sample, such as crystal structure, morphology, etc.³⁸ Compared to SEM and TEM, STEM is a new analysis method that combines the principles and characteristics of scanning and ordinary transmission electron analysis. TEM is the object image created by focusing and amplifying the electron beam through the sample. At present, the commercial field emission STEM can not only obtain the high-resolution Z-contrast image and the energy loss spectrum of atomic resolution, but also all the other ordinary TEM (such as diffraction imaging, ordinary high-resolution phase contrast image, selective electron diffraction, converged-electron diffraction, micro-region composition analysis, etc.) can be completed in one experiment. In the recent 20 years, with the continuous development of electron microscopy, STEM with atomic resolution has become the most popular and widely used electron microscopic characterization and testing method.³⁹ In this paper, the STEM of dielectric ceramics will be reviewed.

1.3. Operational regimes for STEM-basic methods

As shown in Fig. 3, STEM uses electron beams excited by field emission electron guns, which are converged into atomic-scale electron beam spots after passing through a complex concentrating system. The resolution of STEM can be further improved by installing a spherical aberration corrector (Cs-Corr STEM) at the condenser location.⁴⁰ As a highly focused electron probe, point-by-point raster scanning is performed on thin samples under the control of scan coils. The interaction between the electron beam spot and the sample produces scattered electrons, which are collected, converted and imaged according to the scattering angle. In the working process of STEM, the mode of point-by-point scanning and imaging make each point scanned on the sample correspond to the generated image point one by one.

TEM collects transmitted electrons parallel to the electron beam for one-time imaging, while STEM uses the converging electron beam for point-by-point scanning imaging of the sample. For example, placing a SE detector above a thin sample can produce a SE image,^41,42,43,44 and placing a backscatter detector can produce a backscattered electron (BSE) image.⁴⁵ A ring detector placed under a thin sample can receive transmitted electrons scattered at a large angle. Such an image is called a HAADF image, which is sensitive to the square of the atomic number. This image is also called a Z-contrast image.⁴⁶ Electron Energy Loss Spectroscopy (EELS) can be used in combination with TEM to obtain the chemical information for qualitative and quantitative analysis. In EELS, after an electron beam with known kinetic energy is incident on the material to be tested, some electrons interact with the atom and give off inelastic scattering, losing part of the energy and causing small deflection of the path at random. The magnitude of the energy loss in this process is measured and analyzed by the electron spectrometer. By studying the energy loss distribution of inelastic scattered electrons, the space environment information of electrons in atoms can be obtained, so as to study a variety of physical and chemical properties of samples.

STEM has the characteristics of both scanning imaging and transmission analysis, and its instrument structure can be regarded as a combination of an SEM and a TEM.³⁸ It differs from the TEM mainly by the addition of a scanning attachment, and differs from the SEM by the location of the electron signal detector under the sample.

STEM images can also be obtained by SEM and TEM.⁴⁷ However, it is usually obtained by TEM because it requires the sample to be thin. In order to improve the resolution of ultrahigh resolution electron microscopy, it uses a sample table that can place the sample inside the objective lens. The sample table can only use thin samples a few millimeters thick, allowing STEM images to be obtained.

2. Morphology Test

Morphology testing is the basic function of the STEM. The STEM image is usually combined with corresponding EDS element mapping to illustrate how each element gets distributed. The elemental distribution of the KNN system was studied by Wang et al. using TEM and EDS.⁴⁸ For the (1–x)-(0.9(K $_{0.5}$ $_{0.5}$ Na $_{0.5}$ $_{0.5}$ )NbO₃–0.1Bi(Zn $_{2 ∕ 3}$ $_{2 ∕ 3}$ Nb $_{1 ∕ 3}$ )O₃)–xZnO sample, all elements are uniformly distributed in the $x = 0$ sample. With the increase of doping Zn content ( $x = 0.1$ ), the content of Zn in the grain interior gradually decreases due to the dynamics and low solubility in the solid state. When the doping Zn content is increased to $x = 0.125$ , the Zn content is enriched in the grain boundary. Compared to the $x = 0$ sample, the dispersion behavior of the $x = 0.1$ sample increases from 1.47 to 1.79. This is due to the inhomogeneity of charge and structure caused by the Zn $^{2 +}$ cations gradient distribution, which destroys the development of polar nano-regions (PNRs) and enhances the local random field. In 0.925(K $_{0.5}$ Na $_{0.5}$ )NbO₃–0.075Bi(Zn $_{2 ∕ 3}$ (Ta $_{0.5}$ Nb $_{0.5}$ ) $_{1 ∕ 3}$ )O₃ sample,⁹ all elements are uniformly distributed after 20h of high-energy ball milling. Although the crystallinity of the grain and the grain boundary is different, a chemically uniform microstructure is also achieved.

In thin film materials, the thickness of the sample can be more accurately obtained from cross-sectional STEM image.⁴⁹ In the 0.85BaTiO₃–0.15Bi(Mg $_{0.5}$ Zr $_{0.5}$ )O₃ (BT–BMZ) system, the thickness of BT–BMZ in BT–BMZ/Si (left), BTBMZ/HfO₂/Si (middle), and BT–BMZ/G/HfO₂/Si (right) is 340, 380 and 370nm, respectively, (Fig. 4). The corresponding EDS elemental mapping reveals the even distribution and high-quality interface, which would greatly benefit its performance in energy storage applications. Besides, the multilayer film can be clearly distinguished by the Z contrast of the image. Ba(Zr $_{0.15}$ Ti $_{0.85}$ )O₃ (BZT15) and Ba(Zr $_{0.35}$ Ti $_{0.65}$ )-O₃ (BZT35) is clearly revealed in epitaxially grown BZT15/BZT35 multilayers with repetition periods (N) of 6–12.⁵⁰ In the Ba $_{0.7}$ Ca $_{0.3}$ TiO₃–BaZr $_{0.2}$ Ti $_{0.8}$ O₃ (BCT–BZT) multilayers (Fig. 5), the multilayer structure can be clearly seen by contrasting the black used for BCT and the white used for BZT.³⁴ Each layer of thin film has good interface quality, which provides guarantee for improving electrical properties.

Fig. 4. (a) Cross-sectional STEM images of BT–BMZ/Si (left), BTBMZ/HfO₂/Si (middle), and BT–BMZ/G/HfO₂/Si (right). (b) Corresponding EDS mappings of BT–BMZ/HfO₂/Si.⁴⁹

Fig. 5. Cross-sectional STEM images of the three BCT/BZT multilayers with (a) $N = 8$ ; (b) $N = 4$ ; (c) $N = 2$ .³⁴

3. Local Structure Heterogeneity

Relaxor ferroelectrics (RFEs) and antiferroelectrics (RAFEs) with hysteresis-free polarization response are generally believed to be the desired materials for energy storage applications, which is tightly associated with their local structural heterogeneity. STEM is an effective method to detect the local symmetry of RFEs and RAFEs due to its high resolution. In this section, STEM is used as a tool to detect the local structure heterogeneity, and the correlation with energy storage properties is also discussed. It is divided into three parts: ion displacement, domain size, domain structure.

3.1. Ion displacement

The most intuitive way to observe local structure heterogeneity through STEM is ion displacement. Significant displacement of cations is observed when cations with different size, mass and electronegativity share the equivalent position. In typical cubic Bi₂Ti₂O₇, the lattice sites are high-symmetric. In Yang’s work (Fig. 6),⁵¹ through high-entropy design, compared with the system error of 0.8 pm, the average value of the cation displacement is about 3.9pm. At the same time, the displacement direction also presents a disordered distribution. Due to the existence of atomic disorder and lattice distortion, the average grain size gradually decreases, while the content of amorphous phase increases. Therefore, a higher resistivity and a larger breakdown strength are obtained through an effective high entropy strategy.

Fig. 6. (a) Magnitude and (b) direction of cation displacements in the high entropy film.⁵¹

Luo et al.⁵² provided the oxygen–oxygen distances and cation–cation distances to study the local heterogeneity after Ta incorporation (Fig. 7). The mean values of oxygen–oxygen distances are 394.2pm and 393.1pm for AN and ANT55, respectively. Compared to minor change in the oxygen–oxygen distance, the standard deviation of the distances changes significantly, which is 40pm for AN and 25pm for ANT55. In a similar way to the O–O distances, the A–A distances are reduced from 276.8pm to 276.2pm after 55% Ta incorporation. With the mixture of Nb/Ta, the standard deviation of the A–A distances increased significantly, which is 2.9pm for AN and 4.5pm for ANT55. All the data show that the incorporation of Ta increases the heterogeneity of the local structure in ANT55, which effectively stabilizes the relaxor AFEs. Thus, high $W_{rec}$ with $η$ of 90% can be achieved at the same time in ANT55.

3.2. Domain size

The energy storage properties of RFE are considered to be related to its domain size, which generally fall into two categories: microscale domains and nanoscale domains. In the solid solution (BiFeO₃) $_{1 - x}$ –(SrTiO₃)_x (denoted ad BFSTO) films, Pan et al.⁵³ believe that the STO incorporated in the BFO induces compositional and chemical disorder, leads microscale domain transform to nanoscale domains and STEM is used to validate this (Fig. 8). With increasing STO content, large-scale domains are gradually being transformed into nanoscale PNRs, while the distance between neighbouring PNRs also increases. Benefiting from the transformation from microscale domain to nanoscale domains, the remnant polarization ( $P_{r}$ ) is effectively suppressed. In addition, a large maximum polarization ( $P_{m}$ ) value is maintained, which is due to the large amount of highly polarized BFO, thus, inducing high energy density ( $W_{rec}$ ) and energy efficiency ( $η$ ).

Fig. 8. (a)–(d) HAADF–STEM images of the atomic-scale ferroelectric domain structure of the BFSTO films (scale bars: 10nm).⁵³

This transformation was also observed in Pan’s another work.⁵⁴ With the increase of Sm doping, the diameter of domain size decreases from $\sim 2$ –5nm to $\sim 1$ –2nm, and the domain volume fraction decreases from $\sim 54$ % to $\sim 15$ %. This domain scale transition is related to the strong local heterogeneity by Sm doping, which also observed by many researchers. For example, the domain size with the same polarization direction is decreased to $\sim 1$ –3nm through high-entropy design, which ultrahigh $η$ $\sim 90.8$ % is realized in high entropy sample.¹²

3.3. Domain structure

The generally accepted method of obtaining RFEs/RAFEs is to break the long-range microscale domain into the nanoscale domain by doping or forming solid solution. The drawback of this method is the reduction of polarization. A structural design strategy was proposed by Pan et al. through the judicious introduction of rhombohedral (R) and tetragonal (T) nanodomains, which coexist in a cubic (C) paraelectric matrix.⁵ In comparison to the RFEs with only the R or T nanodomains, polymorphic nanodomains with competitive free energies lead to a smoother domain-switching pathway. According to the projected shift of the lattice canter of the B site cation with respect to its four nearest neighbouring cations of the A site, the domain structure is determined, which are shown in Fig. 9. All nanodomains for the BFO–STO binary film exhibit R phase, while both R-phase and T-phase nanodomains are observed with BTO incorporation. Through this strategy, ferroelectric hysteresis is minimized while maintaining a high degree of polarization, leading to an overall improvement in the performance of the energy storage system.

Fig. 9. HAADF STEM images of (a)–(c) $x = 0.0$ and (d)–(g) $x = 0.3$ .⁵

4. Atomic Resolution Structure Phase Evolution

PNRs are considered to have high activity and response speed to external electric fields, low loss and high thermal breakdown strength.¹² Nanoscale PNRs can be characterized by high angle annular dark field STEM (HAADF–STEM) at the atomic resolution level. In Chen’s work,¹² in order to improve the overall energy storage performance of dielectric ceramics, it is proposed to introduce multiple ions to design “local polycrystalline distortion”. Chen et al.¹² observed R phase, T phase, orthorhombic (O) phase and C phase in K $_{0.2}$ Na $_{0.8}$ NbO₃ based high-entropy system. As shown in Fig. 10(a), the T phase, $R ∕ O$ phase and C phase can be confirmed by the arrows with the [001]c-direction, [011]c-direction and no direction. The transformations from T phase to $R ∕ O$ phase and then to C phase are enlarged in Fig. 10(b). Besides, the enlarged image (Fig. 10(c)) and schematic projection (Fig. 10(d)) of the unit cell along [100]c show the atomic-scale detail of the perovskite structure and the difference in the direction of the arrows between the T and $R ∕ O$ phases. In order to further distinguish between the R and O phases, polarization vector image along [110]c is performed by STEM. As shown in Fig. 10(c), both the R, T, O and C phases have observed. Furthermore, the polarization vectors (Fig. 10(f)), the magnified image (Fig. 10(g)) and the schematic projection (Fig. 10(h)) show the detailed perovskite structure of the high-entropy system along [110]c. The local structural heterogeneity is further confirmed by the random distribution of the polarization magnitude and polarization angle along [100]c (Figs. 10(i) and 10(j)). This is attributed to the fact that the large number of ions with a variety of ionic radii and valence states share the equivalent position.

Fig. 10. (a)–(g) Atomic-resolution HAADF STEM polarization vector image and (i) polarization magnitude mapping and (j) polarization angle mapping.¹²

In addition to the high entropy strategy, Chen’s other work⁵⁵ has proposed a strategy to improve the energy storage density in super-paraelectrics by constructing local diverse polarization. Both R phase, T phase, O phase and C phase were observed in 1/3BaTiO₃–1/3Bi $_{0.5}$ Na $_{0.5}$ TiO₃–1/3NaNbO₃ (BT–BNT–NN) ceramics (Figs. 11(a) and 11(b)). The gradual transition of the polarization vector from T toR to OtoC of the relaxed BT–BNT–NN sample is found in Fig. 11(c). Regions with the same direction of polarization form the PNRs, which measures $\approx$ 1–3nm. The size of the PNRs is smaller than the size reported in some literature, which can provide higher activity and response to an external electric field, thereby reducing losses and enhancing thermal breakdown strength, resulting in excellent energy storage properties with high energy density and high energy efficiency.

Fig. 11. (Color online) Atomic-resolution HAADF STEM polarization vector image along [100]c (a) and [110]c (b). (c) enlarged image of the marked area (orange rectangle) in (b).⁵⁵

5. Conclusion

This review focuses on the use of STEM as a tool to characterize the atomic-scale structure of dielectric ceramics. For decades, the research of dielectric ceramics has focused on energy storage properties and the relationship between the structure and properties at the microscale, sub-microscale and nanoscales. However, few studies focus on atomic scales, which is mainly due to the limitation of resolution. STEM not only plays an important role in elucidating the local structural heterogeneity through ion displacement, domain size and domain structure, but also proves its effectiveness in elucidating the atomic resolution structural phase evolution. STEM provides unique insights at the atomic scale and clarifies the relationship between atomic-scale structure and properties, which is difficult to detect by other technologies. A deep understanding of the relationship between atomic-scale structure and properties will create opportunities for obtaining dielectric ceramics with high energy density and efficiency.

Acknowledgments

This work was supported by the National Key Research Program of China under Grant No. 2021YFB3800601, the Basic Science Center Project of the National Natural Science Foundation of China (NSFC) under Grant No. 51788104. Min Zhang and Rui Wei contributed equally to this work.

References

1. Q. Li, F. Z. Yao, Y. Liu, G. Zhang, H. Wang and Q. Wang, High-temperature dielectric materials for electrical energy storage, Annu. Rev. Mater. Res. 48, 219 (2018). Crossref, Google Scholar
2. H. Palneedi, M. Peddigari, G. T. Hwang, D. Y. Jeong and J. Ryu, High-performance dielectric ceramic films for energy storage capacitors: Progress and outlook, Adv. Funct. Mater. 28, 1803665 (2018). Crossref, Google Scholar
3. Q. Li, K. Han, M. R. Gadinski, G. Zhang and Q. Wang, High energy and power density capacitors from solution-processed ternary ferroelectric polymer nanocomposites, Adv. Mater. 26, 6244 (2014). Crossref, Google Scholar
4. X. Hao, A review on the dielectric materials for high energy-storage application, J. Adv. Dielectr. 3, 1330001 (2013). Link, Google Scholar
5. H. Pan, F. Li, Y. Liu, Q. Zhang, M. Wang, S. Lan, Y. Zheng, J. Ma, L. Gu and Y. Shen, Ultrahigh-energy density lead-free dielectric films via polymorphic nanodomain design, Science 365, 578 (2019). Crossref, Google Scholar
6. Q. Li, L. Chen, M. R. Gadinski, S. Zhang, G. Zhang, H. U. Li, E. Iagodkine, A. Haque, L.-Q. Chen and T. N. Jackson, Flexible high-temperature dielectric materials from polymer nanocomposites, Nature 523, 576 (2015). Crossref, Google Scholar
7. Z. Yao, Z. Song, H. Hao, Z. Yu, M. Cao, S. Zhang, M. T. Lanagan and H. Liu, Homogeneous/inhomogeneous-structured dielectrics and their energy-storage performances, Adv. Mater. 29, 1601727 (2017). Crossref, Google Scholar
8. F.-Z. Yao, Q. Yuan, Q. Wang and H. Wang, Multiscale structural engineering of dielectric ceramics for energy storage applications: From bulk to thin films, Nanoscale 12, 17165 (2020). Crossref, Google Scholar
9. X. Wang, Y. Huan, P. Zhao, X. Liu, T. Wei, Q. Zhang and X. Wang, Optimizing the grain size and grain boundary morphology of (K, Na) NbO₃-based ceramics: Paving the way for ultrahigh energy storage capacitors, J. Materiomics 7, 780 (2021). Crossref, Google Scholar
10. Z. Cai, C. Zhu, H. Wang, P. Zhao, Y. Yu, L. Li and X. Wang, Giant dielectric breakdown strength together with ultrahigh energy density in ferroelectric bulk ceramics via layer-by-layer engineering, J. Mater. Chem. A 7, 17283 (2019). Crossref, Google Scholar
11. P. Zhao, L. Chen, L. Li and X. Wang, Ultrahigh energy density with excellent thermal stability in lead-free multilayer ceramic capacitors via composite strategy design, J. Mater. Chem. A 9, 25914 (2021). Crossref, Google Scholar
12. L. Chen, S. Deng, H. Liu, J. Wu, H. Qi and J. Chen, Giant energy-storage density with ultrahigh efficiency in lead-free relaxors via high-entropy design, Nat. Commun. 13, 3089 (2022). Crossref, Google Scholar
13. Z. Yang, H. Du, L. Jin and D. Poelman, High-performance lead-free bulk ceramics for electrical energy storage applications: Design strategies and challenges, J. Mater. Chem. A 9, 18026 (2021). Crossref, Google Scholar
14. Q. Hu, L. Jin, T. Wang, C. Li, Z. Xing and X. Wei, Dielectric and temperature stable energy storage properties of 0.88BaTiO₃–0.12Bi (Mg $_{1 ∕ 2}$ Ti $_{1 ∕ 2}$ )O₃ bulk ceramics, J. Alloys Compd. 640, 416 (2015). Crossref, Google Scholar
15. Z. Yang, H. Du, L. Jin, Q. Hu, H. Wang, Y. Li, J. Wang, F. Gao and S. Qu, Realizing high comprehensive energy storage performance in lead-free bulk ceramics via designing an unmatched temperature range, J. Mater. Chem. A 7, 27256 (2019). Crossref, Google Scholar
16. N. Qu, H. Du and X. Hao, A new strategy to realize high comprehensive energy storage properties in lead-free bulk ceramics, J. Mater. Chem. C 7, 7993 (2019). Crossref, Google Scholar
17. J. Castle and P. Zhdan, Characterization of surface topography by SEM and SFM: Problems and solutions, J. Phys. D 30, 722 (1997). Crossref, Google Scholar
18. D. Cui, J. M. MacLeod and F. Rosei, Probing functional self-assembled molecular architectures with solution/solid scanning tunnelling microscopy, Chem. Commun. 54, 10527 (2018). Crossref, Google Scholar
19. Y. X. Tong, Q. H. Zhang and L. Gu, Scanning transmission electron microscopy: A review of high angle annular dark field and annular bright field imaging and applications in lithium-ion batteries, Chin. Phys. B 27, 066107 (2018). Crossref, Google Scholar
20. H. Vahidi, K. Syed, H. Guo, X. Wang, J. L. Wardini, J. Martinez and W. J. Bowman, A review of grain boundary and heterointerface characterization in polycrystalline oxides by (scanning) transmission electron microscopy, Crystals 11, 878 (2021). Crossref, Google Scholar
21. X. Wei, H. Yan, T. Wang, Q. Hu, G. Viola, S. Grasso, Q. Jiang, L. Jin, Z. Xu and M. J. Reece, Reverse boundary layer capacitor model in glass/ceramic composites for energy storage applications, J. Appl. Phys. 113, 024103 (2013). Crossref, Google Scholar
22. T. Lewis, Interfaces: Nanometric dielectrics, J. Phys. D 38, 202 (2005). Crossref, Google Scholar
23. M. Arbatti, X. Shan and Z. Y. Cheng, Ceramic–polymer composites with high dielectric constant, Adv. Mater. 19, 1369 (2007). Crossref, Google Scholar
24. S. Komarneni, Nanocomposites, J. Mater. Chem. 2, 1219 (1992). Crossref, Google Scholar
25. J. Li, L. Zhang and S. Ducharme, Electric energy density of dielectric nanocomposites, Appl. Phys. Lett. 90, 132901 (2007). Crossref, Google Scholar
26. Y. Bai, Z.-Y. Cheng, V. Bharti, H. Xu and Q. Zhang, High-dielectric-constant ceramic-powder polymer composites, Appl. Phys. Lett. 76, 3804 (2000). Crossref, Google Scholar
27. H. Wu, F. Zhuo, H. Qiao, L. Kodumudi Venkataraman, M. Zheng, S. Wang, H. Huang, B. Li, X. Mao and Q. Zhang, Polymer-/ceramic-based dielectric composites for energy storage and conversion, Energy. Environ. Mater. 5, 486 (2022). Crossref, Google Scholar
28. W. Sarjeant, I. W. Clelland and R. A. Price, Capacitive components for power electronics, Proc. IEEE 89, 846 (2001). Crossref, Google Scholar
29. E. Baer and L. Zhu, 50th anniversary perspective: dielectric phenomena in polymers and multilayered dielectric films, Macromolecules 50, 2239 (2017). Crossref, Google Scholar
30. D. Zheng and R. Zuo, Enhanced energy storage properties in La (Mg $_{1 ∕ 2}$ Ti $_{1 ∕ 2}$ ) O₃-modified BiFeO₃-BaTiO₃ lead-free relaxor ferroelectric ceramics within a wide temperature range, J. Eur. Ceram. Soc. 37, 413 (2017). Crossref, Google Scholar
31. R. G. Chaudhuri and S. Paria, Core/shell nanoparticles: Classes, properties, synthesis mechanisms, characterization, and applications, Chem. Rev. 112, 2373 (2012). Crossref, Google Scholar
32. X. H. Zhang, Z. H. Sun, R. Jin, C. W. Zhu, C. L. Zhao, Y. Lin, Q. Q. Guan, L. Cao, H. W. Wang, S. Li, H. C. Yu, X. Y. Liu, L. L. Wang, S. Q. Wei, W. X. Li and J. L. Lu, Conjugated dual size effect of core-shell particles synergizes bimetallic catalysis, Nat. Commun. 14, 530 (2023). Crossref, Google Scholar
33. M. Liu, M. Cao, F. Zeng, J. Qi, H. Liu, H. Hao and Z. Yao, Fine-grained silica-coated barium strontium titanate ceramics with high energy storage, Ceram. Int. 44, 20239 (2018). Crossref, Google Scholar
34. Z. Sun, C. Ma, M. Liu, J. Cui, L. Lu, J. Lu, X. Lou, L. Jin, H. Wang and C. L. Jia, Ultrahigh energy storage performance of lead-free oxide multilayer film capacitors via interface engineering, Adv. Mater. 29, 1604427 (2017). Crossref, Google Scholar
35. E. P. Gorzkowski, M. J. Pan, B. A. Bender and C. C. Wu, Effect of additives on the crystallization kinetics of barium strontium titanate glass–ceramics, J. Am. Ceram. Soc. 91, 1065 (2008). Crossref, Google Scholar
36. F. Guan, Z. Yuan, E. W. Shu and L. Zhu, Fast discharge speed in poly (vinylidene fluoride) graft copolymer dielectric films achieved by confined ferroelectricity, Appl. Phys. Lett. 94, 052907 (2009). Crossref, Google Scholar
37. J. S. Hardy, J. W. Templeton, D. J. Edwards, Z. Lu and J. W. Stevenson, Lattice expansion of LSCF-6428 cathodes measured by in situ XRD during SOFC operation, J. Power Sources 198, 76 (2012). Crossref, Google Scholar
38. A. Ponce, S. Mejía-Rosales and M. José-Yacamán, Scanning transmission electron microscopy methods for the analysis of nanoparticles, in Nanoparticles in Biology and Medicine: Methods and Protocols, Methods in Molecular Biology, Vol. 906 (Humana Press, 2012), p. 453. Crossref, Google Scholar
39. R. R. Zamani, F. S. Hage, S. Lehmann, Q. M. Ramasse and K. A. Dick, Atomic-resolution spectrum imaging of semiconductor nanowires, Nano Lett. 18, 1557 (2017). Crossref, Google Scholar
40. H. Sawada, Y. Tanishiro, N. Ohashi, T. Tomita, F. Hosokawa, T. Kaneyama, Y. Kondo and K. Takayanagi, STEM imaging of 47-pm-separated atomic columns by a spherical aberration-corrected electron microscope with a 300-kV cold field emission gun, J. Electron. Microsc. 58, 357 (2009). Crossref, Google Scholar
41. R. M. Allen, Secondary electron imaging as an aid to STEM microanalysis, Ultramicroscopy 10, 237 (1982). Crossref, Google Scholar
42. G. Hembree, P. Crozier, J. Drucker, M. Krishnamurthy, J. Venables and J. Cowley, Biassed secondary electron imaging in a UHV-STEM, Ultramicroscopy 31, 111 (1989). Crossref, Google Scholar
43. J. Liu and J. Cowley, Imaging with high-angle scattered electrons and secondary electrons in the STEM, Ultramicroscopy 37, 50 (1991). Crossref, Google Scholar
44. H. Inada, D. Su, R. Egerton, M. Konno, L. Wu, J. Ciston, J. Wall and Y. Zhu, Atomic imaging using secondary electrons in a scanning transmission electron microscope: Experimental observations and possible mechanisms, Ultramicroscopy 111, 865 (2011). Crossref, Google Scholar
45. L. Gignac, C. Beslin, J. Gonsalves, F. Stellari and C.-C. Lin, High energy BSE/SE/STEM imaging of 8um thick semiconductor interconnects, Microsc. Microanal. 20, 8 (2014). Crossref, Google Scholar
46. S. J. Pennycook, Z-contrast STEM for materials science, Ultramicroscopy 30, 58 (1989). Crossref, Google Scholar
47. C. Sun, E. Müller, M. Meffert and D. Gerthsen, On the progress of scanning transmission electron microscopy (STEM) imaging in a scanning electron microscope, Microsc. Microanal. 24, 99 (2018). Crossref, Google Scholar
48. Y. Huan, T. Wei, X. Wang, X. Liu, P. Zhao and X. Wang, Achieving ultrahigh energy storage efficiency in local-composition gradient-structured ferroelectric ceramics, Chem. Eng. J. 425, 129506 (2021). Crossref, Google Scholar
49. J. Jin, C. Ma, G. Hu, L. Lu, Z. Liang, W. Zhao, Y. Liu, C. Ma and M. Liu, Enhancing energy storage performances in an ultra-wide temperature range via interface engineering and thermal management for silicon-integrated dielectric capacitors, Appl. Phys. Lett. 119, 242902 (2021). Crossref, Google Scholar
50. Q. Fan, M. Liu, C. Ma, L. Wang, S. Ren, L. Lu, X. Lou and C.-L. Jia, Significantly enhanced energy storage density with superior thermal stability by optimizing Ba(Zr $_{0.15}$ Ti $_{0.85}$ )O₃/Ba(Zr $_{0.35}$ Ti $_{0.65}$ )-O₃ multilayer structure, Nano Energy 51, 539 (2018). Crossref, Google Scholar
51. B. Yang, Y. Zhang, H. Pan, W. Si, Q. Zhang, Z. Shen, Y. Yu, S. Lan, F. Meng and Y. Liu, High-entropy enhanced capacitive energy storage, Nat. Mater. 21, 1074 (2022). Crossref, Google Scholar
52. N. Luo, K. Han, M. J. Cabral, X. Liao, S. Zhang, C. Liao, G. Zhang, X. Chen, Q. Feng and J.-F. Li, Constructing phase boundary in AgNbO₃ antiferroelectrics: Pathway simultaneously achieving high energy density and efficiency, Nat. Commun. 11, 4824 (2020). Crossref, Google Scholar
53. H. Pan, J. Ma, J. Ma, Q. Zhang, X. Liu, B. Guan, L. Gu, X. Zhang, Y.-J. Zhang and L. Li, Giant energy density and high efficiency achieved in bismuth ferrite-based film capacitors via domain engineering, Nat. Commun. 9, 1813 (2018). Crossref, Google Scholar
54. H. Pan, S. Lan, S. Xu, Q. Zhang, H. Yao, Y. Liu, F. Meng, E.-J. Guo, L. Gu and D. Yi, Ultrahigh energy storage in superparaelectric relaxor ferroelectrics, Science 374, 100 (2021). Crossref, Google Scholar
55. L. Chen, N. Wang, Z. Zhang, H. Yu, J. Wu, S. Deng, H. Liu, H. Qi and J. Chen, Local diverse polarization optimized comprehensive energy-storage performance in lead-free superparaelectrics, Adv. Mater. 34, 2205787 (2022). Crossref, Google Scholar

Vol. 14, No. 03

Metrics

Downloaded 439 times

History

Received 23 February 2023

Revised 8 April 2023

Accepted 14 April 2023

Published: 25 May 2023

Information

This is an Open Access article published by World Scientific Publishing Company. It is distributed under the terms of the Creative Commons Attribution 4.0 (CC BY) License which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

Keywords

PDF download