Electroceramics for High-Energy Density Capacitors: Current Status and Future Perspectives

Abstract

Materials exhibiting high energy/power density are currently needed to meet the growing demand of portable electronics, electric vehicles and large-scale energy storage devices. The highest energy densities are achieved for fuel cells, batteries, and supercapacitors, but conventional dielectric capacitors are receiving increased attention for pulsed power applications due to their high power density and their fast charge–discharge speed. The key to high energy density in dielectric capacitors is a large maximum but small remanent (zero in the case of linear dielectrics) polarization and a high electric breakdown strength. Polymer dielectric capacitors offer high power/energy density for applications at room temperature, but above 100 °C they are unreliable and suffer from dielectric breakdown. For high-temperature applications, therefore, dielectric ceramics are the only feasible alternative. Lead-based ceramics such as La-doped lead zirconate titanate exhibit good energy storage properties, but their toxicity raises concern over their use in consumer applications, where capacitors are exclusively lead free. Lead-free compositions with superior power density are thus required. In this paper, we introduce the fundamental principles of energy storage in dielectrics. We discuss key factors to improve energy storage properties such as the control of local structure, phase assemblage, dielectric layer thickness, microstructure, conductivity, and electrical homogeneity through the choice of base systems, dopants, and alloying additions, followed by a comprehensive review of the state-of-the-art. Finally, we comment on the future requirements for new materials in high power/energy density capacitor applications.

1. Introduction

To limit global warming to <1.50 °C, as set out in the Paris agreement, carbon dioxide emissions need to decrease ∼45% by 2030 and reach net-zero by 2050.^1,2 Technologies based on renewable resources such as sun, wind, and tides will play a pivotal role to meet these targets. Although the increasing deployment of renewable energies is encouraging, there still are many barriers to the replacement of power generation from traditionally high CO₂-emitting sectors based on coal and gas, which is still a critical and large portion of the energy generation, due to the intermittent nature of renewables. Hence, to simultaneously move away from fossil fuels and to circumvent the unpredictability inherent in clean energy resources, it is necessary to integrate energy-harvesting technologies with energy storage devices.

Energy storage, therefore, is emerging as a key enabler for sustainable renewable technologies, particularly for the electrification of transportation but also in more specialized applications such as heart defibrillators and active armor.³ Technologies already exist to store energy, such as batteries, electrochemical supercapacitors, and electrostatic capacitors.⁴⁻¹⁶ The latter are electrical energy-storage devices belonging to the category of passive components, which are ubiquitous in electronics. Indeed, every year more than 3 trillion multilayer ceramic capacitors (MLCCs) are manufactured from BaTiO₃ (BT), the prototypical ferroelectric (FE) ceramic.¹⁷⁻²²

In comparison with Li-ion batteries or fuel cells, the nonpolarized electrostatic or dielectric capacitors possess high power density (∼10⁴–10⁵ W/kg) resulting from their faster charging/discharging characteristics (∼μs), which are advantageous for power electronics in electrical vehicles (EVs) and pulse power applications (Figure 1a).^4,23−27 Hence, electrostatic capacitors are emerging as promising candidates for energy storage devices, where high power density in combination with high energy density are important technological requirements, as illustrated by the exponential rise in publications devoted to energy storage involving electrostatic ceramic capacitors, Figure 1b. Apart from high energy density and fast charging–discharging rate, other properties such as temperature/frequency stability, fatigue resistance, lifetime reliability, equivalent series resistance, and manufacturing cost are equally important for dielectric capacitors used in practical applications. New electroceramics are, therefore, required to facilitate near-engine power electronics, exhibit ultrafast charging, and have more durable EV performance at high temperature and voltage. Thus, future electroceramics must (i) deliver high energy density (W_rec > 10 J cm^–3) and conversion efficiencies (η > 90%); (ii) endure wider temperature ranges (−50–250 °C) and frequency ranges (1–1000 Hz); (iii) exhibit greater reliability (>10⁵ cycles) and fatigue resistance (<5% change over capacitor lifetime); and (iv) be compatible with cost-effective internal electrodes and be easily integrated with other components.

Figure 1.

(a) Applications for energy storage capacitors. *EMP: electromagnetic pulse. (b) Number of annual publications on lead-based ceramics, lead-free ceramics, ceramic multilayers, and ceramic films for energy storage capacitors from 2010 to 2020. (Collected from Web of Science, search “energy storage/density lead-based ceramic, lead-free ceramic, multilayer ceramic, ceramic capacitor, ceramic films but NOT polymer”). Reproduced with permission from PixaBay, Creative Commons License.

Historically, many different dielectric materials, ranging from paper and plastic to ceramics, have been employed in the fabrication of electrostatic capacitors. Nowadays, capacitors are fabricated from either polymers or ceramics because they offer the best combination of properties in terms of capacitance, dielectric loss, breakdown strength (BDS), and for the latter, thermal stability.

The prospects of employing ceramic capacitors for energy storage can be traced back to the 1960s work by Jaffe²⁸ from the Clevite Corp., USA. One decade later, Burn and Smyth²⁹ from Sprague Electric Company evaluated the energy storage performance in SrTiO₃ (ST) and BT with applied electric fields up to 400 kV cm^–1. Until that point, quantitative data of energy storage on these materials were limited to fields generally smaller than 150 kV cm^–1 due to the relatively low dielectric BDS of the fabricated ceramics. They emphasized that the maximum energy density for a ceramic should be obtained for thinner dielectric layers due to the lower probability for the occurrence of defects (such as pores, voids, or microcracks), which are well-known sources of dielectric breakdown. Later in 1990, Love,³⁰ also from Sprague Electric Company, revisited energy storage in ceramic capacitors and highlighted empirical design principles to achieve enhanced energy storage in capacitors, as shown in Table 1. Commercial C0G-type capacitors are manufactured from low relative permittivity (ε_r) linear dielectrics but may achieve an energy storage of 1 J cm^–3, by virtue of their intrinsically high BDS. The significance of the BDS, to achieve high energy storage becomes apparent in the case of X7R-type capacitors, fabricated from high ε_r BT. An important correlation between dielectric BDS and the thickness (t) can be extracted from Table 1. Indeed, by halving the t of the dielectric layers, the energy storage appears to increase >3 fold. This effect has been recently captured by Yang and co-workers,³¹ who compiled BDS data from literature for several dielectric materials of different t and observed decay inversely proportional to (t)^a, where a was determined as 0.5. Finally, when comparing the energy storage of Z5U and X7R, it becomes apparent that high ε_r alone is not a sufficient parameter to achieve high energy storage. Interestingly, Love³⁰ stressed that the capacitor industry was rather conservative in terms of perfecting the BDS of ceramics to reach values near those of single-crystals, which would significantly enhance the energy storage in ceramic capacitors.

Table 1. General Characteristics of Commercial Type Ceramic Materials Relevant for Energy Storage (Adapted from Love³⁰) Using Electronic Industries Alliance (EIA) Classifications^a.

dielectric type	dielectric BDS (V μm–1)	relative permittivity, εr	t (μm)	energy at 1 kV (J cm–3)
C0G (temperature coefficient 0 with tolerance ±30 × 10–6/K)	65	75	18.5	0.88
Z5U (+10/+85 °C, ΔC/C0 = +22/–56%)	13.2	7500	95	0.02
X7T (−55/+125 °C, ΔC/C0 = +22/–33%)	16	2800	70	0.71
X7R (−55/+125 °C, ΔC/C0 = ± 15%)	30	2000	38	1.40
X7R	40	2000	30	1.34
X7R	90	1800	20	4.82

^aClass I ceramic capacitors are accurate, temperature-compensating capacitors, C0G will have 0 drift with a tolerance of ±30 × 10^–6/K. Class II ceramic capacitors have a dielectric with a high permittivity. C and C₀ are represented capacitance value and capacitance value at 25 °C.

Love³⁰ proposed that maximum energy storage density can be achieved in intermediate rather than high ε_r materials since they exhibit larger BDS. Fletcher and co-workers³² convincingly postulated that greater energy storage densities can indeed be achieved in FE materials, whose Curie temperature (T_c) is adjusted to ensure that the material is operated in the paraelectric regime, where it shows a relatively small zero-field ε_r, an approach already mentioned by Jaffe in 1961.²⁸

In 2009, Ogihara and co-workers³³ proposed the use of so-called weakly coupled relaxors, such as 0.7BaTiO₃–0.3BiScO₃ (0.7BT–0.3BS), to fabricate energy storage devices. This new conceptual approach aimed at exploiting the extraordinary temperature stability of ε_r exhibit by this family of materials. When compared with commercial X7R capacitors, 0.7BT–0.3BS capacitors displayed superior performance, reaching a recoverable energy density (W_rec) of 6.1 J cm^–3 at 730 kV cm^–1. Again, the large dielectric BDS played a decisive role in this performance. More recently, in 2019 Wang, Reaney and co-workers³⁴ unveiled a novel approach to enhance energy storage characteristics via the fabrication of chemically heterogeneous but electrically homogeneous ceramics, with W_rec reaching 10.5 J cm^–3, as detailed later in this review.

Here, we present the principles of energy storage performance in ceramic capacitors, including an introduction to electrostatic capacitors, key parameters for evaluating energy storage properties, microstructural considerations, and critical electrical factors. Second, we will review the current state-of-the-art for lead and lead-free electroceramics for energy storage capacitors with bulk ceramics, ceramic multilayers (MLs), ceramic films and glass ceramics evaluated separately. Third, we will describe strategies for optimizing energy storage in electroceramics. Finally, we will demonstrate, with appropriate examples, a guide to the future development for electroceramics in energy storage capacitors.

2. Principles of Energy Storage in Electroceramics

2.1. Electrostatic Capacitors

The simplest dielectric capacitor consists of two parallel metallic plates separated by an insulator, which becomes polarized under the application of an electric field. This is the defining behavior of a dielectric material. The actual capacitance, C (i.e., ability to store charge), of an ideal capacitor is given by the ratio of the charge, Q, stored on each metallic plate and the applied voltage, V, as shown by eq 1.

1

Nevertheless, from a practical viewpoint, a more useful equation to compute the C of a real device, as illustrated in Figure 2, encompassing a dielectric material between two parallel plates of area, A, separated by a distance, d, subject to a V, can be obtained through the application of Gauss’s law

2

where ε is the permittivity of the dielectric, and a measure of its polarizability. Combination of eqs 1 and 2 provides the relationship:

3

From eq 3, it becomes immediately apparent that the ability of dielectric capacitor to charge and, therefore store energy, is ultimately associated with ε of the dielectric.

Figure 2.

Schematic representation of an electrostatic capacitor, where D, P, and ε₀ are electric displacement, polarization, and electric permittivity of free space (electric constant), respectively.

2.2. Key Parameters for Evaluating Energy Storage Properties

During the application of a V, the electrostatic energy stored, W, in the dielectric can be estimated by

4

where Q_max is the maximum charge achieved at the end of the charging cycle and dq is the incremental charge increase during the charging cycle. The volumetric energy density, W_st (i.e., the energy stored per volume unit, A d), is a common key performance indicator, expressed by

5

where E is the electric field and D_max is the electric displacement in the material under the maximum applied field, E_max. The electrical displacement (D) corresponds to the charge density (Q/A) on the metallic plates and is expressed by D = ε₀E + P (Figure 2), where P is the polarization (surface charge density).

For high ε materials, D is approximately equal to P, and it follows that D = εE = ε₀ε_rE, where ε₀ is the permittivity of free space (= 8.854 × 10^–12F m^–1) and ε_r is the relative permittivity, which is the ε/ε₀ ratio. This approximation allows stored energy density (W_st) to be defined in terms of P, as follows

6

where P_max is the maximum polarization reached at the E_max. From a practical viewpoint, eq 6 is prevalent in the calculation of W_st because several experimental methods exist to determine P under an applied E. In 1961, Jaffe²⁸ pointed out that the recoverable energy (W_rec) corresponds to the area above the discharging curve, whose upper limit is given by the P_max. Essentially, the mathematical integration of the area above a polarization-electric (P–E) loop provides an estimate of W_rec, as schematically illustrated in Figure 3 for four distinct types of polarization response.

Figure 3.

Four distinctive P–E hysteresis loops and their energy storage behavior: (a) linear, (b) FE, (c) relaxor-ferroelectric (with the schematic of energy storage calculation), and (d) antiferroelectric materials. *W_loss is loss energy density.

For linear dielectrics such as Al₂O₃, where ε_r is independent of the applied E. The calculation of W_rec from the P–E response illustrated in Figure. 3a, is given by

7

which clearly shows that W_rec is dependent on ε_r and E. Parts b–d of Figure 3 show cases where polarization responses deviate from linearity, and consequently, the computation of W_rec needs to be carried out using eq 6. The response illustrated in Figure 3b is typical of a classical FE material, such as BT, where the hysteresis is linked to polarization switching of macroscopic FE domains, as explained in detail in the review by Damjanovic.³⁵ Already in 1961, Jaffe²⁸ stressed that in FEs, charging energy is mainly absorbed by domain switching and is retained as remanent polarization (P_r). The typically high remanence of classical FEs can be effectively minimized via chemical doping, giving rise to the response shown in Figure 3c, which is characteristic of relaxor-ferroelectrics (RFEs), such as doped-BT and Pb(Mg_1/3Nb_2/3)O₃.³⁶

It is now generally accepted that relaxor behavior originates from the response of polar nanoregions (PNRs) to an alternating E. RFEs remain unsaturated at high applied E, and therefore, any increment of the E will have a contribution to energy storage. Remanence-free materials are therefore, preferable for achieving high W_rec. Linear dielectric materials meet this requirement but due to their low ε_r, energy storage is limited. Antiferroelectrics (AFEs) display low-remanence under low E but at large E the P–E loop opens due to the stabilization of an FE with respect to AFE phase and they display a saturated polarization, as illustrated in Figure 3d. In principle, therefore, as suggested by Jaffe,²⁸ AFEs should afford advantages for high energy storage, providing that dielectric breakdown issues are eliminated (i.e., the BDS should be high enough to induce the AFE-FE phase transition).

From the above, it becomes evident that nonlinear dielectric materials such as FEs, RFEs, and AFEs exhibit energy dissipation (W_loss); therefore, the W_rec is actually the most important parameter, as schematically illustrated in Figure 3c (red area). Hence, W_rec becomes

8

Energy conversion efficiency of a capacitor can then be calculated as

9

where W_loss is the energy loss during discharging, which correlates to the area enclosed by the P–E loop (Figure 3c green area).

Electric-field induced polarization can be determined via the measurement of charge, current, and voltage responses, typically achieved using either the Sawyer–tower, the virtual ground, the shunt or the current step methods. Each presents advantages and disadvantages as listed in Table 2. For details of each method, the reader is referred to Prume and co-workers.³⁷ Prume, Schmitz, and Tiedke proposed that overall the virtual ground method offers the highest precision for the measurement of FEs.

Table 2. Comparison of Different Hysteresis Measurement Methods for FEs³⁷.

method	measured quantity	reference component	integration necessary	bandwidth requirement	influence of parasitics
Sawyer–tower	charge Q	capacitor	no	moderate	high
virtual ground	current I	no	yes	high	low
shunt	current I	resistor	yes	high	high
current step	voltage V	no	no	moderate	moderate

2.3. Key Factors for Optimizing Energy Density

The microstructural features of electroceramics, such as density, grain size, secondary phases and core–shell structures, play an important role in energy storage properties. Simultaneously, the intrinsic electrical response, e.g., band gap, alongside the electrical microstructure, i.e., the distribution of conductive and resistive elements, are equally critical factors for the optimization of energy density. The following section reviews these factors, and gives examples of where and how they may be optimized.

2.3.1. Intrinsic Band Gap

The band gap (E_g) is the forbidden energy between the top of the valence band and bottom of the conduction band. E_g is commonly used to define insulator (E_g > 4.0 eV), semiconductor (0.0 eV < E_g < 4.0 eV), and metal (E_g = 0.0 eV). For semiconductor, the intrinsic BDS can be defined as

10

where BDS is direct proportional to E_g.³⁸ Thus, semiconductors with wider E_g have higher intrinsic BDS. The electronic structure and band gaps of semiconductor can be studied theoretically using, e.g., linear discriminant analysis, or experimentally, e.g., absorbance spectroscopy and diffuse reflectance spectroscopy.³⁹ A general rule of thumb is that the activation energy (E_a) for conduction is approximately half E_g. Both may be increased by doping or through the formation of solid solutions, often delivering higher BDS and W_rec.^40,41

For example, the highest E_g ∼ 3.58 eV among all different kinds of lead-free electroceramics was found in NaNbO₃ (NN), as shown in Figure 4a.⁴² Thus, NN was introduced into Na_0.5Bi_0.5TiO₃ (NBT) and BiFeO₃–BaTiO₃ (BF–BT) to enhance E_g. The E_g for BF–BT–xNN ceramics increased from 2.5 eV up to 2.95 eV for x ≤ 0.15, as shown in Figure 4b accompanied by significant enhanced W_rec ∼ 8.12 J cm^–3 under electric field ∼400 kV cm^–1, along with greater thermal stability (±10%, −50 to +250 °C) and ultrafast discharge rate (t_0.9 < 100 ns), Figure 4.⁴³

Figure 4.

(a) Comparison of the E_g among dielectric perovskites and a schematic of electronic breakdown. (b) Variation of average grain size and E_g as a function of NN concentration. (c) P–E loops and dP/dE under different E, (d) W_rec and η values, and (e) pulsed overdamped discharging energy density (W_D) of the BF–BT–0.10NN ceramic. (a) Reproduced with permission from ref (42). Copyright 2019 John Wiley and Sons. (b-e) Reproduced with permission from ref (43). 2020 John Wiley and Sons.

2.3.2. Electrical Microstructure

The distribution of regions with different conductivity and ε_r are important aspects of the so-called “electrical microstructure” of electroceramics.⁴⁴ In many instances, such as the core–shell microstructure or grain boundary response of BT based ceramics, the distributions are markedly heterogeneous and lead to localization of the electrical field strength in lower ε_r regions or pathways for breakdown through interconnected conducting regions. In 2019, electrical homogeneity was for the first-time proposed by Wang, Reaney and co-workers in the BF–BT system as a key factor to optimize BDS and as a consequence W_rec. Electrical heterogeneity was effectively eliminated by alloying with a third end-member so that it became more difficult to form a conductive pathway at high field, resulting in higher BDS and W_rec.³⁴

A homogeneous electrical microstructure may be obtained in many different ways such as heat-treatment in the appropriate atmosphere (N₂, Air, O₂) provided the type and magnitude of electrical conductivity is affected by oxygen stoichiometry. Practically, however, in production, a suitable dopant strategy is utilized once the conduction type is known (p vs n type). For example, the conductivity of BF-ST-based compositions is suppressed by doping with 3 mol % Nb on the B-site to compensate for Bi volatilization and the formation of oxygen vacancies (V_O^..), through variation of the Fe valence (Fe³⁺ to Fe⁴⁺).⁴⁵

For materials with more than one bulk-like region, e.g., phase mixtures, core–shell microstructures, or surface layers, alternating current (AC) impedance spectroscopy (IS) is able to show multiple responses and the resistance (R) and C can be extracted.⁴⁶⁻⁵³ Both the volume fraction and difference in magnitude of R and C for multiple electrical responses are equally important in influencing energy storage performance. Given the importance of the electrical microstructure, a brief outline of the role of IS is described and its advantages with respect to direct current methods are emphasized.

Direct current (DC) electrical measurements are the most commonly employed technique to evaluate the electrical characteristics of materials. However, they merely give the overall response instead of the properties of specific regions (e.g., grains and grain boundaries) unless microprobe techniques are employed.^54,55 Such techniques are useful but the sample volume is small, which casts doubt on their ability to represent global behavior and they are difficult to implement experimentally.

An alternative and much more convenient technique is IS. In IS measurements, an AC signal with small voltage over a wide range of frequency, typically 10^–2 to 10⁷ Hz, is applied on the sample.^44,56 The small voltage prevents any permanent change to the sample as well as yielding a (near) linear relationship between input and output. The wide range of frequencies allows separation of the response of different electro-active regions according to their relaxation times. For energy storage capacitors, impedance is capable of: (i) establishing the contributions to the electrical microstructure (grains, grain boundaries, core–shell structure and electrode–sample interface) and determine their individual conductivity and ε_r which give an insight into the distribution of electrical components within the sample; (ii) verifying the origin of the dominant electrical behavior (i.e., grains, grain boundary or interfacial layer response);^57,58 and (iii) determining the conduction mechanism and charge carrier type which helps further interpret the electrical response of the material.⁴⁷

Impedance can be defined as a complex number which usually contains both resistive and reactive (capacitive and/or inductive) components:

11

Different electro-active regions of a material are characterized by a R and a C, usually in parallel. Then the electric relaxation time or time constant, τ, of each region can be expressed as its R and C

12

at the frequency of maximum loss, ω_max, it holds the relation:

13

Due to their different R and C values, electro-active regions can be separated in the frequency domain. Once the value of R and C are extracted, they can then be assigned to appropriate regions of the sample.

Normally the impedance measurement needs to be taken over a temperature and/or oxygen partial pressure (pO₂) range to gain a better understanding of the conduction mechanism and the charge carrier. The associated activation energy, E_a, can be estimated using Arrhenius equation

14

where σ is the conductivity, σ₀ is pre-exponential factor, k is the Boltzmann constant, and T is temperature. E_a may be related to predominant charge carrier and conduction mechanism. The type of charge carrier may also be determined to some extent by the pO₂ dependence of conductivity, i.e., p-type: conductivity increases with increasing pO₂; n-type: conductivity decreases with increasing pO₂; ionic charge carrier: conductivity is independent with pO₂.

2.3.3. Density and Porosity

The density of the ceramic materials plays an essential role on electrical performance, especially BDS. Ceramics with higher density tend to support higher E closer to the intrinsic/theoretical BDS. In contrast, low density ceramics exhibit conductive pathway composed of pores/voids which result in short circuit under modest field strengths. The relationship between the voltage across the pore and the external E based on a “slab” model is shown below

15

where V_c and V_ext are the voltage applied cross the cavity pore and external applied voltage, ε_c and ε_d are the permittivity of the cavity and the dielectric, respectively,^59,60 and td and tc are the thicknesses of the dielectric and cavity, respectively. Thus, the local E increases markedly for materials with larger pores and pore volumes, resulting in lower BDS.

High density electroceramic materials are commonly obtained by optimization of the sintering conditions, including sintering temperature/time and heating/cooling rate. For ceramics that are difficult to densify using a conventional approach, sintering aids are often added.⁶¹⁻⁶⁴ Higher density ceramics may be obtained by the addition of ZnO,⁶⁵ CuO,⁶⁶ and MgO,⁶² which enhances BDS and W_rec. For K_0.5Na_0.5NbO₃–Bi(Mg_2/3Nb_1/3)O₃ (KNN–BMN), small amounts of CuO help densify ceramics through the formation of a transient liquid phase, as reported by Qu and co-workers (Figure 5).⁶⁴ The sintering temperature was also reduced from 1150 to 930 °C, allowing compatibly with Cu or Ag/Pd internal electrode in MLs and giving rise to W_rec ∼ 4.02 J cm^–3 at 400 kV cm^–1 for 0.9KNN-0.1BMN with 1% mol CuO.⁶⁶

Figure 5.

Scanning electron microscopy (SEM) images of the 0.9KNN-0.1BMN-x mol % CuO ceramics with (a) x = 0.25; (b) x = 0.5; (c) x = 1.0; (d) x = 1.5, as liquid phase is circled in red. (e) Unipolar P–E hysteresis loops and (f) Calculated W and W_rec under different E of 0.9KNN-0.1BMN-1 mol % CuO ceramics.⁶⁶ Reproduced with permission from ref (66). Copyright 2017 John Wiley and Sons.

Different sintering technologies, such as spark plasma sintering (SPS), two-step sintering,⁶⁷ and the formation of coatings using chemical methods,⁶⁸⁻⁸¹ have also been shown to improve density and give rise to higher BDS and W_rec.

2.3.4. Grain Size

The effect of grain size (G) on energy storage properties has been discussed for several electroceramics because of the relationship between BDS and G, expressed in eq 16

16

where a is the exponent values being in the range of 0.2–0.4.^31,82−84 Waser explained that leakage current in fine-G ceramics is lower than coarse-G ceramics due to the high grain boundary density which act as barriers for charge carriers.⁸⁵ Thus, dielectric materials with high density and fine-G are required to optimize energy storage. G may be tailored by chemical doping and the formation of solid solution. It may also be modified by the application of an ultrathin coating on the primary particles prior to sintering via chemical coating methods, e.g., SiO₂ on BT ceramics.^{67,77,86−89} The optimization on E_max and W_rec via grain size-engineering for several materials is illustrated in Figure 6.

Figure 6.

Relationship between energy storage properties of ceramics and G: (a) G vs E_max and (b) G vs W_rec. *AN: AgNbO₃.

For example, an average G ∼ 10 μm was reported for BF–BT ceramics, which was reduced to <2 μm after A-site Nd doping, as shown in Figure 7. Meanwhile, improved W_rec ∼ 1.8 J cm^–3 and η ∼ 88% were obtained for 15 mol % Nd–BF–BT and 40 mol % Nd–BF–BT, respectively.⁹⁰ Similar optimization behavior has also been found in KNN–BMN and KNN–ST ceramics, resulting in BDS ∼ 400 kV cm^–1 and W_rec > 3.5 J cm^–3.^91,92

Figure 7.

SEM images of x mol % Nd-doped BF–BT with different Nd concentrations: (a) BF–BT, (b) 2.5 mol % Nd–BF–BT, (c) 5 mol % Nd–BF–BT, (d) 7.5 mol % Nd–BF–BT, (e) 10 mol % Nd–BF–BT, (f) 15 mol % Nd–BF–BT, (g) 20 mol % Nd–BF–BT, (h) 30 mol % Nd–BF–BT, and (i) 40 mol % Nd–BF–BT; the G distributions of Nd-doped BF–BT are shown in the insets of the SEM images.⁹⁰ (j) G, density and (k) energy storage performance at 170 kV cm^–1, as a function of x(Nd) mol % in BF–BT ceramics. Reproduced with permission from ref (90). Copyright 2017 Royal Society of Chemistry.

2.3.5. Core–Shell Structure

Core–shell subgrain microstructures are observed in many lead-free ceramics, due to either kinetic limitations of the diffusion process (typical for BT based ceramics) or immiscibility on cooling from high temperature for perovskite end members with dissimilar ion size and bonding (BF based ceramics).^22,34,45 The effect of core–shell microstructures on energy storage performance is still unclear. In BT-based ceramics, the cores are often more conducting than the doped shells and core to core conductive pathways lead to breakdown.⁹³⁻⁹⁶ For BF based ceramics, the defect chemistry of the cores and shells remains to be elucidated, but initial work suggests that further dopants are needed to create electrical homogeneity and thus eliminate the conducting pathways.^34,45 The theoretical modeling has reported a positive influence of core–shell microstructure but none have been unambiguously validated experimentally.⁹⁷

3. State-of-the-Art in Electroceramics for Energy Storage

3.1. Bulk Ceramics

3.1.1. Lead-Based Ceramics

Lead-based ceramics are used commercially as energy storage materials for high-power pulsed capacitors due to their excellent W_rec and η.⁹⁸⁻¹⁰¹ The energy storage properties of RFE and AFE lead-based ceramics are summarized in Table 3.

Table 3. Summary of Energy Storage Properties for Lead-Based Ceramics^a.

compositions	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
(Pb0.89Ba0.08La0.02)(Zr0.7Sn0.27Ti0.03)O3	135	22.6	2.1	76.5	(150)
(Pb1.067La0.02)(Zr0.95Ti0.05)O3	90	39.5	2.12	92.98	(141)
0.90(Pb0.97La0.02)(Zr0.65Sn0.30Ti0.05)O3–0.10Bi(Zn2/3Nb1/3)O3	115	29	2.19	95.6	(162)
Pb0.97La0.02(Zr0.58Sn0.35Ti0.07)O3	118	29.0	2.35	86.1	(158)
Pb0.91La0.02Ba0.06(Zr0.65Sn0.3Ti0.05)O3	150	29.5	2.4	82	(159)
(Pb0.93Ba0.04La0.02)(Zr0.65Sn0.3Ti0.05)O3–0.005Mn2O3	308	31.5	2.64	73	(161)
Pb0.97La0.02(Zr0.33Sn0.55Ti0.12)O3@0.05SiO2	238	34.6	2.68	83.5	(87)
(Pb0.87Ba0.1La0.02)(Zr0.65Sn0.3Ti0.05)O3–0.75Y	130	46.5	2.75	71.5	(149)
(Pb0.88La0.08)(Zr0.91Ti0.09)O3	170	31.5	3.04	92	(112)
1.7 mol % Pr3+ doped 0.24Pb (In1/2Nb1/2)O3–0.42Pb(Mg1/3Nb2/3)O3-0.34PbTiO3	50	20	3.1	90	(109)
0.92Pb(Tm0.5Nb0.5)O3–0.08Pb(Mg1/3Nb2/3)O3	310	17.03	3.12		(136)
(Pb0.87Ba0.1La0.02)(Zr0.68Sn0.24Ti0.08)O3	180	58.2	3.2		(151)
Pb0.97La0.02(Zr0.50Sn0.46Ti0.04)O3	150	43	3.2	86.5	(124)
0.55(Pb0.97La0.02)(Zr0.93Sn0.05Ti0.02)O3–0.45(Pb0.93Ba0.04La0.02) (Zr0.65Sn0.3Ti0.05)O3	180	25	3.2	74.4	(160)
Pb0.97La0.02(Zr0.56Sn0.35Ti0.09)O3	175	39.4	3.3	80	(166)
(Pb0.895La0.07)(Zr0.9Ti0.1)O3	175	42.3	3.38	86.5	(115)
0.9PbHfO3–0.1Pb(Mg0.5W0.5)O3	155	43.5	3.7	72.5	(134)
Pb0.94La0.04(Lu0.5Nb0.5)O3	681		3.85		(137)
(Pb0.955La0.03)(Zr0.50Sn0.42Ti0.08)O3	180	41	3.99	79.2	(127)
Pb0.97La0.02(Zr0.60Sn0.35Ti0.05)O3	200	34.48	4.1		(121)
(Pb0.97La0.02)(Zr0.5Sn0.44Ti0.06)O3	250	29.3	4.2	82	(117)
(Pb0.97La0.02)(Zr0.5Sn0.44Ti0.06)O3	250	29.3	4.2	82	(118)
Pb0.955La0.03(Zr0.5Sn0.43Ti0.07)O3	200	36	4.2	78	(126)
(Pb0.97La0.02)(Zr0.8Sn0.145Ti0.055)O3	225	34	4.38	73	(124)
(Pb0.858Ba0.1La0.02Y0.008)(Zr0.65Sn0.3Ti0.05)O3- (Pb0.97La0.02)(Zr0.9Sn0.05Ti0.05)O3	200	46.8	4.65	60	(152)
La0.02Pb0.97[(Yb0.5Nb0.5)0.92Ti0.08]O3	240	34	5.18	65	(135)
(Pb0.97La0.02Zr0.85Sn0.12Ti0.03O3)−0.5 wt % Al2O3	315	35.5	5.3	88.3	(72)
(Pb0.955Sr0.015La0.02)(Zr0.75Sn0.195Ti0.055)O3	350	33.5	5.56	70	(156)
Pb0.97La0.02(Zr0.5Sn0.45Ti0.05)O3	400	36.2	5.6	63	(116)
(Pb0.858Ba0.1La0.02Y0.008)(Zr0.65Sn0.3Ti0.05)O3–(Pb0.97La0.02)(Zr0.9Sn0.05Ti0.05)O3	306	48.5	6.4	62.4	(173)
Pb[(Lu0.5Nb0.5)–(Mg0.5W0.5)]O3	340	46	6.4	71	(132)
Pb0.91La0.06(Zr0.552Sn0.368Ti0.08)O3@1 wt %PbO–B2O3–SiO2–Al2O3–ZnO–MnO2	380	43	7.4	91.6	(77)
PbHfO3	270	44.5	7.6	80.8	(133)
Pb0.98La0.02(Hf0.45Sn0.55)0.995O3	380	36	7.63	94	(138)
(Pb0.91La0.06)(Zr0.96Ti0.04)O3–1.0 mol % MnCO3	300	43.5	7.65	87	(145)
(Pb0.98La0.02)(Zr0.55Sn0.45)0.995O3	400	41.5	10.4	87	(131)
(Pb0.94La0.02Sr0.04)(Zr0.9Sn0.1)0.995O3	400	44	11.18	82.2	(130)

^at of the bulk ceramics is commonly >0.1 mm.

3.1.1.1. Lead-Based Relaxor-Ferroelectrics

Many lead-based RFEs, including Pb(Mg_1/3Nb_2/3)O₃–PbTiO₃ (PMN–PT), Pb(Zn_1/3Nb_2/3)O₃–PbTiO₃ (PZN–PT), and (Sr,Pb,Bi)TiO₃ (SPBT)-based materials, have been reported as potential candidates for energy storage capacitors.^102−110 Zhang and co-workers investigated the relaxation behavior and energy storage properties of (1–x)PMN–xPT ceramic, obtaining W_rec ∼ 0.47 J cm^–3 at room temperature.¹⁰² Li and co-worker probed the effect of domain structure on W_rec and thermal stability of 0.2PMN–0.8Pb(Sn_xTi_1–x)O₃ (PMN–PS_xT_1–x) ceramics, as illustrated in Figure 8. 0.2PMN–0.8PST ceramics exhibited W_rec ∼ 0.85 J cm^–3 with excellent thermal stability which was attributed to the coexistence of ferroelectric domains and PNRs.¹¹¹

Figure 8.

W_rec of 0.2PMN–0.8PS_xT_1–x ceramics with different Sn (x) contents at 70 kV cm^–1. The insets show the mechanism of enhanced energy storage due to coexistent-phase structure and the W_rec for PMN–PS_xT_1–x ceramics under different electric fields. Reproduced with permission from ref (111). Copyright 2018 Elsevier.

3.1.1.2. Lead-Based Antiferroelectrics

PbZrO₃ (PZ) is the first known AFE and exhibits a double P–E hysteresis loop below T_C. However, the high critical switching field required for an AFE–FE phase transition at room temperature limits applications for energy storage. Chemical substitution to reduce switching field is an effective strategy to overcome the problem and three well-known PbZrO₃ based compositions are reviewed: (i) (Pb,La)(Zr,Ti)O₃ (PLZT);^112−115 (ii) (Pb,La)(Zr,Sn,Ti)O₃ (PLZST);^116−129 and (iii) (Pb,La)(Zr,Sn)O₃ (PLZS).^130,131 Additionally, some new AFEs have also been identified based on PbHfO₃, Pb(Lu_0.5Nb_0.5)O₃, Pb(Yb_0.5Nb_0.5)O₃, and Pb(Tm_0.5Nb_0.5)O₃.^132−138

3.1.1.2.1. (Pb,La)(Zr,Ti)O₃ (PLZT)

According to the phase diagram (Figure. 9), PLZT exists as homogeneous compositions over a wide range of mol % La in the PbZrO₃–PbTiO₃ solid solution.^139,140

Figure 9.

Phase diagram of PLZT at room temperature.¹³⁹ Reproduced with permission from ref (139). Copyright 2014 Elsevier.

When Pb ions are replaced by ≤30 mol % La on the A-site in accordance with a lead vacancy (V_pb^..) ionic compensation model, an orthorhombic AFE phase similar to PbZrO₃ occurs for Zr rich compositions. However, only PLZT with <10 mol % La is typically utilized for energy storage applications^141−147 since higher concentrations have lower polarization and therefore lower W_rec. Li and co-workers prepared (Pb_0.97La_0.02)(Zr_0.95Ti_0.05)O₃ ceramics via a solid-state reaction route, yielding W_rec ∼ 0.83 J cm^–3 and η ∼ 70% under an electric field of 55 kV cm^–1.¹¹⁴ Jo and co-workers found that AFE and RFE behavior can both be obtained by substitution of La and excess PbO in PLZT, resulting in the enhancement of W_rec by promoting a slim and slanted hysteresis loop. Both high W_rec ∼ 3.04 J cm^–3 and η ∼ 92% were obtained along with no performance degradation after 10⁵ cycles.¹¹² Tuning the Zr/Ti ratio has also shown to be an effective way to improve W_rec of PLZT ceramics. Qiao and co-workers reported slimmer P–E loops giving a W_rec ∼ 3.38 J cm^–3 in (Pb_0.895La_0.07)(Zr_xTi_1–x)O₃ ceramics by changing the Zr/Ti ratio (Figure 10), which was attributed to the reduction of tolerance factor and “enhancement of antiferroelectricity”.¹¹⁵ Mn doping is also suggested to improve W_rec of PLZT by “enhancing antiferroelectricity”.^143−145

Figure 10.

Effect of Zr/Ti ratio on P–E loops and energy storage properties of PLZT.¹¹⁵ Reproduced with permission from ref (115). Copyright 2019 Elsevier.

3.1.1.2.2. (Pb,La)(Zr,Sn,Ti)O₃ (PLZST)

To further optimize the energy storage properties of PLZT ceramics, Sn may be substituted on the B-site of PLZT, which broadens the AFE compositional range,^{72,77,87,116,122,148−166} in accordance with phase diagram from 1989 (Figure 11).¹⁶⁷ (Pb_0.97La_0.02)(Zr,Sn,Ti)O₃ with 2 mol % La has been most commonly studied in which the Zr/Ti/Sn ratio is varied to give a complex phase diagram that contains FE tetragonal (F_T), high-temperature nontilted FE rhombohedral (F_R(HT)), low temperature FE rhombohedral (F_R(LT)) AFE tetragonal (A_T), and AFE orthorhombic (A_O) phases.

Figure 11.

Phase diagram of (Pb_0.97La_0.02)(Zr,Sn,Ti)O₃, where T, R, and O represent the tetragonal, rhombohedral, and orthorhombic structure, respectively, and HT and LT represent high and low temperature, respectively. Reproduced with permission from ref (167). Copyright 2005 John Wiley and Sons.

We note that the A_T phase in both PLZT and PLZST has been shown to be incommensurate by a number of researchers and might be better described as a A_O phase in which there is disorder of antipolar coupling.^168−170 PLZST is AFE for concentrations with <15 mol % Ti. For compositions with A_O structure, the critical phase switching fields are above BDS but ceramics with the A_T phase undergo electric field-induced AFE-FE switching at room temperature, for which the switching field increases with increasing Sn concentration (Figure 12).¹¹⁸ Adjusting the Zr/Sn/Ti ratio leads to optimization of W_rec in PLZST ceramics^{116,118,148,157} with an increase in Sn concentration leading to a reduction in switching field (forward threshold electric field, E_F, and backward threshold electric field, E_B) and weakening ferroelectricity.¹⁷¹W_rec ∼ 5.6 J cm^–3 and high thermal stability have been obtained in PLZST ceramics with a Zr/Sn/Ti ratio of 0.5:0.45:0.05¹¹⁶ while Wang and co-workers reported, in their study of the (Pb_0.97La_0.02)(Zr_0.5Sn_0.5-xTi_x)O₃ solid solution, superior W_rec of 4.2 J cm^–3 with η of 82% for (Pb_0.97La_0.02)(Zr_0.5Sn_0.44Ti_0.06)O₃ ceramics with good temperature stability.¹¹⁸ Recently, a ferrielectric (FIE) configuration was reported in PLZST which consists of ferroelectric ordering segments with either magnitude or angular modulation of dipoles.¹⁷² The net polarization of field-induced FE order can be tailored by adjusting the Sn/Ti ratio.

Figure 12.

AFE-type P–E hysteresis loops of Pb_0.97La_0.02(Zr_0.5Sn_0.5–xTi_x)O₃ with x = (a) 0.10, (b) 0.08, (c) 0.06, and (d) 0.04.¹¹⁸ Reproduced with permission from ref (118). Copyright 2016 Elsevier.

The performance of PLZST ceramics is also influenced by the Pb/La ratio.^117,122,126 The AFE phase becomes more stable with a commensurate increase in the AFE–FE switching field as La concentration increases. The energy storage properties of (Pb_1–1.5xLa_x)(Zr_0.5Sn_0.43Ti_0.07)O₃ ceramics were optimized (W_rec of 4.2 J cm^–1) by Dan and co-workers for compositions with x = 0.03 due to a large switching electric field and high BDS.^3,126 Furthermore, doping with Ba and Sr (A-site) improves fatigue behavior and temperature stability, suppresses the stress sensitivity, and enhances energy storage.^{148,150−153,155,161,173}

3.1.1.2.3. (Pb,La)(Zr,Sn)O₃ (PLZS)

Wang and co-workers reported a unique E-induced multiphase transition in PLZS for which a conventional AFE–FE phase transition at low E, followed by a second FE-FE phase transition at a higher E, leads to an increase in polarization.¹³¹W_rec of 10.4 J cm^–3 and η of 87% were achieved at 400 kV cm^–1 for (Pb_0.98La_0.02)(Zr_0.55Sn_0.45)_0.995O₃ ceramics, along with superior discharge current density of 1640 A cm^–2 and ultrafast discharge speed (75 ns discharge period) (Figure 13a,b).¹³¹ Subsequently, a record-high W_rec∼ 11.2 J cm^–3 and a high η ∼ 82% were realized in (Pb_0.98–xLa_0.02Sr_x)(Zr_0.9Sn_0.1)_0.995O₃ ceramics, as illustrated in (Figure 13c,d). The substitution of Pb by Sr gave rise to an increase in BDS and AFE/FE switching fields, leading to further enhancement of energy storage performance.¹³⁰

Figure 13.

(a) Bipolar P–E hysteresis loops and (b) energy storage properties of (Pb_0.98La_0.02)(Zr_0.55Sn_0.45)_0.995O₃ ceramics under different applied fields.¹³¹ (c) Unipolar P–E hysteresis loops of the (Pb_0.94La_0.02Sr_0.04)(Zr_0.9Sn_0.1)_0.995O₃ ceramic under different applied fields. (d) W_rec of (Pb_0.98-xLa_0.02Sr_x)(Zr_0.9Sn_0.1)_0.995O₃ with Sr concentration (x = 0–0.06) as a function of the E.¹³⁰ (a, b) Reproduced with permission from ref (131). Copyright 2019 John Wiley and Sons; (c, d) Reproduced with permission from ref (130). Copyright 2019 Royal Society of Chemistry.

3.1.2. Lead-Free Ceramics

In the last decades, extensive research has focused on lead-free electroceramics due to concerns over the toxicity of lead/lead oxide-based compounds.^174−176 As a result, there has been a steady but continuous improvement in their energy storage performance,^{31,103,177−180} with a view to replacing existing lead-based materials. Several lead-free ceramic systems are considered as potential candidates to replace PLZT for energy storage applications, including those based on BT, ST, KNN, BF, NBT, AgNbO₃ (AN), and NN.

3.1.2.1. BaTiO₃-Based Ceramics

BT-based dielectric ceramics have been studied for decades and dominate the commercial market of ceramic capacitors.^81,181 Several studies have reported improvements in energy storage performance of BT-based ceramics through (i) substituting oxides to improve BDS, such as Al₂O₃, La₂O₃, MgO, SiO₂;^{79,80,182−184} (ii) employing different sintering techniques, such as SPS, citrate precursor, and cold sintering (CS) to increase ceramic density or control grain growth;^68,76,185 (iii) adding sintering aids such as ZnNb₂O₆ and NiNb₂O₆ to increase density;^186,187 (iv) introducing further end-members in the solid solution, Bi(Mg_1/2Ti_1/2)O₃,^188,189 BiYbO₃,¹⁹⁰ BMN,^191−193 NBT–Na_0.73Bi_0.09NbO₃,¹⁹⁴ Nd(Zn_1/2Ti_1/2)O₃,¹⁹⁵ Bi(Zn_2/3Nb_1/3)O₃,¹⁹⁶ Bi(Li_1/2Nb_1/2)O₃,^197,198 Bi(Zn_1/2Zr_1/2)O₃,¹⁹⁹ Bi(Zn_1/2Ti_1/2)O₃,²⁰⁰ YNbO₄,²⁰¹ Bi_0.9Na_0.1In_0.8Zr_0.2O₃,²⁰² Bi(Li_1/2Ta_1/2)O₃,²⁰³ Bi(Mg_1/2Zr_1/2)O₃,²⁰⁴ Bi(Zn_1/2Sn_1/2)O₃,²⁰⁵ K_0.5Bi_0.5TiO₃–KNbO₃,²⁰⁶ and K_0.5Bi_0.5TiO₃–NN²⁰⁷ to promote RFE behavior. The energy storage properties of BT-based materials are summarized in Table 4.

Table 4. Energy Storage Properties of BT-Based Materials^a.

compd	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
0.9BT–0.1BMN	143.5	∼16	1.13	95	(192)
Ba0.85Ca0.15Zr0.1Ti0.9O3	200	∼15	1.153		(182)
0.9BT–0.1(Bi0.9Na0.1)(In0.8Zr0.2)O3	180	∼20	1.33	88	(202)
0.88(Ba0.8Sr0.2)TiO3–0.12 Bi(Zn2/3Nb1/3)O3	225	17	1.62	99.8	(210)
0.9BT–0.1BMN–0.3 wt % MnCO3	205	16.28	1.7	88.5	(191)
0.92(0.65BT–0.35NBT)–0.08(Na0.73Bi0.09NbO3)	172	∼25	1.7	82	(194)
0.92(0.92BT–0.08K0.5Bi0.5TiO3)–0.08NN	220	∼23	1.96	67.4	(207)
0.88BT–0.12Bi(Li0.5Nb0.5)O3	270	∼19	2.03	88	(197)
0.9BT–0.1Bi(Li0.5Ta0.5)O3	280	∼11.9	2.2	88	(203)
0.85BT–0.15Bi(Zn0.5Sn0.5)O3	280	∼23	2.41	91.6	(205)
0.9BT–0.1Bi(Zn0.5Zr0.5)O3	264	∼25	2.46		(199)
0.85BT–0.15Bi(Mg0.5Zr0.5)O3	280	∼23	2.9	86.8	(204)
0.9Ba0.65Sr0.35TiO3–0.1BMN	400	23	3.34	85.71	(211)
0.65(Ba0.98Li0.04)Ti0.98O3–0.35(Sr0.7Bi0.2)TiO3	410	35	3.54	77	(212)
0.6BT–0.4Bi(Mg0.5Ti0.5)O3	340	∼40	4.49	93	(208)
BT–0.06Bi2/3(Mg1/3Nb2/3)O3	520	25	4.55	91	(209)

^at of the bulk ceramics is commonly >0.1 mm.

The most effective approach, however, is the introduction of a Bi-based perovskite end member in which the B-site contains multiple cations. Doping with a Bi ion that has a lone pair electronic 6s² configuration into the Ba-site increases P_max. P_r is minimized by forming a so-called “weakly coupled relaxor” state in which long-range polar coupling is disrupted by the multiple B-site ions which have different valence and ionic radius to Ti. Using this principle, Hu and co-workers²⁰⁸ reported high W_rec of 4.49 J cm^–3 with a η of 93% for 0.6BT–0.4 Bi(Mg_0.5Ti_0.5)O₃ (BT–BMT) ceramics that were temperature stable to 170 °C (variation W_rec < 5% and η < 4%).

Of greater commercial potential, Yang and co-workers reported similar properties with BT–0.06 Bi_2/3(Mg_1/3Nb_2/3)O₃ (BT–0.06B_2/3MN) but in compositions compatible with Ag/Pd electrodes due to the presence of only 4 mol % Bi on the A-site.²⁰⁹ Similar energy storage properties, W_rec ∼ 4.6 J cm^–3 and η ∼ 92% (Figures 14a,b) to BT–BMT were obtained for BT–0.06B_2/3MN ceramics which also benefited from the highest BDS, ∼520 kV cm^–1, among all reported BT-based compositions.²⁰⁹ BT–0.06B_2/3MN is not electrically homogeneous but BDS and W_rec were still optimized by reducing, though not completely eliminating, the difference between the bulk and grain boundary responses with respect to undoped BT, Figures 14c,d.²⁰⁹

Figure 14.

(a) Unipolar P–E loops under E_max and (b) calculated W_rec and η at different electric field for BT–0.06B_2/3MN ceramics. (c) Z* plots of (c) BT at 400 °C and (d) BT–0.06B_2/3MN at 550 °C.²⁰⁹ Reproduced from ref (209). Copyright 2020 American Chemical Society.

3.1.2.2. SrTiO₃-Based Ceramics

ST, which is an incipient ferroelectric, is considered as a promising candidate for energy storage applications due to its relatively high permittivity (ε_r ∼ 300) and low dielectric loss (<1%) at room temperature. Extensive efforts have been made to improve the energy storage performance of ST-based ceramics, including (i) doping with Ba,^213−216 Dy,²¹⁷ Mg,²¹⁸ Ce,⁶⁴ and Bi^219−222 on the A-site or Mn²²³ and Sn^224,225 on the B-site; (ii) using sintering aids, such as ZnO,⁶⁵ MgO,^226−228 SiO₂,^63,229 SrO–B₂O₃–SiO₂,²³⁰ ZnO–Li₂O,²³¹ Bi₂O₃–B₂O₃–SiO₂,²³² B₂O₃–SiO₂–Bi₂O₃–CaO–BaO,²³³ Al₂O₃,²³⁴ BaO–TiO₂–SiO₂,⁷⁵ BaCu(B₂O₅),²³⁵ and NiNb₂O₆;²³⁶ (iii) employing different sintering techniques such as microwave sintering^73,75 and SPS;⁷⁰ and (iv) introducing complex end-members, such as NBT,^237,238 NBT–Ba(Al_0.5Nb_0.5)O₃,²³⁹ NBT–BT,^118,240,241 and Bi(Mg_0.5Hf_0.5)O₃.²⁴² The energy storage properties of ST-based materials are summarized in Table 5.

Table 5. Energy Storage Properties of ST-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec(J cm–3)	η (%)	ref
95 wt % Ba0.4Sr0.6TiO3–5 wt %MgO	300	11.88	1.5	88.5	(227)
Ba0.4Sr0.6TiO3–8 mol % SiO2	400	9.11	1.6	90.9	(229)
0.6ST–0.4NBT	210	∼20	1.7	69.93	(237)
0.45ST–0.2NBT–0.35BT	170	∼25	1.78	77.06	(243)
(Sr0.99Mg0.01)TiO3	362	∼12	1.86	89.3	(218)
0.5ST–0.5(0.94 Bi0.54Na0.46TiO3–0.06BT)	180	∼30	1.88	79	(239)
0.5ST–0.5(0.95NBT–0.05BaAl0.5Nb0.5O3)	190	∼28	1.89	77	(240)
Ba0.4Sr0.6TiO3–9 wt %(Bi2O3–B2O3–SiO2)	279	∼17	1.98	90.57	(232)
0.95(Sr0.5Na0.25Bi0.25TiO3)–5 wt %MgO–0.05KNbO3	178.5	∼27	2	76.34	(244)
Ba0.4Sr0.6(Ti0.996Mn0.004)O3–2 wt % MgO	300	∼16	2.014	88.6	(228)
Ba0.3Sr0.6Ca0.1TiO3–2 wt %MgO	194.33	∼23	2.223		(245)
Sr0.985Ce0.01TiO3–3 wt % SiO2	290	∼10	2.23		(64)
SrTi0.985(Zn1/3Nb2/3)0.015O3–4.5 wt %ZnNb2O6	422	9.34	2.35	77	(246)
0.9ST–0.1(Bi0.48La0.02Na0.48Li0.02Ti0.98Zr0.02O3)	323	∼20	2.59	85	(247)
0.8ST–0.2(NBT– Ba0.94La0.04Zr0.02Ti0.98O3)	320	∼22	2.83	85	(248)
0.995(0.6ST−0.4NBT)−0.005ZrO2	285	∼25	2.84	71.54	(74)
0.9(Sr0.7Bi0.2)TiO3–0.1 Bi(Mg0.5Hf0.5)O3	360	∼22	3.1	93	(242)
Ba0.3Sr0.7TiO3–1.6 wt % ZnO	400		3.9		(65)
98.5 wt %Ba0.4Sr0.6TiO3–1.254 wt %Al2O3–0.246 wt %SiO2	493		5.09		(234)

^aThe t of the bulk ceramics is commonly >0.1 mm.

From a review of the literature, doping commonly increases both ε_r and P_max but decreases the BDS, sintering aids increase BDS but lower the P_max. The highest energy densities have therefore been achieved through either dopants and/or alloying with “relaxor end-members” which also act as sintering aids. Adopting these protocols, a W_rec of 3.1 J cm^–3 with η ∼ 93% was obtained for 0.9(Sr_0.7Bi_0.2)TiO₃–0.1 Bi(Mg_0.5Hf_0.5)O₃ ceramic at 360 kV cm^–1, Figures 15a,b,²⁴² which was also fatigue-resistant up to 10⁵ cycles and temperature stable from 25 to 200 °C, Figures 15c–f.²⁴²

Figure 15.

(a) Unipolar P–E loops and (b) W, W_rec, and η of 0.9(Sr_0.7Bi_0.2)TiO₃–0.1Bi(Mg_0.5Hf_0.5)O₃ ceramic as functions of E. (c) Unipolar P–E loops and (d) W_rec and η of 0.9(Sr_0.7Bi_0.2)TiO₃–0.1Bi(Mg_0.5Hf_0.5)O₃ ceramic as functions of cycle numbers up to 10⁵. (e) Unipolar P–E loops, with the inset shows the P_max and P_r as functions of temperature, and (f) W_rec and η of 0.9(Sr_0.7Bi_0.2)TiO₃–0.1Bi(Mg_0.5Hf_0.5)O₃ ceramics as a function of temperature. Reproduced with permission from ref (242). Copyright 2019 John Wiley and Sons.

3.1.2.3. K_0.5Na_0.5NbO₃-Based Ceramics

In 2016, the energy storage properties of KNN–(Bi,Na)HfO₃ solid solutions were first investigated, and W_rec ∼ 0.54 J cm^–3 was achieved at 129 kV cm^–1.²⁴⁹ ZnO and CuO were introduced as sintering aids improved W_rec in KNN-based ceramics^66,250 by decreasing porosity and restricting grain growth. BDS of 400 kV cm^–1 was obtained in 0.8KNN-0.2Sr(Sc_0.5Nb_0.5)O₃ ceramic with 0.5 mol % ZnO²⁵⁰ and CuO reduced the sintering temperature and enhanced the density of 0.9KNN–0.1BMN ceramics.⁶⁶ A third perovskite end-member, e.g., ST, BF, and BMN, has also been shown to optimize energy storage properties.^91,92,251W_rec of 4.03 J cm^–3 at 400 kV cm^–1 was obtained in 0.85KNN–0.15ST bulk ceramics⁹² with similar energy storage performance realized in 0.90KNN–0.10BMN ceramic with a large P_max (41 μC cm^–2) obtained at 300 kV cm^–1, Figure 16.⁹¹ The energy storage properties of KNN-based materials are summarized in Table 6.

Figure 16.

(a) Unipolar P–E hysteresis loops and (b) calculated W and W_rec and (c) W_loss and η for 0.90KNN–0.10BMN ceramics under different E. Reproduced with permission from ref (91). Copyright 2017 Royal Society of Chemistry.

Table 6. Energy Storage Properties of KNN-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
(K0.48Na0.52)0.88Bi0.04NbO3	189	∼12.7	1.04		(252)
0.93KNN–0.07BMN	150	∼22.5	1.3	58.8	(253)
0.9KNN–0.1BF	206	23.7	2	61	(251)
0.8KNN–0.2Sr(Sc0.5Nb0.5)O3–0.5%ZnO	400	∼11.6	2.6	73.2	(250)
K0.14Bi0.12Na0.5NbO3–1 mol % CuO	300	29	2.89	80	(254)
0.85KNN-0.15 Bi(Zn0.5Zr0.5)O3	325	30	3.5	86.8	(255)
0.9KNN–0.1BMN–1.0 mol % CuO	400	∼21	4.02	57.3	(66)
0.85KNN–0.15ST	400	∼26	4.03	∼52	(92)
0.9KNN–0.1BMN	300	∼34	4.08	∼62.7	(91)

^at of the bulk ceramics is commonly >0.1 mm.

3.1.2.4. BiFeO₃-Based Ceramics

BF-based ceramics are best known as multiferroics but have also been explored for high-temperature ferroelectric and piezoelectric applications due to their high T_C and large spontaneous polarization.^256−259 Compared with other lead-free ceramics, BF-based were not initially considered as good candidates for energy storage applications due to their high leakage current.²⁶⁰p-type electrical conductivity due to the volatilization of Bi and the variation of Fe valence states has been reported frequently in BF system, which limits the BDS and restricts energy density.^261−263 However, the introduction of dopants and alloying with end-members suppresses the leakage current, reduces the electrical conductivity, increases intrinsic E_g and induces transitions from a FE to RFE state. The energy storage properties of BF-based materials are summarized in Table 7.

Table 7. Energy Storage Properties of BF-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
3 mol % Nb2O5–0.65BF–0.35BT	90	19	0.71		(265)
0.61BF–0.33BT–0.06Ba(Mg2/3Nb1/3)O3	125	32.3	1.56	75	(269)
0.61BF–0.33BT–0.06La(Mg0.5Ti0.5)O3	130	33.3	1.66	82	(267)
0.6BF–0.34BT–0.06Sr(Al0.5Nb0.5)O3	150	35	1.75	81	(270)
15%Nd–0.70BF–0.30BT	170	30.7	1.82	41.3	(90)
0.65BF-0.30BT–0.05Ba(Zr0.5Zn0.5)O3	190	32	2.04	54	(34)
0.65BF–0.30BT–0.05Bi(Zn2/3Nb1/3)O3	180	32.8	2.06	53	(266)
0.56BF–0.30BT–0.14AN	195	26	2.11	84	(268)
0.60BF–0.30BT–0.10Nd(Mg2/3Nb1/3)O3	220	24	2.4	77	(34)
0.7(0.67BF–0.34BT)–0.3(Sr0.7Bi0.2)TiO3	180	37	2.4	90.4	(271)
0.62BF–0.3BT–0.08Nd(Zr0.5Zn0.5)O3	240	26	2.45	72	(34)
0.6BF–0.34BT–0.06Ba(Zn0.5Ta0.5)O3	160	41	2.56	71	(272)
0.67Bi0.9Sm0.1FeO3–0.33BT	200	48	2.8	55.8	(273)
0.25Bi0.83Sm0.17Fe0.95Sc0.05O3−0.75[0.85BT-0.15Bi(Mg0.5Zr0.5)O3]	206	36	3.2	92	(274)
0.61BF–0.33BST–0.06La(Mg2/3Nb1/3)O3	230	36.5	3.38	59	(275)
0.57BF–0.30BT–0.13 Bi(Li0.5Nb0.5)O3	280	30	3.64	74	(276)
0.57BF–0.33BT–0.10NN	360	51	8.12	90	(42)
0.5BF–0.4ST–0.1BMN-0.03Nb	460	45	8.2	74.1	(45)

^aThe t of the bulk ceramics is commonly >0.1 mm.

BF–BT-based materials are purported as potential energy storage compositions due to their excellent BDS and high P_max.^34,45,90 Undoped BF–xBT ceramics exhibit optimized ferroelectric and piezoelectric properties in a mixed-phase region of 0.25≤ x ≤ 0.35.²⁶⁴ The majority of studies have focused on this region and modified compositions either by (i) A and/or B-site chemical doping, including Nd, Nb, Zn_0.5Zr_0.5, Zn_2/3Nb_1/3, and Li_0.5Nb_0.5 or (ii) alloying with a third end-member, such as Nd(Zn_0.5Zr_0.5)O₃, Nd(Mg_2/3Nb_1/3)O₃, La(Mg_0.5Ti_0.5)O₃, Ba(Mg_1/3Nb_2/3)O₃, AN, and NN.^{34,43,90,265−269}

The importance of electrical homogeneity was first postulated in 2019 by Wang, Reaney, and co-workers as a key factor to optimize BDS as well as W_rec.³⁴ Two electrical components with similar capacitance value, C∼ 8 × 10^–10 F cm^–1 and 1.4 × 10^–9 F cm^–1, were found in 0.7BF–0.3BT ceramics, as illustrated Figure 17a,b. However, the associated resistances, R₁ ∼ 3.8 kΩ cm and R₂ ∼ 1.3 MΩ cm, were vastly different. The presence of a large volume fraction of conductive cores (R₁) easily led to electrical breakdown at lower electrical field. In contrast, only one single electrical component with C ∼ 1.87 × 10^–9 F cm^–1 and resistivity of ∼2.3 MΩ cm was observed for 8% Nd(Zr_0.5Zn_0.5)O₃ (NZZ)-doped BF–BT ceramics (Figures 17c,d). The conductive electrical component was effectively eliminated by forming a solid solution with NZZ, inhibiting the formation of conductive pathways at higher electric field (>500 kV cm^–1). As a result, optimized W_rec ∼ 2.45 and 10.5 J cm^–3 were realized at E_max ∼ 240 and 700 kV cm^–1 for 0.08NZZ-BF–BT bulk ceramics and MLs, respectively.

Figure 17.

Combined Z′′ and M′′ spectroscopic plots at 275 °C for (a) BF–BT and (b) BF–BT–0.08NZZ. Unipolar P–E loops of BF–BT–0.08NZZ (c) bulk ceramics and (e) ceramic MLs, with cross-section SEM image as shown in inset figure. Calculated energy storage properties of BF–BT–0.08NZZ (d) bulk ceramic and (e) ceramic MLs. Reproduced from ref (34). Copyright 2019 Royal Society of Chemistry.

Recently, superior energy density through tailored dopant strategies was achieved in BF–ST–xNb–yBMN ceramics, by promoting electrical homogeneity, enhancing E_a and suppressing the p-type conduction, all of which resulted in significantly enhanced BDS. The ceramic without Nb (BF–ST–0.06BMN, x = 0) exhibits a broadened arc in Z* at room temperature (Figure 18a) with at least two electrical components observed in combined Z′′ and M′′ spectroscopic plots (Figure 18b). By donor doping Nb on the B-site, only one ideal semicircle in the Z* plots with a single Debye peak at the same frequency in both Z′′ and M′′ spectroscopic plots was observed for all doped samples (Figures 18c,d), suggesting electrical homogeneity. Nb doping suppresses the formation of Fe⁴⁺ associated with the loss of Bi₂O₃ or/and V_O^.. during ceramic processing, thus reducing electrical conductivity with respect to x = 0 by several orders of magnitude, coupled with enhanced E_a (Figure 18e). A reduction in Seebeck coefficient from ∼ 600 μV K^–1 to zero indicates a commensurate decrease in charge carrier concentration as Nb concentration increases. At x = 0.03, the BDS increases to 360 kV cm^–1 (Figures 18g) which is insufficient on its own to optimize W_rec but when combined with an increase in BMN (y) concentration to reduce polar coupling, results in W_rec ∼ 8.2 J cm^–3 at E_max ∼ 460 kV cm^–1 for BF–ST–0.03Nb–0.1BMN ceramics, Figure 18h.⁴⁵

Figure 18.

(a) Z* plots and (b) Combined Z′′ and M′′ spectroscopic plots of 0.54BF–0.4ST–0.06BMN–xNb (x = 0); (c) Z* plots and (d) Combined Z′′ and M′′ spectroscopic plots of x = 0.01–0.05; (e) Arrhenius plots, (f) Seebeck coefficients, and (g) unipolar P–E loops under E_max of x = 0.01–0.05. (h) W_rec and η of (0.6–y)BF–0.4ST–0.03Nb–yBMN. Reproduced from ref (45). Copyright 2020 Royal Society of Chemistry.

3.1.2.5. Na_0.5Bi_0.5TiO₃–Based Ceramics

NBT-based ceramics are promising candidates of lead-free dielectrics due to their high P_max and T_c. However, their large hysteresis and low BDS are not ideal for high energy density capacitor applications.^277−284 Attempts to improve their properties generally fall into the following approaches: (i) doping on the A-site (Ba, Sr, K, Li, La, Dy, Nd)^{222,239,243,285−296} and B-site (commonly Nb)²⁹⁷ and codoping (K,Sr/Nb; K,La/Zr; Li,K,Sr/Ta,Nb; K,Mg/Nb; Ba/Nb; Ba/Sn; Ba/Sn,Zr; Ba/Ta; Ba/Zr; Ba,Ca/Zr; Ba,K,Ca/Nb,Zr; Ba/Mg,Nb; Ba/Mg; Ba,La/Al,Nb; Ba,Sr/Yb,Nb; Ba/Hf; La/Al; La,Ba/Nb; Sr/Sn; Sr/Zr; Sr/Mg; Sr/Mg,Nb);^{194,298−321} (ii) forming solid solution with other end-members, such as AN, NN, and SBT;^{297,322−325} (iii) using additives such as MnO, Fe₂O₃, MgO, SnO₂, ZnO, CaO, and ZrO₂^{62,74,238,326−334} and (iv) employing different processing methods such as hot-pressing⁵⁹ and synthesis using sol–gel derived powders.^{69,71,74,335,336} The energy storage properties of NBT-based materials are summarized in Table 8.

Table 8. Energy Storage Properties of NBT-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
0.95(0.76NBT–0.24ST)-0.05AN	120	43.5	2.03	61.8	(324)
Bi0.41Na0.35Sr0.21TiO3	135	∼35.8	2.04	82.4	(290)
0.75Na0.25Sr0.5Bi0.25TiO3–0.25MgO	200	∼35	2.06	84	(317)
0.8(0.775NBT–0.225BaSnO3)–0.2BaZrO3	245	20	2.08	88.8	(315)
0.85NBT–0.15BaHfO3	175	29	2.1	66.1	(310)
0.775NBT–0.225BaSnO3–5 wt %MgO	215	∼28	2.13	67.8	(326)
Sn–0.45ST–0.2NBT–0.35BT	240	∼21	2.25	79.51	(328)
0.9NBT–0.1Li2TiO3	200	∼25	2.3	74.2	(291)
(0.5NBT–0.5ST) −1%SnO2	180	37.37	2.35	∼80	(238)
(Na0.5Bi0.5)0.8Ba0.2Ti0.8Sn0.8O3	195	∼34	2.35	71.04	(338)
{Bi0.5[(Na0.8K0.2)0.9Li0.1]0.5}0.96Sr0.04(Ti0.975Ta0.025)O3	143	∼43	2.42	∼64	(339)
0.55NBT–0.45(Bi0.2Sr0.7TiO3)	200	25	2.5	95	(222)
0.94(BNT–Bi0.2Sr0.7TiO3)–0.06KNN	180	37	2.65	84.6	(298)
0.95(0.6ST–0.4NBT)–0.05Zr	285	∼25	2.84	71.54	(74)
0.85(0.95NBT–0.05SrZrO3)–0.15NN	210	∼30	2.93	72	(325)
0.96(0.65NBT–0.35Sr0.85Bi0.1TiO3)–0.04NN	220	50.46	3.08	81.4	(323)
0.6(Bi0.51Na0.47)TiO3–0.4Ba(Zr0.3Ti0.7)O3	280	∼25	3.1	91	(299)
0.9(0.76NBT−0.24NN)−0.1SBT	200	43	3.12	75	(340)
0.93NBT–0.07LaAlO3	210	43	3.18	60	(320)
(Na0.25Bi0.25Sr0.5)(Ti0.8Sn0.2)O3	310	26.5	3.4	90	(321)
0.85(Na0.5Bi0.5)0.7Sr0.3TiO3−0.15BMN	250	38	3.45	88.01	(316)
0.95(0.6 Bi0.5Na0.5TiO3–0.4Sr0.7Bi0.2TiO3)–0.05AN	246	41	3.62	89	(322)
0.65(0.84NBT–0.16K0.5Bi0.5TiO3)–0.35(Bi0.2Sr0.7TiO3)	350	∼33.99	4.06	87.3	(288)
0.55NBT–0.45SBT	315	19.1	4.14	92.2	(295)
0.90(Na0.5Bi0.5)0.7Sr0.3TiO3–0.10 Bi(Ni0.5Sn0.5)O3	270	47	4.18	83.64	(341)
0.75Bi0.58Na0.42TiO3–0.25ST	535	41	5.63	94	(342)
0.78NBT–0.22NN	390	45	7.02	85	(297)
0.62NBT–0.30SBT–0.08BMN	470	48	7.5	92	(337)

^at of the bulk ceramics is commonly >0.1 mm.

Notably, Li and co-workers reported that 0.55NBT–0.45(Sr_0.7Bi_0.2)TiO₃(SBT) achieved W_rec of 2.5 and 9.5 J cm^–3 with η > 90% for bulk ceramic and MLs at 200 and 720 kV cm^–1 (Figure 19a,b), respectively.²²² Superior W_rec ∼ 7.02 J cm^–3 and η ∼ 85%, were also reported for 0.78NBT–0.22NN ceramics at E_max ∼ 360 kV cm^–1, Figures 19c,d, with <10% variation from 25–250 °C and from 0.1 to 100 Hz.²⁹⁷

Figure 19.

(a) Unipolar P–E loop and the current–E curve for NBT–0.45SBT bulk ceramics. (b) W_rec and η for NBT–0.45SBT ceramic MLs; inset SEM image of the ceramic MLs (c) Bipolar P–E loops and (c) calculated W_rec and η of 0.78NBT–0.22NN ceramic.^222,297 (a, b) Reproduced with permission from ref (222). Copyright 2018 John Wiley and Sons; (c, d) Reproduced with permission from ref (297). Copyright 2019 Royal Society of Chemistry.

Recently, Ji and co-workers³³⁷ proposed that the key factors for designing an ideal RFE with high energy density were as follows: (i) utilization of a highly polar base system (e.g., NBT); (ii) disruption of long-range polar coupling through forming solid solutions with, e.g., SBT and BMN without sacrificing average ionic polarizability, Figure 20a, and (iii) simultaneously inducing or retaining electrical homogeneity with a highly resistive single component in IS (∼250 kΩ cm at 660 °C), Figure 20b. These factors combined to give E_max ∼ 470 kV cm^–1, W_rec ∼ 7.5 J cm^–3, and η ∼ 92% for 0.62NBT–0.30SBT–0.08BMN, Figure 20c,d.

Figure 20.

(a) Tolerance factor and average ionic polarizability per unit cell of NBT–SBT–xBMN as a function of BMN concentration. (b) Combined Z′′ and M′′ spectroscopic plots for NBT–SBT–0.08BMN ceramics at 660 °C. (c) P–E loops at the E_max, and (d) W_rec and η for NBT–SBT–xBMN ceramics. Reproduced with permission from ref (337). Copyright 2021 Elsevier.

3.1.2.6. AgNbO₃–Based Ceramics

AFEs have long been considered as the prime candidate for energy storage capacitors due to their large P_max and small P_r. There are only a handful of lead-free AFE systems, with AN showing particular promise because it possesses a large saturation polarization of 52 μC cm^–2 under an E_max ∼ 220 kV cm^–1.³⁴³ Recent research on AN ceramics has focused on stabilizing the AFE phase so that switching field is moved to higher fields while simultaneously optimizing P_max.^344,345

There have been a number of recent studies on AN focusing on (i) substitution of aliovalent B-site oxides such as MnO₂ and WO₃;^346,347 (ii) doping Ba, Sr, Ca, Bi, La, Sm, and Gd on the A-site^348−356 often with isovalent Ta doping on the B-site;^357−359 and (iii) forming solid solutions with end-members, such as BiMnO₃ and Bi(Zn_2/3Nb_1/3)O₃.^360,361 Most dopants reduce the G of AN which maximizes BDS but delay the onset of the AFE-FE transition to higher field while simultaneously narrowing hysteresis in the induced FE phase. W_rec of 4.4, 4.5, and 5.2 J cm^–3 with η of 70, 63, and 69.2% has been obtained for La, Gd, and Sm A-site doped AN ceramics,^{348,351,355,359,362} respectively, and B-site Ta-doped AN was reported to exhibit W_rec of 4.2 J cm^–3 with η of 69% (Figure 21a,b).³⁵⁷ A-site doping with ions smaller in radius than Ag is suggested to decrease tolerance factor and enhance AFE stability, while donor doping is compensated by A-site vacancies which reduce antipolar and polar coupling of the AFE and field induced FE phases, respectively. Some authors postulate that substituting B-site ions with a lower polarizability than Nb also stabilizes the AFE phase and moves the switching field higher.^{348,350,359,363} The underlying principles are schematically represented in Figure 21c. The energy storage properties of AN-based materials are summarized in Table 9.

Figure 21.

(a) Bipolar P–E loops of AN and Ag(Nb_0.85Ta_0.15)O₃ ceramics. (b) Energy storage performance of Ag(Nb_1–xTa_x)O₃ ceramics prior to their breakdown. (c) Schematic of the underlying principles for enhancing energy storage property in AN-based materials. (a, b) Reproduced with permission from ref (357). Copyright 2017 John Wiley and Sons; (c) Reproduced with permission from ref (359). Copyright 2019 Royal Society of Chemistry.

Table 9. Energy Storage Properties of AN-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
AN	150	∼34	2.0	46	(345)
AN	175	∼33	2.1	40–50	(343)
Ag0.96Ba0.02NbO3	∼180	∼34	2.3	46	(356)
0.6 mol % BiMnO3–AN	175	∼36	2.4	54	(360)
0.1 wt % Mn–AN	150	∼37.2	2.5	57	(346)
Ag0.91Bi0.03NbO3	200	∼30	2.6	86	(353)
Ag0.90Sr0.05NbO3	190	∼38	2.9	56	(354)
Ag0.94La0.02NbO3	230	∼28	3.12	63	(352)
0.3 wt % Mn-doped Ag0.97La0.01NbO3	142	37.8	3.2	62	(351)
0.1 wt % W-AN	200	∼42.5	3.3	50	(347)
AN–0.03NBT	220	33	3.4	62	(364)
Ag0.90Ca0.05Nb0.95Ta0.05O3	210	37	3.36	58	(363)
Ag0.92Ca0.04NbO3	220	∼37.6	3.55	63	(349)
Ag(Nb0.8Ta0.2)O3	270	∼30	3.7	∼65	(358)
AgNb0.85Ta0.15O3	233	∼35.1	4.2	69	(357)
2 mol % La-doped AN	273	∼30	4.4	70	(350)
Ag0.88Gd0.04NbO3	290	∼32.5	4.5	∼63	(355)
Ag0.94Sm0.02NbO3	310	31	4.5	63	(362)
0.99AN–0.01 Bi(Zn2/3Nb1/3)O3	220	46.8	4.6	57.5	(115)
(Sm0.02Ag0.94)(Nb0.9Ta0.1)O3	280	∼36	4.87	63.5	(359)
Sm0.03Ag0.91NbO3	290	∼36	5.2	69.2	(348)
AgNb0.45Ta0.55O3	460	∼29	6.3	90	(365)
Ag0.97Nd0.01Nb0.80Ta0.20O3	370	38	6.5	71	(366)
Ag0.76La0.08NbO3	476	33	7.01	77	(367)

^at of the bulk ceramics is typically >0.1 mm.

3.1.2.7. NaNbO₃-Based Ceramics

Recently, AFE NN has received attention as a potential candidate for energy storage applications. An AFE double hysteresis loop is difficult to observe in NN because (i) the energy difference between the AFE phase and field-induced FE phase is very small and (ii) E-induced FE phase is metastable. Thus, AFE behavior in NN based materials is commonly stabilized by chemical substitution with end members such as BS and CaHfO₃.^{344,368−370}

E_max > 250 kV cm^–1 and W_rec > 2.5 J cm^–3 have been reported for NN in solid solution with BMN, ST, Bi(Mg_2/3Ta_1/3)O₃, and Bi(Mg_0.5Ti_0.5)O₃ by stabilizing the AFE phase or inducing relaxor behavior,^371−374 as summarized in Table 10. Zuo and co-workers⁴² have also proposed the concept of an “AFE relaxor” to explain the energy storage properties of 0.78NN–0.22NBT ceramics. They argue that the local AFE phase transforms reversibly into an FE monoclinic phase at∼ 400 kV cm^–1, giving a large ΔP (P_max > 50 μC cm^–2 and P_rem < 5 μC cm^–2). W_rec of ∼12.2 J cm^–3 was reported with η ∼ 69%, at 680 kV cm^–1, Figure 22.⁴² However, the term “AFE relaxor” has little physical significance since an antipolar phase cannot form short-range polar features characteristic of a relaxor. 0.78NN–0.22NBT may, therefore, be better described as either a relaxor or a short-range AFE phase that undergoes a field induced transition. This intriguing behavior is interesting, but it is the large E_max (680 kVcm^–1) that is most likely responsible for the exceptional W_rec rather than the intrinsic crystal chemistry. The underpinning reasons for the large E_max most likely relate to the defect chemistry, band gap and electrical homogeneity, consistent with the key factors proposed by Ji and co-workers.³³⁷

Table 10. Energy Storage Properties of NN-Based Materials^a.

compounds	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
NN–0.09Bi(Zn0.5Ti0.5)O3	200	29	2.1	76	(375)
0.9NN–0.10BMN	300	∼23	2.8	82	(371)
0.8NN–0.2ST	323	34.5	3.02	80.7	(372)
0.78NN–0.22Ba(Mg2/3Nb1/3)O3	540	18.7	3.51	87	(376)
Na0.7Bi0.1NbO3	250	∼30	4.03	85.4	(377)
NN–MnO2	360	33	4.3	90	(378)
Na0.84Bi0.08Nb0.92Zr0.08O3	430	30	4.9	88	(67)
0.9NN–0.1 Bi(Ni0.5Sn0.5)O3	550	25	5	68	(379)
0.78NN–0.22 Bi(Mg2/3Ta1/3)O3	620	17	5.01	86.8	(373)
0.92NN–0.08 Bi(Mg0.5Ti0.5)O3	480	38	5.57	71	(374)
(Na0.91La0.09)(Nb0.82Ti0.18)O3	550	42	6.5	65.9	(380)
0.75[0.9NN–0.1 Bi(Mg0.5Ta0.5)O3]–0.25(Bi0.5Na0.5)0.7Sr0.3TiO3	800	22	8	90.4	(381)
0.76NN–0.24NBT	680	∼55	12.2	69	(42)

^at of the bulk ceramics is commonly >0.1 mm.

Figure 22.

(a) Bipolar P–E loops with corresponding current density-field (J–E) curves and (b) W_rec and η values of under different E for the 0.76NN–0.24BNT ceramic at 10 Hz (c) a comparison of W_rec, η, and E_max among the recently reported bulk ceramics; (d) temperature-dependent P–E hysteresis, (e) temperature- and frequency-dependent ε_r and (f) W_rec and η as a function of temperature for the 0.76NN-0.24NBT ceramic at 450 kV cm^–1. Reproduced with permission from ref (42). Copyright 2019 John Wiley and Sons.

3.1.3. Glass Ceramics

Glass-ceramics are composed of one or more crystallized phases (ceramics) dispersed uniformly in amorphous phase (glass). They often exhibit the combined properties of ceramics and glass depending on the induced crystalline phases and their microstructures. Glass-ceramics are prepared by melting the requisite raw materials, cooling to room temperature to form a glass, followed by two step annealing to induce crystal nucleation (approximately at the glass transition temperature, T_g) and growth > T_gFigure 23.^179,382 The microstructure of a glass ceramic is typically dominated by a largely 2D and 3D defect-free (e.g., no grain boundaries) glass phase and a uniformly distributed (provided the system undergoes homogeneous rather than heterogeneous nucleation) ceramic phase. W_rec and η are both large due to the high BDS associated with the absence of 2D and 3D defects accompanied by a near zero value of P_r. The energy storage properties of glass-based glass ceramics are summarized in Table 11.

Figure 23.

Schematic of the processing step of glass-ceramics.

Table 11. Energy Storage Properties of Glass Ceramics^a.

compounds	εr (1 kHz, 300 K)	E (kV cm–1)	Wrec (J cm–3)	ref
14.4SrO–17.6BaO–32Nb2O5–36B2O3	117	1050	5.71	(386)
25.6BaO–6.4Na2O–32Nb2O5–36SiO2	∼90	1248	∼6.2	(409)
14.3SrO–17.5BaO–31.9Nb2O5–35.8B2O3–0.5ZnO + 0.5La2O3	131	1127	7.1	(400)
20BaO–20SrO–20Nb2O5-5Al2O3–1.5B2O3–33.5SiO2 + 0.2La2O3	92.4	1326	7.2	(406)
20SrO–20BaO–10Nb2O5–10TiO2–32SiO2–8Al2O3	52.9	1817	7.73	(410)
14.3SrO–17.5BaO– 31.9%Nb2O5–35.8%B2O3–0.5ZnO + 0.5Sm2O3	143.8	1132	8.15	(402)
14.4SrO–17.6BaO–32Nb2O5–36B2O3 + 1%Yb2O3	98.3	1398	8.5	(403)
15K2CO3–15SrCO3–30Nb2O5–32SiO2–4Al2O3–4B2O3	102 (10 kHz)	1411	8.99	(407)
20BaO–20SrO–20Nb2O5–5Al2O3–1.5B2O3–33.5SiO2 + 0.05MnO2	95.8	1471	9.2	(411)
42[0.2Na2O–0.8SrO]–28Nb2O5–30SiO2	53 (100 kHz)	2074	10.09	(412)
9.6K2O–22.4BaO–32Nb2O5–36SiO2	75	1937	12.06	(393)
25.6BaO–6.4K2O–32Nb2O5–36SiO2 + 1Gd2O3	83	1818	12.14	(405)
20BaO–12K2O–32Nb2O5–36SiO2	83	1859	12.7	(413)
15.16SrO–6.736BaO–10.104K2O–32Nb2O5–28B2O3–8P2O5	85.2	1844	12.83	(395)
6.4K2O–25.6SrO–32Nb2O5–36SiO2 + 3CaF2	114	1623	13.5	(414)
31.2SrO–7.8Na2O–26Nb2O5–35SiO2	91	1941	15.2	(390)
15 Bi2O3–15Nb2O5–40SiO2–30Al2O3	100	1861	15.3	(391)
65(48SrO–12Na2O–40Nb2O5)–35SiO2	124	1669	15.3	(392)
24BaO–6Na2O–30Nb2O5–10Al2O3–30SiO2	∼70	2322	16.6	(388)
15.4Na2O–15.4PbO–23.1Nb2O5–46.2SiO2	175	1486	17	(385)
25.6(0.4SrO–0.6BaO)–6.4K2O–32Nb2O5–36SiO2	118	1828	17.45	(394)
21.25BaO–1PbO–12.75Na2O–34Nb2O5–32SiO2	154	1638	18.29	(389)
25.6BaO–3.2Na2O–3.2K2O–32Nb2O5–36SiO2	22	4433	19	(396)
21.6BaO–2.4PbO–6Na2O–30Nb2O5–10Al2O3–30SiO2	137	1848	20.7	(384)
63SiO2–12BaO–16B2O3–9Al2O3	6	12,000	38.5	(397)

^at of the bulk ceramics is commonly >0.1 mm.

As discussed above, the crystallization of glass ceramics is controlled by the annealing procedure, where the annealing temperature and time are critical for nucleation and growth of the ceramic phase, the microstructure and the properties. Generally, the volume fraction of crystalline phase increases with increasing annealing temperature and time, accompanied by an increase of ε_r and decrease of BDS. The optimized W_rec is a balance between ε_r and BDS. Chen and co-workers³⁸³ reported that tungsten bronze structured, Ba_0.27Sr_0.75Nb₂O_5.78 phase formed from the Na₂O–BaO–SrO–Nb₂O₅–SiO₂–B₂O₃ glass matrix at 800 °C and a secondary phase NaSr_1.2Ba_0.8Nb₅O₁₅ emerged when crystallization temperature exceeded 850 °C. Remarkably high BDS ∼ 1400 kV cm^–1 with ε_r of ∼50 were obtained, leading to a W_rec = 4 J cm^–3. Besides, Wang and co-workers reported ultrahigh W_rec of 20.7 J cm^–3 in BaO–PbO–Na₂O–Nb₂O₅–SiO₂–Al₂O₃ (BPNN-AS) glass ceramics at the optimized crystallization temperature of 900 °C, as shown in Figure 24.³⁸⁴ With increasing crystallization temperature from 850 to 1000 °C, the BDS decreased from 1890 to 1440 kV cm^–1 and the crystallinity increased from 64.5 to 97.3% (Figure 24b). Similar results were also reported in other glass ceramic systems, including Na₂O–PbO–Nb₂O₅–SiO₂,³⁸⁵ SrO–BaO–Nb₂O₅–B₂O₃,³⁸⁶ K₂O–SrO–Nb₂O₅–SiO₂–Al₂O₃–B₂O₃,³⁸⁷ BaO–Na₂O–Nb₂O₅–SiO₂–Al₂O₃,³⁸⁸ BaO–PbO–Na₂O–Nb₂O₅–SiO₂,³⁸⁹ SrO–Na₂O–Nb₂O₅–SiO₂,³⁹⁰ and Bi₂O₃–Nb₂O₅–SiO₂–Al₂O₃.³⁹¹

Figure 24.

(a) X-ray diffraction (XRD) patterns of BPNN-AS glass ceramics annealed from 850 to 1000 °C; (b) ε_r and M₂NaNb₅O₁₅ + NN phase proportion with increasing annealing temperature; (c) BDS Weibull distribution plots; (d) W_rec of 900 °C annealed BPNN-AS glass ceramics, compared with other kinds of ferroelectric glass ceramics. Reproduced with permission from ref (384). Copyright 2018 Royal Society of Chemistry.

Each constituent oxide in the glass matrix has an important effect on the crystal phase, microstructure, BDS, and energy storage properties. For example, SiO₂ is an important and active studied constituent oxide in glass matrix. With increasing SiO₂ content, ε_r of SrO–Na₂O–Nb₂O₅–SiO₂ (SNN-Si) glass ceramics first increased and then decreased as shown in Figure 25a, which was attributed to the change of volume fraction Sr₆Nb₁₀O₃₀ (Figure 25b). The optimal ε_r of 120 and BDS ∼ 1700 kV cm^–1 were obtained with 35 mol % SiO₂ (Figures 25a,c), resulting in the highest theoretical W_rec of 15.2 J cm^–3.³⁹² Wang and co-workers reported that, as K₂O concentration increased in K₂O–BaO–Nb₂O₅–SiO₂ glass ceramics, grain boundary R and activation energy decreased, indicating the decrease of interfacial polarization, leading to the enhancement of BDS to ∼1900 kV cm^–1 and W_rec ∼ 12 J cm^–3.³⁹³ They also reported that substitution of Sr for Ba in SrO–BaO–K₂O–Nb₂O₅–SiO₂ led to the formation of solid phase Sr_0.5Ba_0.5Nb₂O₆ and improvement of dielectric properties.³⁹⁴ The highest BDS of ∼1800 kV cm^–1 and W_rec of 17.5 J cm^–3 were achieved with Sr = 0.4 due to a uniform and dense microstructure and lower interfacial polarization. Li and co-workers reported substitution of K with Ba in SrO–BaO–K₂O–Nb₂O₅–B₂O₃–P₂O₅ glass ceramics transformed Ba_0.5Sr_0.5Nb₂O₆ to a solid solution of K_2xyBa_(1-x)ySr_5–yNb₁₀O₃₀ and then KSr₂Nb₅O₁₅ phase, leading to a decrease in ε_r.³⁹⁵ A maximum theoretical W_rec of 12.8 J cm^–3 was obtained under BDS ∼ 1800 kV cm^–1, along with dielectric loss <0.3%. Liu and co-workers studied the effect of R₂O (R = Li, Na, K) on the phase structure, dielectric properties and BDS in BaO–R₂O–Nb₂O₅–SiO₂ glass ceramics, where the highest W_rec ∼ 19 J cm^–3 was achieved with ε_r ∼ 22 and superior BDS of ∼4400 kV cm^–1 in composition of BaO–Na₂O–K₂O–Nb₂O₅–SiO₂.³⁹⁶

Figure 25.

(a) Dielectric properties, (b) XRD patterns, and (c) BDS of SNN-Si glass ceramics as a function of SiO₂ concentration (β). Reproduced with permission from ref (392). Copyright 2017 Elsevier.

Compared with alkali based compositions, alkali-free glass compositions were found to deliver lower dielectric loss and fewer defect microstructure. For example, Smith and co-workers reported both ultrahigh E_max ∼ 12000 kV cm^–1 and W_rec ∼ 35 J cm^–3 in BaO–B₂O₃–Al₂O₃–SiO₂ glass.³⁹⁷ The effect of Al/Si ratio on the modification of the microstructure and properties of SrO–BaO–Nb₂O₅–SiO₂–Al₂O₃ glass ceramics was also studied by Xiu and co-workers.³⁹⁸

Additionally, rare-earth oxides, such as La₂O₃,^399−401 Sm₂O₃,⁴⁰² Yb₂O₃,^403,404 and Gd₂O₃,⁴⁰⁵ are also commonly substituted into glass formulations for energy storage applications. Rare-earth oxides are mainly reported to act as nucleating agents^399,402 or crystal growth inhibitors.^400,406 Zhang and co-workers revealed that La₂O₃ leads to a homogeneous microstructure in the BaO–SrO–TiO₂–Al₂O₃–SiO₂ glass-ceramics which improved BDS ∼ 1600 kV cm^–1 and W_rec ∼ 3.2 J cm^–3 (2.5 times of the glass-ceramics without La₂O₃).³⁹⁹ Zheng and co-workers reported that 0.5 mol % La₂O₃ in SrO–BaO–Nb₂O₅–B₂O₃–ZnO glass-ceramics also optimized ε_r (∼130) and W_rec ∼ 7.1 J cm^–3 through achieving a BDS ∼ 1100 kV cm^–1 due to a reduction in crystallite size and precipitation of high ε_r phase, Sr_0.5Ba_0.5Nb₂O₆.⁴⁰⁰ A similar effect was reported for Sm₂O₃ by Chen and co-workers in the SrO–BaO–Nb₂O₅–B₂O₃–ZnO glass ceramics with W_rec of 8.2 J cm^–3 at 1100 kV cm^–1.⁴⁰² Moreover, Yb₂O₃ is reported to eliminate the impurity phases and form a uniform microstructure in BaO–SrO–TiO₂–Al₂O₃–B₂O₃–SiO₂ glass-ceramics, leading to W_rec of 3.5 J cm^–3, ∼ 1.8 times higher than undoped compositions.⁴⁰⁴

Apart from the conventional annealing, novel methods such as microwave treatment have been reported to improve the energy storage properties of glass ceramics. Zhang and co-workers found that microwave treatment restrained the formation of the dendritic microstructure in Ba_xSr_1-xTiO₃−(Ba−B−Al−Si−O) (BST−BBAS) glass-ceramics (Figure 26), leading to the improvement of BDS from 1200 kV cm^–1 to 1500 kV cm^–1, corresponding to W_rec of 2.8 J cm^–3 (950 °C anneal).⁴⁰⁸ Xiao and co-workers further reported that the precipitation of impurity phases in the K₂O–SrO–Nb₂O₅–SiO₂–Al₂O₃–B₂O₃ glass-ceramics was limited by controlling the crystallization time using microwave sintering, with optimum ε_r, BDS of 1400 kV cm^–1 and maximum theoretical W_rec (∼9 J cm^–3) obtained after 10 min.⁴⁰⁷

Figure 26.

SEM images of BST-BBAS samples annealed at 950 °C by (a) conventional method and (b) microwave treatment.⁴⁰⁸ SEM images of BNNS samples annealed at 1000 °C by (c) conventional method and (d) microwave treatment. (e, f) Temperature dependence of dielectric properties of BNNS samples. The BDS plot is inset in (f).⁴⁰⁹ (a, b) Reproduced with permission from ref (408). Copyright 2014 Elsevier; (c−f) Reproduced with permission from ref (409). Copyright 2017 Elsevier.

3.1.4. Summary of State-of-the-Art in Ceramics

The debate over whether lead-free electroceramics can replace their lead-based counterparts has been ongoing for over two decades. Lead based compositions generally outperform their lead-free counterparts on most metrics. Moreover, lead-free compositions are disparate with a large number of different formulations potentially required to cover the properties achieved with essentially doped PZT. Provided, however, that the performance, reliability and cost of lead-free are competitive with PZT, it is highly likely that lead-free electroceramics will begin to replace their lead-based equivalents and attain large scale production in the coming years as a consequence of environmental legislation.^31,103,415

Of all the applications, lead-free high energy density capacitors are the most likely to see large-scale production since (i) the performance of lead-free compositions is approaching that of lead-based; (ii) reduction in intrinsic electrical properties may be compensated by increasing the BDS often through decreasing layer thickness (see section 3.1.2); and (iii) the capacitor industry is dominated by lead-free BT-based MLCCs, and thus, there is an expectation that the related products will not contain lead.^{34,45,90,209,276,416} This latter statement does not hold for piezoelectric ceramics market which is dominated by PZT and its derivatives.^{174,417−419}

The energy storage performances, E_max, ΔP, W_rec, and η, for lead-based and lead-free ceramics are summarized and plotted in Figure 27 (note: glass ceramics are not included). A comparison of W_rec vs E_max for different lead-based/lead-free bulk ceramics is displayed in Figure 27a. Lead-based bulk ceramics have the advantage of both high E_max (up to ∼400 kV cm^–1) and W_rec (up to ∼12 J cm^–3) with respect to lead-free candidates. NN-based ceramics currently offer the highest W_rec under high E (>350 kV cm^–1) for lead-free compositions, followed by AN-, NBT-, BF-, and KNN-based materials. BT and ST-based ceramics display the lowest W_rec in spite of their high E_max of ∼450 kV cm^–1, but they are perhaps the most appealing dielectrics commercially since they are the current basis of MLCC production.

Figure 27.

Comparison of (a) E_max vs W_rec; (b) ΔP vs W_rec; and (c) W_rec vs η for lead-based/lead-free bulk ceramics. *t of the bulk ceramics is commonly >0.1 mm.

ΔP vs W_rec is compared in Figure 27b. Lead-, NBT-, and BF-based materials exhibit extraordinarily high ΔP (up to ∼60 μC cm^–2), followed by AFE NN- and AN-based (up to ∼40 μC cm^–2) and KNN-based materials (up to ∼35 μC cm^–2) with the lowest (up to ∼25 μC cm^–2) for BT- and ST-based materials. Bicontaining electroceramics such as NBT and BF, have been heavily studied recently as potential lead-free electroceramic materials due to their large polarization.^{260,266,281,282,284,420−423} A high ΔP (60 μC cm^–1) is obtained in NBT and BF based by reducing the P_r through chemical substitution with other perovskite end members to form a relaxor with an ultraslim P–E loop. To some extent therefore, the advantage of a high intrinsic polarization end member such as BF is weakened. Intermediate ΔP ∼ 40 μC cm^–2 values are observed for AFEs such as AN- and NN-based materials but P_max is often limited as compositions exhibit polarization saturation as a function of applied field.

Figure 27c compares W_rec vs η for a wide range of compositions. ST-based materials display the best η (∼90%) due to their linear-like dielectric behavior. For BT-, NBT-, and lead-based materials, η varies with the material composition since it is a function of many factors. Dielectric loss associated with defects such as V_O^.. play a role but primarily at high field and high frequency, energy is dissipated during the transition to a field induced long-range ordered state which is manifested by the opening of the P–E loop. If the transition is smeared over for the operational range through alloying or doping to create a so-called “weakly coupled relaxor state”, η > 90% can be achieved.^43,209,337

High leakage current and electrical conductivity are considered as major challenges in BF and KNN-based materials but are addressed by appropriate doping, e.g. donor doping to mitigate p-type conductivity in BF-based ceramics.⁴⁵ In addition, AFE based materials generally suffer from opening of the polarization loop above switching field to form a field induced FE phase which is detrimental to η. Compositional modifications to AN and NN ceramics aim not only to push the AFE-FE transition to higher field and stabilize the AFE phase but also to disrupt the long-range ordering in the field induced FE phase, thereby creating a slimmer portion of the P–E loop (higher η) than that being observed in unmodified materials.³⁶⁶

Apart from the parameters discussed above (E_max, ΔP, W_rec, and η), temperature and frequency stability are also important for practical applications. In the future, high energy density ceramic capacitors will be placed closer to the core engine electronics to optimize the equivalent circuit resistance. Therefore, the temperature requirement for energy storage ceramics is anticipated to increase. According to the white paper “Multilayer Ceramic Capacitors for Electric Vehicles” published by Knowles capacitors in 2017,⁴²⁴ the explosive development of EVs has prompted the appearance of new 200 °C-stable C0G type I dielectric ceramic capacitor on the market. However, these materials still do not fulfill the required high power/voltage, energy density, and temperature requirements (∼250 °C) to facilitate use near-engine. Better frequency stability from 100 Hz to 100 kHz is required to reduce power fluctuations when capacitors are used for DC/DC conversion for battery charging and DC/AC conversion for propulsion.^425−429 Enhanced frequency stability also enables the capacitor to be compatible with diodes and thyristors for power switching and control.⁴³⁰

The temperature and frequency stabilities of many high energy density ceramics are evaluated, as shown in Figure 28. Only two compositions to date deliver W_rec > 3.5 J cm^–3 up to 250 °C, 0.57BF–0.33BT–0.1NN⁴³ and 0.78NBT–0.22NN.²⁹⁷ Most other compositions either do not sustain or do not have properties reported >200 °C. Typical issues associated with operating at higher temperature include, widening of the P–E loop or early breakdown due to high leakage current and electric field and/or temperature-induced phase transitions. High leakage currents above 200 °C typically arise from oxygen vacancy diffusion.⁴⁵

Figure 28.

(a) Temperature-^{42,43,130,131,208,242,297,337,348,352,357,359} and (b) frequency-dependent W_rec for some reported electroceramic materials for high energy density capacitors.^{276,295,297,337,416}

Most compositions have been shown to deliver W_rec at a few hundred Hz but higher frequencies (>kHz) are rarely reported. Wang and co-workers, for example, discussed frequency stability from 10^–2 to 10² Hz for 0.57BF-0.3BT–0.13BLN with W_rec ∼ 8 J cm^–3 and η ∼ 81% at 400 kV cm^–1. All electroceramics for capacitors appear to deliver a charging–discharging speed at or faster than 1 μs. Short times of τ_0.9 ∼ 0.15 μs (90% of energy discharge in 0.15 μs) were reported by Li and co-workers in BT based ceramics at 150 kV cm^–1197 while Qi et al. described even faster charging–discharging speeds (τ_0.9 ∼ 97 ns) in BF-based ceramics at 200 kV cm^–1.⁴² Increasing attention has also been focused on fatigue-resistant behavior as a performance metric in practical applications with 10⁴ cycles (W_rec variation <10%) reported in BF-based ceramic multilayers from 20 to 100 °C, coupled with a low value of electrostrain (<0.03%).²⁷⁶

3.2. Ceramic Multilayers and Films

3.2.1. Ceramic Multilayers

Ceramic MLs are fabricated by a series of processing steps which include slurry preparation, tape-casting, screen printing, lamination, cosintering, and termination, as shown in Figure 29.^{18,176,422,431,432} This fabrication technology is a powder-based approach that accommodates scale-up from laboratory research to commercial manufacturing. The market of ceramic MLs ∼ $5.3 billion in 2017 but will reach ∼ $7.8 billion by 2024 for electronic applications, including but not limited to mobile phones, laptops and motor vehicles.⁴³³ Advanced high energy density ceramic MLs, based on AFEs and RFEs materials, are being developed to facilitate power electronics within hybrid electric vehicles which require higher W_rec and operating temperature. Simultaneously, research into low cost internal electrodes is required so that the highest performant ceramics can be developed. The energy storage properties for different ceramic MLs are summarized in Table 12.

Figure 29.

Ceramic MLs fabrication process (MLs cofire technology). Reproduced with permission from ref (18). Copyright 2010 IEEE.

Table 12. Summary of the Energy Storage Performance of Ceramic MLs.

materials	t (μm)	electrode	E (kV cm–1)	ΔP (μC cm–2)	Wrec J cm–3	η (%)	ref
[0.94(0.75NBT–0.25NN)–0.06BT]–0.1CaZrO3	30	70Ag/30Pd	120		0.35	77	(440)
BT + BT@SiO2 *layer structure	10/20		301.4	18	1.8	71.5	(441)
ST + Li2CO3)/(0.94NBT–0.06BT) *layer structure	50/12		237	30	2.41	68	(442)
(Pb0.88Ba0.05La0.02Dy0.04) (Zr0.68Sn0.27Ti0.05)O3	11	95Pd/5Ag	300	28	2.7	67.4	(435)
0.6NBT–0.4ST	27	75Ag/25Pd	270		2.83	85	(443)
Pb(Zr0.95Ti0.05)0.98Nb0.02O3	32				3		(434)
Nb and Mn co-doped 0.9BT–0.1NBT	30	60Ag/40Pd	480	18	3.33	80	(444)
Ca(Zr0.80Ti0.20)O3	10	Pt	1500	6	4		(445)
BT–0.12 Bi(Li0.5Ta0.5)O3	30	Pt	466	25	4.05	95.5	(203)
BT–0.12 Bi(Li0.5Nb0.5)O3	29	Pt	450	25	4.5	91.5	(198)
0.7BT–0.3BiScO3	25	Pt	730	24	6.1		(33)
15% Nd doped BF–BT	33	Pt	540	39	6.74	77	(90)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3	11	60Ag/40Pd	790	27	7.8	88	(446)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3	11	60Ag/40Pd	750	27	8.13	95	(437)
NBT–0.45(Sr0.7Bi0.2)TiO3	30	Pt	720	30	9.5	92	(222)
BF–0.3Ba0.8Sr0.2TiO3–0.06La(Mg2/3Nb1/3)O3	8	Pt	720	43	10	77	(416)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3	5	60Ag/40Pd	1047	30	10.12	89.4	(438)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3	9	60Ag/40Pd	1000	30	10.5	93.7	(447)
BF–0.3BT–0.08Nd(Zr0.5Zn0.5)O3	17	Pt	700	34	10.5	87	(34)
Pb0.98La0.02(Zr0.7Sn0.3)0.995O3	20	Pt	560	50	12.6	80	(436)
BF–0.3BT–0.13 Bi(Li0.5Nb0.5)O3	8	Pt	953	45	13.8	81	(276)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3	4.8	Pt	1500	36	14.1	69.7	(448)
0.50BF–0.40ST–0.10BMN	8	Pt	1000	50	15.8	80	(45)
NBT–0.30(Sr0.7Bi0.2)TiO3–0.08BMN	8	Pt	1000	50	18	93	(337)
BT−0.13Bi[Zn2/3(Nb0.85Ta0.15)1/3]O3@SiO2	4.7	60Ag/40Pd	1755	35	18.24	94.5	(89)
⟨111⟩-textured NBT–0.30SBT	20	Pt	1030	>65	21.5	80	(439)

Several lead-based AFE ceramic MLs have been reported using different internal electrodes. A giant power density ∼2000 kW cm^–3 and W_rec∼ 3 J cm^–3 was obtained in Pb(Zr_0.95Ti_0.05)_0.98Nb_0.02O₃ MLs using Pt as internal conductive electrode.⁴³⁴ Optimized performance of W_rec ∼ 3.8 J cm^–3 was also reported by Hao and co-workers for (Pb_0.88Ba_0.05La_0.02Dy_0.04)(Zr_0.68Sn_0.27Ti_0.05)O₃ MLs using 5%Ag/95%Pd internal electrode. However, this was accompanied by a large electric field-induced strain ∼0.71% at 300 kV cm^–1,⁴³⁵ a major drawback for practical applications due to inferior mechanical stability in operation. To date, Hao and co-workers have reported the record-high W_rec ∼ 12.6 J cm^–3 and η ∼ 80% under E_max ∼ 560 kV cm^–1 for Pb_0.98La_0.02(Zr_0.7Sn_0.3)_0.995O₃ ceramic ML, as shown in Figure 30.⁴³⁶

Figure 30.

(a) P–E loops under electric field up to E_max (b) calculated W_rec and η values for Pb_0.98La_0.02(Zr_0.7Sn_0.3)_0.995O₃ ceramic ML. Reproduced with permission from ref (436). Copyright 2020 Royal Society of Chemistry.

Although lead-free AFE ceramics show great promise, there are no reports of ceramic MLs in the literature. Instead, lead-free RFEs have dominated recent ceramic ML research due to their excellent W_rec and fatigue-resistant behavior accompanied by negligible electric field induced strain. 0.87BT–0.13 Bi(Zn_2/3(Nb_0.85Ta_0.15)_1/3)O₃ ceramic MLs have been fabricated with a dielectric layer t of 11 μm using 60Ag/40Pd internal electrodes that exhibit excellent W_rec ∼ 8.13 J cm^–3 and η ∼ 95% at 750 kV cm^–1.⁴³⁷ The energy storage properties of 0.87BT–0.13 Bi(Zn_2/3(Nb_0.85Ta_0.15)_1/3)O₃ were further enhanced by Zhao and co-authors by decreasing the dielectric layer to ∼5 μm, achieving W_rec ∼ 10.12 J cm^–3 at 1012 kV cm^–1, Figure 31a,b, as well as demonstrating good temperature stability from 75 to 175 °C.⁴³⁸ NBT–0.45SBT ceramic MLs have also been reported to exhibit W_rec ∼ 9.5 J cm^–3 with η ∼ 95% at 720 kV cm^–1,²²² which were further improved by forming a solid solution with a third perovskite end-member, 10%BMN to give W_rec ∼ 18 J cm^–3 with η > 90% at 1000 kV cm^–1, by Ji and co-workers, Figures 31c,d.³³⁷ The highest among all lead/lead-free ceramic MLs, W_rec ∼ 21.5 J cm^–3 at ∼1030 kV cm^–1, however, was achieved for textured NBT–0.3SBT ceramic MLs.⁴³⁹ Texturing was achieved through the use of <111>-oriented ST platelets which reduced field-induced strain at high field thus enhancing the BDS greatly. Combination of texturing with alloying with a third end member (BMN) in NBT–SBT may well represent an exciting path to achieve yet higher energy densities.

Figure 31.

(a) Bipolar P–E loops and (b) calculated W_rec and η of 0.87BT−0.13Bi[Zn_2/3(Nb_0.85Ta_0.15)_1/3]O₃ MLs under different electric fields.⁴³⁸ (c) Unipolar P–E loop, with inset SEM micrograph of ceramic MLs, and (d) calculated W_rec and η of NBT–0.45SBT–0.08BMN ceramic MLs under different E. (a, b) Reproduced with permission from ref (438). Copyright 2019 John Wiley and Sons; (c, d) Reproduced with permission from ref (337). Copyright 2021 Elsevier.

Wang, Reaney, and co-workers have employed a range of different chemical dopants and alloying additions to investigate BF-(B,S)T-based ceramic multilayers.^{34,45,90,276,416} In 2018, Nd-doped BF-0.3BT was reported to exhibit W_rec ∼ 6.74 J cm^–3 (more than 3 times higher than the bulk value) and η ∼ 77% at 540 kV cm^–1 with a layer t of 33 μm.⁹⁰ On the other hand, alloying BF–0.3BT with 8 mol % Nd(Zr_0.5Zn_0.5)O₃ resulted in W_rec ∼ 10.5 J cm^–3 with η of 87% at 700 kV cm^–1 with a dielectric layer t of ∼17 μm, Figure 17.³⁴ Further studies focused on promoting electrical homogeneity, which was considered to prevent conductive pathways developing in these composition, thereby avoiding the breakdown at high field and facilitating the improvement of W_rec by reducing the dielectric layer thickness. Finally, W_rec ∼ 13.8 J cm^–1 at 953 kV cm^–1 was achieved for BF–0.3BT–0.13 Bi(Li_0.5Nb_0.5)O₃ (BF–BT–0.13BLN) ceramic MLs at a dielectric layer t of 8 μm, as shown in Figure 32.²⁷⁶

Figure 32.

(a) Backscattering electron (BSE) cross section micrographs of BF–BT–0.13BLN ceramic MLs; (b) energy dispersive X-ray (EDX) mapping of all elemental distribution (c) transmission electron microscopy (TEM) micrograph obtained from an interface between a BF–BT–0.13BLN grain and a Pt grain (electrode); inset shows a high resolution TEM (HRTEM) image (filtered) obtained from the grain at a higher magnification. (d) Unipolar P–E loops and (e) calculated energy storage properties for BF–BT–0.13BLN ceramic MLs. Reproduced from ref (276). Copyright 2020 Royal Society of Chemistry.

3.2.2. Ceramic Films

Higher BDS and W_rec compared to MLs fabricated through powder-based technology have been reported for ceramic films deposited on LaNiO₃/Si₍₁₀₀₎ or Pt/Ti/SiO₂/Si substrates by physical vapor or chemical deposition techniques, such as radio frequency magnetron sputtering,⁴⁴⁹ spin coating,^101,450,451 pulsed laser deposition,⁴⁵² and chemical solution deposition.⁴⁵³ The BDS of ceramic films is significantly improved due to the reduction of t (<1 μm) often attributed to fewer defects (grain boundaries) and/or pore/void concentration. Not only have higher figures of merit been reported for ceramic films, but several researchers have proposed novel underlying mechanisms (beyond reduction in defect density) behind the enhancement, such as the formation of polymorphic nanodomains.⁴⁵² However, nonpowder based techniques are difficult to scale up into MLCCs, which may constrain these extraordinary results to lab-based fundamental research rather than promising practical output for commercial exploitation. The energy storage properties for ceramic films are summarized in Table 13.

Table 13. Energy Storage Properties for Different Ceramic Films.

materials	t (μm)	E (kV cm–1)	ΔP (μC cm–2)	Wrec (J cm–3)	η (%)	ref
Pb0.97La0.02(Zr0.97Ti0.03)O3	1.7	1158	∼90	20.1	64	(460)
1 mol %Fe–doped 0.72NBT–0.18KBT–0.10ST	0.9	1200	58	20.34	35.17	(461)
Pb0.97Y0.02[(Zr0.6Sn0.4)0.925Ti0.075]O3	0.5	1300	70	21	91.9	(462)
Pb0.92La0.08(Zr0.52Ti0.48)O3	0.4	1600	35	22	77	(463)
Pb0.85La0.1ZrO3	0.45	1500	45	23.1	73	(464)
0.942(Na0.535K0.480NbO3)−0.058LiNbO3	5	1400	∼52	23.4	70	(465)
(Sr0.85Bi0.1)Ti0.99Mn0.01O3	0.25	1982	35	24.4		(466)
Pb0.97La0.02(Zr0.98Ti0.02)O3	2	984	∼120	25.2	52	(467)
Bi(Mg0.5Ti0.5)O3	0.11	900	70	26	55	(468)
(Pb0.98La0.08)(Zr0.52Ti0.48)O3	0.25	2200	40	27.5	62.2	(469)
6 mol % BF-doped (K0.5Na0.5)(Mn0.005Nb0.995)O3	1	2000	∼40	28	90	(470)
PbZrO3/PbZr0.52Ti0.48O3	0.35	2615	60	28.2	50	(471)
(Pb0.92La0.08)(Zr0.65Ti0.35)O3	0.32	3000	40	29.7	50.8	(472)
1 mol % Mn-doped NBT	1.2	2310	∼40	30.2	48	(455)
(K0.5, Na0.5)(Mn0.005, Nb0.995)O3-6 mol % BF	1	1900	45	31	90.3	(470)
0.9Pb(Mg1/3Nb2/3)O3–0.1PbTiO3	0.375	2640	35	31.3	40	(473)
0.9(0.94NBT–0.06BT)–0.1NN	0.3	3170	20	32	90	(474)
0.4 Bi(Mg0.5Zr0.5)O3–0.6PbTiO3	0.5	2000	43	32.3	51.4	(475)
2 mol % Fe-doped (Na0.85K0.15)0.5Bi0.5TiO3	1.15	2300	∼35	33.3	51	(456)
0.95NBT–0.05ST	1.5	1950	∼55	36.1	41	(476)
Ba2Bi4Ti5O18	0.41	2150	38	37.1	91.5	(477)
Pb0.82La0.12Zr0.85Ti0.15O3	1	2100	∼52	38	71	(478)
5.8 mol % SiO2 doped HfO2	0.01	∼5800		40	72	(479)
Pb0.91La0.09 (Zr0.65Ti0.35)0.9775O3	1	1998	∼65	40.2	62	(449)
0.89NBT–0.06BT–0.05BF	0.28	1750	150	42.9	65.7	(480)
BaBi4Ti4O15	0.45	2000	40	43.3	87.1	(481)
Bi3.25La0.75Ti3O12	0.5	2040	50	44.7	78.4	(482)
Hf0.3Zr0.7O2	0.0092	4500	30	46	53	(458)
Pb0.97La0.02Zr0.66Sn0.23Ti0.11O3	0.65	4000	55	46.3	84	(453)
0.5 Bi(Ni1/2Ti1/2)O3-0.5PT	0.455	2250	62	46.7		(483)
Pb0.94La0.04(Zr0.98Ti0.02)O3	2	3699	∼48	47.4	25	(454)
0.4ST-0.6 Bi3.25La0.75Ti3O12	0.3	2470	38	47.7	87.4	(484)
0.9 Bi0.2Sr0.7TiO3-0.1BF	0.46	4800	25	48.5	47.57	(485)
Si doped Hf0.5Zr0.5O2	0.01	∼3500	∼30	50	80	(486)
0.6NBT–0.4 Bi(Ni0.5Zr0.5)O3	1	2200	∼63	50.1	64	(457)
0.5 Bi(Ni1/2Ti1/2)O3-0.5PT-excess 20% PbO	0.45	2250	75	50.2		(487)
(Pb0.97La0.02)(Zr0.7Sn0.25Ti0.05)O3	1.8	3750	∼75	56	45	(488)
0.6BT–0.4 Bi3.25La0.75Ti3O12	0.3	3200	45	61.1	84.2	(489)
6 mol % Si-doped HfO2	0.01	4500		61.2	65	(490)
8 atom % Al doped HfO2	0.05	4900	14	63	90	(459)
Na0.485Bi0.5(Ti0.96W0.01Ni0.03)O3	0.6	2500	30	63.1	55	(491)
BF/Bi3.25La0.75Ti3O12	0.14	2753	50	65.5	74.2	(492)
Bi0.525Na0.5(Ti.96W0.01Ni0.03)O3	0.4	2500	78	65.8	52.9	(493)
0.25BF–0.75ST	0.5	4460		70	70	(494)
0.97(0.93NBT–0.07BT)–0.03BF	0.35	2285	∼120	81.9	64.4	(495)
(Ba0.95Sr0.05)(Zr0.2Ti0.8)O3	0.1	6200		102	87	(496)
0.25BF–0.3BT–0.45ST	0.45	4900	90	112	80	(452)
0.68PMN–0.32PT	0.15	5800	∼100	130	75	(497)

Lead-based ceramic films have been studied heavily in the past decade using different preparation methods, particular for PLZT. (Pb_0.94La_0.04)(Zr_0.98Ti_0.02)O₃ ceramic film with t of 2 μm was reported to exhibit 47.4 J cm^–3 under electric field of 3700 kV cm^–1 on a Pt₍₁₁₁₎/TiO₂/SiO₃/Si₍₁₀₀₎ substrate.⁴⁵⁴ Further increasing La concentration resulted in superior W_rec ∼ 40.2 J cm^–3 and η ∼ 62%, achieved at E_max∼ 1998 kV cm^–1 for Pb_0.91La_0.09(Zr_0.65Ti_0.35)_0.9775O₃ on a LaNiO₃/F-Mica substrates.⁴⁴⁹ In addition, for Pb_0.97La_0.02Zr_0.66Sn_0.23Ti_0.11O₃, an improved W_rec ∼ 46.3 J cm^–3 was achieved at ∼4000 kV cm^–1, accompanied by excellent temperature stability (up to 380 K) and cyclic reliability (up to 10⁵), Figure 33.⁴⁵³

Figure 33.

(a) Unipolar P–E loops, inset image is cross-section SEM image of ceramic film deposited on the substrate, and (b) W_rec and η of Pb_0.97La_0.02Zr_0.66Sn_0.23Ti_0.11O₃ ceramic film under various E. (c) Temperature and (d) cyclic unipolar P–E loops, W_rec, and η of Pb_0.97La_0.02Zr_0.66Sn_0.23Ti_0.11O₃ ceramic film under E∼ 2200 kV cm^–1. Reproduced with permission from ref (453). Copyright 2018 Elsevier.

In the past few years, there has been increased focus on lead-free ceramic films due to concerns over toxicity of PbO. NBT and BF based relaxor compositions have dominated research and have competitive or even superior energy storage performance to lead-based films. Mn doped NBT ceramic films on a LaNiO₃/Si₍₁₀₀₎ substrates with t = 1.2 μm were reported to exhibit excellent W_rec ∼ 30.2 J cm^–3 under E_max ∼ 2310 kV cm^–1.⁴⁵⁵ Similar properties, W_rec ∼ 33.3 J cm^–3 under E_max ∼ 2300 kV cm^–1, were also obtained for Fe doped NBT–K_0.5Bi_0.5TiO₃ ceramic film.⁴⁵⁶ Recently, even higher W_rec ∼ 50.1 J cm^–3, η ∼ 63.9% accompanied by fast charge–discharge speed (∼210 ns) were achieved simultaneously at ∼2200 kV cm^–1 in relaxor 0.6NBT–0.4 Bi(Ni_0.5Zr_0.5)O₃ films.⁴⁵⁷ In addition, an ultrahigh W_rec ∼ 112 J cm^–3 with η ∼ 80% was reported by Pan and co-workers in BF–BT–ST ceramic films recently (Figure 34). Polymorphic nanodomains with competitive rhombohedral and tetragonal phases with competitive free energy were considered critical for the extraordinary electrical properties. BF was chosen as a main component due to its large spontaneous polarization. BT was introduced to form a solid solution to encourage coexistence of rhombohedral and tetragonal phases and finally ST was incorporated to further disrupt the long-range polar coupling and induce polymorphic nanodomains. By tuning the ratio of BF, BT and ST, a highly disordered composition is produced with rhombohedral and tetragonal nanodomain. The experimental observations were validated by phase-field simulations in the optimized composition, 0.20BF–0.25BT–0.55ST.⁴⁵² Apart from these perovskite lead-free relaxor candidates, HfO₂-based ceramic films have also been explored and demonstrate promising energy storage properties, stabilities/reliabilities, scalability, and integration. W_rec ∼ 46 J cm^–3 with excellent temperature stability (up to 175 °C) and cyclic fatigue resistant (up to 10⁹ time) was reported by Park in a 9.2 nm thick Hf_0.3Zr_0.7O₂ film⁴⁵⁸ and W_rec ∼ 63 J cm^–3 with η ∼ 85% were realized in 50 nm thick Al doped HfO₂ ceramic films with excellent temperature and frequency stability.⁴⁵⁹

Figure 34.

(a) Comparative display of Landau energy profiles and P–E loops of an FE with micrometer-size domains, an RFE with nanodomains, and an RFE with polymorphic nanodomains. Phase-field-simulated three-dimensional domain structures of (b) 0.45BF-0.55ST with rhombohedral nanodomains and (c) 0.20BF-0.25BT–0.55ST with coexisting rhombohedral and tetragonal nanodomains. (d) Simulation of the two-dimensional multiple nanodomain structure of 0.20BF-0.25BT–0.55ST in the cubic matrix. Contour plots of the simulated (e) W_rec and (f) η of BF–BT–ST solid solutions. Reproduced with permission from ref (452). Copyright 2019 The American Association for the Advancement of Science.

3.2.3. Summary of State-of-the-Art in Ceramic MLs and Films

The energy storage performances, W_rec, η, and ΔP, between bulk ceramics, ceramic MLs, and ceramic films are shown in Figure 35. The highest W_rec (∼130 J cm^–3) and E_max (∼5800 kV cm^–1) are obtained in ceramic films, followed by ceramic MLs (∼21 J cm^–3 and E_max ∼ 1000 kV cm^–1) and bulk ceramics (W_rec (∼12 J cm^–3 and E_max ∼ 650 kV cm^–1) and scale primarily with t of the dielectric layer (Figure 35d). W_rec in ceramic films is also improved by higher ΔP (up to ∼120 μC cm^–2) with respect to ceramic MLs (up to ∼70 μC cm^–2) and bulk ceramics (up to ∼60 μC cm^–2), as shown in Figure 35b. η for bulk ceramics, ceramic MLs and ceramic films varies significantly with composition and relates to factors such as, energy dissipation through a field induced transition to a long-range polar state, domain switching, polarization rotation, and leakage current relating to the presence of V_O^.. and associated defect dipoles.

Figure 35.

(a) E_max vs W_rec, (b) ΔP vs W_rec, (c) W_rec vs η, and (d) t vs W_rec and E_max between bulk ceramics, ceramic MLs and ceramic films.

From a commercial perspective, energy storage performance of lead-free ceramic MLs has improved significantly in the past few years with BF and NBT based ceramic MLs now rivalling lead-based ceramic MLs, delivering W_rec > 15 J cm^–3 at ∼1000 kV cm^–1. However, the selection of inner electrodes for these materials is limited to Pt, currently too costly for mass production. Commercial focus therefore, currently remains mainly on modified BT compositions which are compatible with Ni, Ag and Ag/Pd electrodes depending on composition and pO₂ during fabrication. A huge stride forward in the industry would be the development of a low cost equivalent to Ag/Pd electrodes which could permit the fabrication of a wider range of MLs at higher pO₂ which would inhibit the formation of V_O^.. and maintain high BDS.

4. Strategies for Improving Energy Storage Properties

The review of the state-of-the-art of ceramics, MLs, and films presented in section 3 clearly points to a set of criteria that are required to optimize energy storage performance. Though some of these have been alluded to in section 3 as part of the review of the state-of-the-art, it is worth collating these principals for RFE and AFE materials to act as guide for future materials development.

4.1. Optimization through an Induced Relaxor State

The most commonly utilized strategy to optimize energy storage properties is through inducing a relaxor state within a system that contains highly polarizable ionic species. It is typically carried out through strategic doping or alloying to form a pseudoternary solid solution. Levin and co-workers proposed that long-range ferroelectric correlations can be effectively “blocked” by using designed dopants with enhanced local polarizability.⁴⁹⁸ If this concept is married to dopant strategies to induce or maintain a homogeneous electrical microstructure, ΔP and BDS can be optimized leading to a large W_rec.

These two simple precepts can be applied to most systems; e.g., a frequency dispersion of ε_r was observed after doping xBi_2/3(Mg_1/3Nb_2/3)O₃ (B_2/3MN) into BT ceramics, along with reduction on maximum dielectric constant (ε_m) and associated temperature (T_m).²⁰⁹ After fitting the ε_r and T_m using modified Curie–Weiss law (as follows)

17

γ was found to be within the range of 1.61 for BT–0.06B_2/3MN ceramics with a Burns temperature ∼154 °C. In addition, XRD revealed a transformation from tetragonal to an average pseudocubic structure as increasing x concentration, coupled with a reduction in P_r, confirming a relaxor state at room temperature, Figure 36a,b.

Figure 36.

(a) XRD patterns with representative peaks and (b) Bipolar P–E loops for xB_2/3MN–BT ceramics with x = 0.00–0.10. (c) BSE surface micrographs of Ag–Pd cofired 0.06B_2/3MN–BT ceramics. (d) EDX mapping distribution of Ag, Pd, Ba, and Ti elements. Reproduced from ref (209). Copyright 2020 American Chemical Society.

An optimum W_rec ∼ 4.55 J cm^–3 at 520 kV cm^–1 was recorded for BT–0.06B_2/3MN (Figure 36)²⁰⁹ with similar properties reported for Bi(Mg_0.5Ti_0.5)O₃,²⁰⁸ Bi(Li_0.5Ta_0.5)O₃,²⁰³ Bi(Zn_0.5Zr_0.5)O₃,¹⁹⁹ and K_0.73Bi_0.09NbO₃⁴⁹⁹ doped BT ceramics. However, the advantage of BT–B_2/3MN is in the comparatively low concentration of dopant required to induce a relaxor state. The polar coupling is disrupted through a combination of A-site (Bi and V_O^..) and B-site (Nband Mg) dopants which provide a range of difference in ionic size, charge and electronegativity. However, only 4 mol % Bi is present on the A-site which minimizes reaction with Ag/Pd (Figure 36c). MLs with Ag–Pd internal electrodes may then be fired in ambient pO₂, minimizing V_O and reducing leakage current at high fields.

Similar compositional modifications to induce a relaxor state have been adopted in BF-based solid solution. For example, in FE BF–BT compositions, macroscopic herringbone-type domains are observed in (Figure 37a) but doping with 5 mol % Nd, induced a nanodomain state, accompanied by a decrease in P_r and a frequency-dependent permittivity curve, confirming the transformation from FE to RFE. Simultaneously, average G was reduced from ∼10 μm (BF–BT) to ∼2 μm (Nd doped BF–BT). As a result, enhanced W_rec ∼ 1.8 J cm^–3 was obtained for 15 mol % Nd doped BF–BT, which was further improved to 6.74 J cm^–3 by multilayering (Figure 37b,c)⁹⁰ with similar optimization reported in BF doped with Bi(Zn_2/3Nb_1/3)O₃ and Bi(Zr_0.5Zn_0.5)O₃.^34,266 The same design strategy of inducing an RFE state is also used in thin film BF–BT–ST compositions.⁴⁵²

Figure 37.

(a) TEM images of the domain structure in BF–BT, 5 mol % Nd–BT–BT and 10 mol % Nd–BF–BT. (b) The changes of W_rec as a function of electric field for x mol % Nd–BF–BT ceramics. (d) W, W_rec, and η for 15 mol % Nd–BF–BT MLs, with ceramic MLs microstructure as inset figure. Reproduced with permission from ref (90). Copyright 2018 Royal Society of Chemistry.

In summary, the inferior energy storage performance of all FEs can be improved by forcing a RFE state through strategic doping or alloying, Figure 38.^{34,66,90,92,199,204,272}P_max is often unsaturated in RFE and increases with E which means ΔP is not only a function of the slimness of the P–E loop but also of the applied E. RFEs are therefore, among the most promising candidates for capacitors in power electronics (E_max > 300 kV cm^–1).

Figure 38.

(a) E_max vs W_rec, (b) ΔP vs W_rec and (c) W_rec vs η for FEs and RFEs bulk ceramics to demonstrate the relaxor optimization.

4.2. Optimization of Antiferroelectrics

Many recent publications have focused on optimization of energy storage in lead-free AN-based materials,^347,353 although similar strategies date back to early studies of lead-based PLZT and PLZST AFE ceramics.^{118,126,143−145} The most comprehensive study of AN AFE ceramics was performed by Lu and co-workers in which they used A and B-site substitutions to develop, Ag_0.97Nd_0.01Nb_0.80Ta_0.20O₃ which yielded W_rec ∼ 6.5 J cm^–3 at 370 kV cm^–1 with η∼ 71%, Figure 39.³⁶⁶ In their study, Lu and co-workers defined several key points required to optimize AN-based ceramics:

• i)Optimization of P_max through local strain/field coupling around the smaller (with respect to Ag) Nd ion the A-site and its compensating V_A, Figure 39a.
• ii)Stabilization of the AFE structure through a combination of Nd and Ta doping which leads the induced AFE/FE transition to higher fields. Figure 39(a).
• iii)Inducing a slim hysteresis curve in the field induced region of the P–E loop. This was also achieved, through Nd doping that disrupted polar and antipolar coupling which manifested itself as a decrease in domain width from ∼1 to 0.5 μm and streaking of ±1/4(001)_c superstructure reflections in electron diffraction patterns for x = 0.03, Figure 40.

Figure 39.

(a) Schematic illustrating how W_rec is optimized through doping in AN. Gray, yellow, green, dark green, and red spheres represent Ag, Nd, Nb, Ta, and O atoms, respectively. The W_rec and η of (b) AN and (c) Nd_0.01Ta_0.20 codoped AN under the respective electric fields.³⁶⁶

Figure 40.

TEM [210]_c (c = cubic) zone axis diffraction patterns and corresponding dark field images obtained using (001) reflections from (a) and (b) AN and (c) and (d) Ag_0.91Nd_0.03NbO₃ (e) [210]_c zone axis diffraction pattern of Ag_0.97Nd_0.01Ta_0.20Nb_0.80O₃. (f) Bright field TEM image of domains in a grain of Ag_0.97Nd_0.01Ta_0.20Nb_0.80O₃.³⁶⁶

All the above maximize the area of the P–E loop to the left of the curve in the positive quadrant and thus optimize W_rec, Figure 39b,c.

Stabilization of the AFE structure was also confirmed by First-principles calculation and Ginzburg–Landau–Devonshire (GLD) phenomenology, as illustrated in Figure 41.

Figure 41.

Schematic contour diagrams of the free energy difference (ΔG) for (a) AN- and (b) Nd/Ta-codoped AN without electric field. Schematic contour diagrams of GLD phenomenological theory of AFE-to-FIE phase transition for (c) AN- and (d) Nd/Ta-codoped AN under application of electric field.³⁶⁶

Figure 42 summarizes the properties for many AFE systems. Most conventional AFEs exhibit low W_rec (∼1.5–2 J cm^–3) and η (∼40%), which can be optimized to W_rec > 3 J cm^–3 and η > 50% (enhanced-AFEs) by strategies described in the work of Lu and others.^{348,350−352,357,359,363,366} Similar values of ΔP (30–40 μC cm^–2) are found for AFEs and enhanced-AFEs which reflects an intrinsic limitation of AFE materials, i.e. when antiparallel polar coupling is fully switched to polar under electric field, P_max reaches saturation and is difficult to enhance unlike for RFEs (section 4.1). η is also difficult to further improve (>80%) due to the hysteresis above the AFE/FE switching field. However, at intermediate electric fields (∼300 kV cm^–1), much higher ΔP and W_rec are obtained for AFEs compared with RFEs, indicating that AFEs are more suitable for low/intermediate-voltage energy storage applications.

Figure 42.

(a) E_max vs W_rec and (b) ΔP vs W_rec and (c) W_rec vs η for AFEs and stabilized AFEs bulk ceramics to demonstrate the AFE stabilized optimization.

4.3. Other Strategies

4.3.1. Chemical Coating

Chemical coating is commonly reported as a strategy to optimize energy storage properties in BT ceramics. In general, chemical coatings are synthesized and applied using wet-chemical and sol–gel methods.^{69,74,79−81} Smaller grain sizes are typically obtained (on the order of a few nm), leading to the enhanced density and BDS. BT ceramics, for example, with a 4 nm thick SiO₂ coating, have a good value of W_rec (∼1.43 J cm^–3),²³³ in which SiO₂ coating inhibits grain growth, thereby modifying the microstructure and reducing DC current leakage. The effect of SiO₂ layer thickness on BT particles has been systematically investigated (Figure 43a-c). The highest W_rec(∼4.8 J cm^–3) and η (∼99.1%) were obtained for 20 wt % SiO₂ coated BT, as shown in Figure 43d–f.⁸⁸ Other coating materials such as Al₂O₃ and La₂O₃,^80,184 are also reported to optimize energy storage properties, as listed in Table 14.

Figure 43.

(a–c) TEM micrographs of BTnanoparticles coated with SiO₂: (a) BT@10 wt %SiO₂; (b) BT@15 wt %SiO₂; (c) BT@20 wt %SiO₂. The shell region is defined by red arrows and dash lines. Bipolar P–E loops of BT with (d) 10 wt % (e) 15 wt % (f) 20 wt % SiO₂ composite ceramics at the highest applied electric field, measured at 10 Hz and room temperature. Reproduced with permission from ref (88). Copyright 2019 Elsevier.

Table 14. Energy Storage Performance of Chemically Coated Materials.

materials	coating/thickness	Emax (kV cm–1)	Wrec (J cm–3)	η (%)	G (nm)	ref
BT	Al2O3/2 nm	108	0.51	80	∼160	(80)
BT	La2O3 SiO2/20 nm	136	0.54	85	∼250	(184)
BT	BiScO3/ 4 nm	120	0.68	81	∼100	(500)
BT	SiO2 and Al2O3/6 nm	190	0.725	80	∼80	(73)
BT	SiO2/10 nm	200	1.2	53.8	∼200	(501)
BT	SiO2/4 nm	290	1.43	53	∼120	(86)
Pb0.97La0.02(Zr0.33Sn0.55Ti0.12)O3	SiO2/2 nm	238	2.68	95	∼180	(87)
BT	SiO2/21 nm	370	4.799	95	∼220	(88)
Pb0.91La0.06(Zr0.552Sn0.368Ti0.08)O3	PbO-B2O3–SiO2–Al2O3–ZnO-MnO2/4 nm	402	7.4		∼250	(77)

4.3.2. Layered Structure

Layer-structures composed of multiple materials have been reported to optimize energy storage properties and are typically tape cast, followed by lamination. Both BDS and ε_r are optimized with the final properties related to the type of electroceramic material and the thickness of each layer. The BDS of BT-based ceramics was enhanced to >300 kV cm^–1 by laminating layers between BT–x wt % SiO₂ layers (t ∼ 20 μm) and BT layer (t ∼ 25 μm).⁴⁴¹ ε_r decreased but this was compensated by an increased BDS with increasing SiO₂ concentration in the BT–x wt % SiO₂ layers. ST + Li₂CO₃ (t ∼ 50 μm) and 0.93NBT–0.07Ba_0.94La_0.04Zr_0.02Ti_0.98O₃ (t ∼ 33 μm) layered structure were also fabricated via tape-casting with improved W_rec ∼ 2.72 J cm^–3 at 294 kV cm^–1.⁵⁰² Enhanced W_rec ∼ 2.41 J cm^–3 at 237 kV cm^–3 was also obtained for layer-structure ceramics with ST + Li₂O₃ (t ∼ 50 μm) and NBT–0.06BT (t ∼ 50 μm), Figure 44a.b.⁴⁴² The interface between the ST + Li₂O₃ and the NBT–0.06BT layer was further investigated using the finite element analysis. BDS was improved by reducing the breakdown paths between the ST + Li₂O₃ and the NBT–0.06BT layer, Figure 44c–e⁴⁴² with electrical field redistribution and interface blocking playing essential roles.⁵⁰³ BDS was also influenced by the difference in ε_r and thickness ratio between the adjacent layers.

Figure 44.

(a) Bipolar P–E loops of the ST+Li₂O₃, NBT–0.06BT and ST + Li₂O₃/NBT–0.06BT ceramic MLs. (b) Comparison of W_rec and electric field between the ST + Li₂O₃/NBT–0.06BT ceramic MLs and some recently reported lead-free ceramics. (c) Distribution of the electric field at 200 kV cm^–1, model of electrical tree propagation simulated using the finite element method for the ST + Li₂O₃/NBT–0.06BT MLs ceramic under (d) 200 kV cm^–1 and (e) 250 kV cm^–1. Reproduced with permission from ref (442). Copyright 2018 Royal Society of Chemistry.

5. Summary and Perspectives

5.1. Lead-Based Energy Storage Ceramics

Lead-based ceramics have great potential as energy storage materials in modern microelectronics where high voltage and temperature are required, such as in pulsed power and power electronic applications. Lead-based AFE-type ceramics exhibit extremely high energy density but optimizing BDS, η and minimizing electrostrain is problematic. Low BDS (<300 kV cm^–1) is often attributed to the volatilisation of lead/lead oxide which leads the formation of lead vacancy (V_pb^..) and V_O that results in current leakage. Such issues may be partially solved by a combination of improved processing and dopants but achieving the values of BDS observed in lead-free materials has proved elusive. The low η in lead-based AFE-type ceramics (<80%) is mainly a result of opening of the hysteresis loop at high field due to the stabilization of a field induced FE phase. This results in a change in crystal class from tetragonal (AFE_T) to rhombohedral (FE_R) giving large strains (>0.3%) which may prevent long-term cycling through mechanical fatigue.

The lack of popularity in researching lead-based compared with lead-free materials in the academic community has meant that exploration of novel systems is rather limited, but there are, for example, interesting mixed Pb- and Bi-based systems with high ε_r and a spontaneous polarization that would mirror some of the design principles adopted in lead-free ceramics, particularly in solid solutions which combine AFEs and relaxor end members. In addition, further work is required to understand crystal structure and phase transition behavior. Many systems have incommensurate modulations and their influence on AFE/FE switching needs to be explored further using in situ XRD and Raman (temperature/electric field), as well as utilizing advanced aberration corrected TEM to study the local structure.

5.2. Lead-Free Energy Storage Ceramics

Lead-free candidates, including BT, ST, BF, KNN, NBT, AN and NN-based systems, are extensively studied and summarized in this review. Research into lead-free materials far outweighs that in lead-based, due to how the potential environmental legislation surrounding manufacturing and the end use of lead-based products has influenced funding bodies and awards. As a result, the optimization of energy storage properties has progressed rapidly in the last 5 years. Successful strategies to improve properties include, disrupting long-range polar coupling particularly if the average ionic polarizability is increased or unaffected, construction relaxor feature (PNRs) in FEs and AFEs, enhancing E_g and as a consequence E_a, reducing the total electrical conductivity and promoting electrical homogeneity through the use of strategic dopants to modify defect chemistry. If these strategies are married with a reduction in the dielectric layer thickness, high values of W_rec ∼ 20 J cm^–3 and η ∼ 90% can be achieved. Recent work on texturing of ceramic MLs has also proved successful in enhancing W_rec but the complexity of this approach may inhibit commercial uptake. However, the two overriding issues with the majority of lead-free compositions, particularly those whose W_rec are >10 J cm^–3 are (i) the need to find an effective low cost internal electrode system that permits their commercial exploitation (currently almost all ML data is quoted with Pt internal electrodes) and (ii) pushing their operating window to >200 °C and >100 Hz. Interestingly, electrostrain, a major drawback in lead-based materials, is broadly speaking not an issue in most of the lead-free RFEs and AFEs since the measured values of strain are often significantly lower (<0.2%) even at high fields. As with lead containing ceramics, there are only a few comprehensive investigations of the energy storage mechanisms which require high field in situ studies to be performed. A greater understanding of the role of defect chemistry, doping and alloying is also required, particular on how this influences, E_g, resistivity and electrical homogeneity, and thus the W_rec and η, thermal stability and cyclic reliability. In addition, in commercial MLCCs, the ripple current, equivalent series R, failure mode, voltage rating, the reliability in high humidity need to be evaluated and explored.

5.3. Glass Ceramics

Glass-ceramics have the advantages of facile manufacture, high W_rec, ultrahigh η (low energy dissipation), ultrafast charge–discharge speed, excellent temperature/frequency stability. However, there are still challenges/problems. Generally, increasing the volume fraction of the crystal phase will increase the ε/P but decrease BDS. It is critical to balance the ε_r and BDS to obtain the highest W_rec. The mechanism of crystallization and control of crystal phase/microstructure is still ambiguous, which should be further investigated using, advanced TEM and in situ XRD/TEM as a function of applied field and temperature. Furthermore, although the theoretical W_rec (>15 J cm^–3, due to ultrahigh BDS, >1100 kV cm^–1) of glass-ceramics are much higher than other bulk ceramics and even MLs (15–20 J cm^–3), the measured/calculated W_rec by P–E loops and discharging processes is low (<2 J cm^–3) due to the lower applied electric fields. As a result, we recommend using the same test method (P–E loops and discharging process) to evaluate the practical energy storage performance for glass-ceramics, consistent with other dielectrics.

Biographies

Dr. Ge Wang is currently working as a postdoctoral research associate in the Functional Materials and Devices group at the University of Sheffield (UK). He obtained his Ph.D. degree from the Department of Materials at the University of Manchester in 2017, specifically on functional and structural behavior of lead-free ferroelectrics. His research focuses on ferroelectric, piezoelectric, energy density ceramic capacitors, lithium-ion batteries, high entropy oxide materials, and structural analysis using synchrotron X-ray diffraction.

Dr. Zhilun Lu obtained his Ph.D. from the Department of Materials Science and Engineering at the University of Sheffield, in January 2016 with the High-Quality PhD Thesis Prize. He worked as a research scientist in Helmholtz-Zentrum Berlin before joining The Henry Royce Institute and Department of Materials Science and Engineering at the University of Sheffield. His work focuses on functional and energy materials such as thermoelectrics, dielectrics, high permittivity oxides, microwaves, and geometrically frustrated magnets. He has expertise in ceramics synthesis and structural and property characterization. His paper “Superior energy density through tailored dopant strategies in multilayer ceramic capacitors” was selected as a HOT Article in Energy & Environmental Science.

Dr. Yong Li is an Associate Professor at School of Materials and Metallurgy, Inner Mongolia University of Science and Technology. He received his PhD degree in Materials Science and Engineering from Beijing Institute of Technology in 2016. His research focuses on ferroelectric and multiferroic materials for energy storage, photoelectric conversion, and microwave attenuation.

Dr. Linhao Li obtained his Ph.D. in 2017 from the Department of Materials Science and Engineering at the University of Sheffield under the supervision of Professor Derek C. Sinclair. After graduation, he continued his research in Sheffield as a PDRA. Linhao’s expertise is probing the composition–structure–electrical/thermal property relationships of functional materials, particularly using impedance spectroscopy. His current research focuses mainly on developing new materials for energy and electronic applications.

Dr. Hongfen Ji received her Ph.D. degree in 2013 in electronics science and technology from Xi’an Jiaotong University, Shaanxi, China. She has held visiting appointments at Department of Materials Science and Engineering, University of Sheffield, UK (2019–2020). Currently, she works as lecturer in the Laboratory of Thin Film Techniques and Optical Test, and School of Optoelectronic Engineering, Xi’an Technological University, China. Her main research interests are in lead-free piezoelectric and ferroelectric films and ceramics, and functional devices based on piezoelectric and ferroelectric materials.

Prof. Antonio Feteira is currently a Professor of Advanced Functional Materials at Sheffield Hallam University. He received his Ph.D. degree in Materials Science and Engineering from The University of Sheffield, where he holds a visiting position. In the past, he was a Senior Research Fellow at the University of Birmingham, having previously worked in the R&D department of TDK EPC, Austria. He held visiting research positions at Penn State University, Warwick University, Max Planck Institute, and Complutense University of Madrid, UFPeL. His research focuses on advanced ceramics for electronics, including actuators, energy storage devices, temperature sensors and wireless communication resonators. He is an Associate Editor for the Journal of the American Ceramic Society and the Journal of Electronic Materials and Editor for the International Journal of Applied Ceramic Technology.

Prof. Di Zhou is a professor at the School of Electronic and Information Engineering at Xi’an Jiaotong University (China). He received his Ph.D. degree in Electronic Science and Technology from Xi’an Jiaotong University in 2009 under the guidance of Prof. Dr Xi Yao and won the National Excellent Ph.D. Thesis Award nomination. He was a joint Ph.D. student at the Materials Research Institute. Pennsylvania State University from 2008 to 2009 and a research associate from 2015 tp 2018 at the Department of Materials Science and Engineering, University of Sheffield. He is an Associate Editor of the Journal of the American Ceramic Society and International Journal of Applied Ceramic Technology, Editorial Board Member of Materials Research Bulletin and Journal of Advanced Dielectrics, and Editorial Advisory Board member of ACS Applied Materials & Interfaces. His research interests include microwave dielectric materials, low temperature cofired ceramics technology (LTCC), energy storage capacitor materials, microwave absorption materials, and functional composite materials for high frequency communication/energy storage/microwave absorption.

Prof Dawei Wang is a Professor at Shenzhen Institute of Advanced Electronic Materials, Shenzhen Institute of Advanced Technology, Chinese Academy of Sciences. Previously, he was a Research Associate at the Department of Materials Science and Engineering of the University of Sheffield (2014–2020) and a visiting scholar at the Materials Research Institute of Pennsylvania State University (2016). He received his Ph.D. degree in Materials Processing Engineering from Beijing Institute of Technology in 2012. He is an associate editor for Journal of American Ceramic Society and Frontiers in Materials. His research focuses on the advanced electronic ceramics for energy storage/conversion/harvesting and translation of new materials to prototype devices/components for electronic systems.

Prof Shujun Zhang is a Professor at ISEM/AIIM of UOW, prior to which, he was a Professor at MatSE Department and Senior Scientist at MRI of Pennsylvania State University. He is the Associate EIC of IEEE Transaction UFFC and Associate Editor of Science Bulletin, Journal ACerS, Journal Electronic Materials, and section EIC of Crystals. He is a fellow of IEEE and ACerS, elected AdCom member of IEEE-UFFC. He is now focusing on fabrication microstructure–property-performance relationship of functional materials for piezoelectric sensors, transducers, and energy storage/harvesting applications.

Prof. Ian M. Reaney FIMM FRMS holds the Dyson Chair in Ceramics and leads the Functional Materials & Devices Group in the Department of Materials Science & Engineering, University of Sheffield. He is a Fellow of the Royal Microscopical Society and the Institute of Materials, Minerals and Mining (IOM3). He has won numerous awards including the Verulam Medal (2017) in Ceramics from IOM3 and was recently elected to the World Academy of Ceramics. He is an Adjunct Professor at Pennsylvania State University (PSU) and European site director of the Centre for Dielectrics and Piezoelectrics in partnership with PSU and North Carolina State University.

Author Contributions

^‡ G.W., Z.L., and Y. L. contributed equally to this work.

The authors declare no competing financial interest.