Recent advances in perovskite air electrode materials for protonic solid oxide electrochemical cells

Abstract

Intermediate-temperature proton-conducting solid oxide cells (P-SOCs) have emerged as a promising technology for power generation and hydrogen production. They have gained significant attention due to their lower operating temperature, higher efficiency, better safety and durability and simplified water management over conventional high-temperature oxygen-conducting solid oxide cells (O-SOCs). However, the performance of P-SOC air electrodes is hindered by the sluggish kinetics of oxygen reduction and evolution reactions, necessitating efficient conductivities of H⁺, O²⁻, and e⁻. Despite critical advancements, the search for optimal air electrode materials remains challenging. This review provides a comprehensive overview of recent advancements in perovskite materials for P-SOC air electrodes, covering fundamental mechanisms, material development, theoretical modeling, and practical applications. It highlights key progress in reaction kinetics, structure–property relationships, and modification strategies across widely studied perovskite-based systems. Particular emphasis is placed on understanding the correlation between structural characteristics and the electrochemical activity and stability of electrodes, which is essential for the rational design of high-performance, durable P-SOC materials. Additionally, advanced methodologies and mechanistic insights into newly developed air electrode materials are explored, with a focus on the role of theoretical simulations, including artificial intelligence (AI)-driven machine learning (ML) techniques. Finally, perspectives are provided on the future development of high-performance P-SOC air electrodes for industrial applications.

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Broader context

As the climate and environmental crisis deepens, the development of clean energy sources has become increasingly critical. Intermediate-temperature proton-conducting solid oxide cells (P-SOCs) show immense potential for power generation and hydrogen production due to their favorable thermodynamics, efficient kinetics, high energy efficiency, low carbon emissions, high tolerance to fuel impurities, modularity, and reversible operation capabilities. In contrast, commercially available oxygen-ion SOCs (O-SOCs) typically operate at temperatures exceeding 750 °C, where the extreme conditions accelerate material degradation and pose significant reliability challenges. To address these issues, researchers are increasingly focusing on P-SOCs, which operate at reduced temperatures of approximately 400–600 °C, enabled by their higher proton conductivity and lower activation energy. Despite these advances, several challenges persist. Even at 500 °C, the air electrode performance of P-SOCs lags that of O-SOCs due to substantial polarization resistance associated with redox reactions at the air electrode. Among potential air electrode materials, perovskite oxides emerge as promising candidates due to their unique electronic structure, superior catalytic properties for oxygen reactions, high thermal stability, and relative ease of synthesis. This review explores cutting-edge P-SOC technologies utilizing perovskite air electrode materials and provides a detailed outlook on future research directions. More broadly, this review provides a valuable reference for designing advanced air electrodes for next-generation P-SOCs.

1. Introduction

Amid the urgent push toward carbon neutrality and the global transition to sustainable development, electrochemical energy conversion and storage technologies have gained substantial attention.¹ Their unique capability to integrate renewable energy sources into existing infrastructures offers a powerful solution for reducing reliance on fossil fuels and supporting a resilient energy ecosystem.² Among these electrochemical technologies, solid oxide fuel cells (SOFCs) and solid oxide electrolysis cells (SOECs) hold particular promise due to their rapid reaction kinetics, high round-trip efficiency, fuel flexibility, and cost-effectiveness in terms of materials.³ However, the commercialization of their conventional versions, that rely on oxygen-ion conducting electrolytes, faces substantial challenges associated with high operation temperature, high operational costs, and low reliability detrimental to sustainability. Significant efforts have, consequently, been dedicated to enhancing their sustainability through improvement of cell durability and simplification of management system on oxygen-ion-based SOFCs/SOECs (O-SOFCs/SOECs) by reducing the operating temperatures to manageable levels (400–600 °C).^4,5 However, achieving these lower temperatures without compromising efficiency and performance remains challenging, as lower temperatures can hinder reaction kinetics and reduce ionic conductivity. Proton (H⁺) migration is believed to be much faster and has significantly lower activation energy (0.4–0.6 eV) than O²⁻ (>1 eV) as the smaller size and mass of the H⁺ allow it easily hop in the lattice.^6,7 Thus, protonic solid oxide electrochemical cells (P-SOCs) hold a greater promise for lower temperature operation.

P-SOCs is a sustainable hydrogen production and power generation technology that integrates economy, efficiency, and safety (Fig. 1(a)). Compared to low-temperature proton exchange membrane systems (PEM, Fig. 1(b1)) which typically operate at 50–80 °C, P-SOCs provide superior thermodynamic performance—including improved catalytic activity, reaction kinetics, and efficiency. Thermal energy compensation (TΔS) for total energy (ΔH) in P-SOC systems lead to lower electrical energy (ΔG) requirements compared to low-temperature catalysts, as shown by eqn (1).^8–10

ΔH = ΔG + TΔS (1)

Fig. 1 (a) Schematic of steam electrocatalysis and power generation in P-SOCs. (b) Principal schemes of (b1) PEMs, (b2) O-SOCs, and (b3) P-SOCs. (c) Comparison among low-temperature PEM, intermediate-temperature P-SOCs, and high-temperature O-SOCs.

Unlike PEM water electrolyzers, P-SOCs facilitate direct production of pure, dry hydrogen in electrolysis mode without a need for external gas purification and dehumidification, thereby making it substantially cost effective. Furthermore, the application of low-temperature PEM systems is constrained by the high cost of platinum group metal catalysts and the stringent raw material requirements, including specific levels of acidity, conductivity, and resistivity. Compared with O-SOFCs (Fig. 1(b2)), water generation at the P-SOFC (Fig. 1(b3)) prevents fuel dilution which contributes to enhanced efficiency and better fuel utilization. High temperatures, according to Stefan–Boltzmann law, increase thermal radiation losses and impose stricter requirements on the cells, interconnectors, and glass sealant to withstand high thermal stresses.¹¹ Consequently, it is economically imprudent to employ them amid their potential to decreases hydrogen production efficiency, shorten device lifespans, and accelerate degradation, in addition.^12,13 The “intermediate-temperature” P-SOCs with their fairly manageable temperatures are, therefore, better alternatives for efficient hydrogen generation via water electrolysis and power production through fuel cell mode (Fig. 1(c)).¹⁴

A typical P-SOC includes a dense electrolyte that enables proton conduction and is sandwiched between two porous electrodes: an air electrode and a fuel electrode. The electrolyte supports efficient proton transport, while the porous electrodes fulfil specific roles: the air electrode assists in oxygen reduction and evolution reactions (ORR/OER), and the fuel electrode is involved in the oxidation and reduction of fuel (H₂).¹⁵ It is imperative to develop solutions that are cost-effective, reliable, and long-lasting for P-SOC technology to be effectively applied, the solutions that include finding catalytic materials with high feed conversion and stability at typical elevated reaction temperatures. Consequently, advances through material design, e.g., modifying material structure/composition alongside SOC technology design are essential and key to the technology.¹⁶ The simultaneous involvement of O²⁻, H⁺, and e⁻ in reactions at the air electrode necessitates the corresponding conductivities for all these three charge carriers. However, materials with this unique triple-conductive capability is rare, highlighting their development an imperative.^17,18 Engineering of air electrodes for P-SOCs is highly complex as they play a crucial role in determining overall performance. This complexity has drawn significant attention to their development. The sluggish kinetics of the proton-involved ORR and OER at the air electrode within the range of 400 °C to 700 °C severely limit the electrochemical performance and energy efficiency.^5,19 This effectively necessitates development of more electrocatalytically active and stable materials, currently a major challenge, for sustainable future of P-SOCs.^20,21

Researchers have dedicated considerable effort to develop advanced air electrode materials for P-SOCs. Metal oxide electrodes (e.g., LiCoO₂) are cost-effective and chemically stable but suffer from low electronic conductivity and phase instability in humid or CO₂-rich environments, limiting their long-term performance. Precious metal electrodes (e.g., Pt, Ru) offer exceptional catalytic activity and high electronic conductivity, ensuring fast reaction kinetics, but their high cost, limited availability, and susceptibility to sintering make them less viable for large-scale applications.^5,8 In contrast, perovskite oxides are the dominant choice for air electrodes, owing to their distinctive electronic structure, catalytic properties in oxygen-related reactions, high thermal stability, and ease of raw material synthesis. Over time, studies on perovskite air electrodes have covered various categories, including oxide materials exhibiting dual O²⁻/e⁻ conductivity, composites that combine proton conductors with dual O²⁻/e⁻ conductors, and newly developed single-phase materials that achieve triple H⁺/O²⁻/e⁻ conductivity. Herein, the recent progress and challenges associated with perovskite air electrode materials for P-SOCs are systematically reviewed. This review aims to provide an in-depth examination of the recent advancements and ongoing challenges in perovskite air electrode materials for P-SOCs. The state-of-the-art research is categorized into several representative air electrode material types—(Ba/Sr/Pr)MO_3−δ-type, LnBaM₂O_5+δ-type, and Ln₂NiO_4+δ-type—with modification methods spanning bulk structures, interfaces, and composites (Fig. 2). Moreover, the application of advanced computational techniques, particularly artificial intelligence (AI)-driven machine learning (ML), is explored for materials design and theoretical explanation. Commercial considerations and future research directions for P-SOCs air electrodes are also discussed, providing a comprehensive perspective on advancing P-SOCs technology.

Fig. 2 Typical modifications of perovskite air electrode materials.

2. Perovskite air electrode materials for P-SOCs

P-SOCs represent a significant advancement in the fields of hydrogen production and energy conversion, operating efficiently in both fuel cell (FC) and electrolysis cell (EC) modes. Their operational versatility makes them crucial technologies in sustainable energy systems, leading to great interest in recent years. This section discusses the electrochemical processes involved in hydrogen generation and power generation and then analyzes the role and requirements of perovskite air electrode materials and the corresponding evaluation methods.

2.1. Fundamentals of protonic solid oxide electrochemical cells

A typical P-SOC single cell features a sandwich-like structure with the configuration of air electrode/proton-conducting electrolyte/fuel electrode (Fig. 3(a)). The electrolyte functions as a dense electron-blocking layer, preventing direct combustion between the fuel (H₂) and air (primarily O₂) while facilitating rapid proton transport.^22,23 The fuel electrode materials comprising of metal (commonly Ni) and electrolyte component is where H₂ generated or consumed.²⁴ The air electrode serves as a site for ORR or OER.²⁵ During typical FC operation, electricity is generated through the electrochemical conversion of hydrogen (supplied as fuel on the fuel electrode) which serves as the proton source. Oxygen (from the air) in the air electrode reacts with the proton in the present of electrons (ORR) to form water (eqn (2) – FC mode). In EC mode, steam undergoes electrolysis to produce H₂ by flowing it through the air electrode for OER (eqn (2) – EC mode).²⁶ When voltage is applied, water decomposes into H⁺ and O₂, releasing the electrons. H⁺ then migrates through the electrolyte to generate dry H₂ which is then swept off the reaction chamber using inert gas, e.g., N₂ or Ar, while electrons flow through the external circuit to the fuel electrode.^27,28Eqn (2) and (3) show reactions at both electrodes and Fig. 3(a) exhibits their simple schematic versions.

Fig. 3 Fundamentals of perovskite air electrode materials in P-SOCs. (a) Schematic of typical P-SOCs and reactions at the air electrode: ORR reaction and OER reaction (as exemplified by triple-conducting material). (b) Structure of perovskite: (b1) simple perovskite, (b2) and (b3) double perovskite, and (b4) Ruddlesden–Popper phase. (c) Typical elements in perovskite.

Air electrode:

Fuel electrode:

In recent years, new reversible P-SOCs have emerged that can convert fuel into electricity in FC mode and water into hydrogen in EC mode.²⁹ Unlike prevalent batteries, which face limitations such as self-discharge and high storage costs, the reversibility of reversible P-SOCs potentially enables long-term energy storage by coping with seasonal energy storage adjustments.³⁰ It operates electrolytically when additional baseload power is needed to supplement solar and wind.³¹ Since chemical fuels can be stored indefinitely or used immediately as renewable feedstocks in various industrial applications, converting renewable electricity into chemical fuels, aiding the transition to a fully sustainable energy economy.⁹ The operational reversibility enables efficient energy storage during periods of renewable electricity oversupply and supports smooth electricity generation, helping to balance supply and demand while expanding the potential for large-scale SOC applications.

2.2. Key properties of perovskites for air electrodes

Essentially, the electrocatalytic process at air electrode of P-SOCs encompasses both the OER and ORR along with hydration reaction, involving the transport of O²⁻, H⁺, e⁻. By virtue of multi-conducting materials, the active sites of the proton-involved ORR and OER can cover the entire electrode surface instead of being confined to the electrode/electrolyte interface. Perovskites and perovskite-like structures dominate multiphase conducting materials, sensibly making them widely utilized as air electrodes.^32–34

2.2.1. Typical structure of perovskites. Typical perovskite has ABO₃, A₂B₂O₅, or A₂BO₄ structure (Fig. 3(b)), wherein A is a rare earth metal or alkaline-earth metal, e.g., lanthanide with +2 or +3 valence, and B is a catalytically active transition metal (TM) (Fig. 3(c)). In recent years, elements like Ba, Sr, La, and Pr at the A-site, along with Fe, Co, and Ni at the B-site, have become popular choices in perovskite structures. A-Site ions coordinate with B-site ions in a way that stabilizes the framework through a 12-fold cuboctahedral coordination with oxygen ions, while B-site ions are arranged in a 6-fold coordination with surrounding oxygen anion octahedra.³⁵ Although B-site ions primarily govern ionic and electronic conduction, A-site ions are crucial for maintaining structural integrity and can influence ionic conduction through interacting with the BO₆ octahedra and O²⁻.^36,37 Over the past few decades, three primary types of perovskite and perovskite-related oxides with H⁺ conduction capabilities have emerged and are extensively utilized as air electrodes in P-SOCs: simple perovskites, double perovskite, and Ruddlesden–Popper (RP) phases.^15,20,36,38

Simple perovskite (ABO_3−δ, Fig. 3(b1)), such as BaCoO_3−δ (BCO), BaFeO_3−δ (BFO), and SrCoO_3−δ (SCO),^39–41 feature rare earth metal or alkaline-earth metal (A site) typically coordinating with oxygen in a 12-fold cuboctahedral arrangement within BO₆ octahedra.⁴² TMs (B site) are generally found in sixfold coordination within these octahedra. Double perovskite (A′AB₂O_5+δ and A₂BB′O_5+δ, Fig. 3(b2) and (b3)), e.g., PrBaCo₂O_5+δ (PBC), PrBa_0.8Sr_0.2Co₂O_5+δ (PBSC), and PrBa_0.8Ca_0.2Co_1.95Pd_0.05O_5+δ^9,43 consist of alternatively stacked layers in the form of AO|BO₂|A′O|BO₂|. A is typically a lanthanide ion (+3), e.g., Pr, Nd, Sm, or Gd or alkaline-earth ion (Ba or Sr), and B is often a first-row TM ion or a mixture thereof. The RP phase has general formula: A_n+1B_nO_3n+1, where n denotes the number of perovskite-like octahedral layers (Fig. 3(b4)), exemplary materials being La₂NiO₄, La₃Co₂O₇, and La₄Ni₃O₁₀,^44–46 They exhibit a unique layered structure in which a sequence of (ABO₃)_n perovskite layers is interleaved with rock-salt-type AO layers. A₂BO_4+δ represents the simplest RP structure. The perovskite structure is flexible, adaptable, and tunable. By modifying their properties, e.g., type, composition, phase structures, and synthesis methods, electronic conductivity, ionic conductivity, pore morphology etc., can be altered and subsequently, catalytic performance.^36,37

2.2.2. H⁺, O²⁻, and e⁻ conduction mechanism. To identify and develop highly efficient perovskite air electrode materials, it is essential to understand the fundamental principles governing the generation and transport of charge carriers (H⁺, O²⁻, and e⁻). The e⁻ conduction in perovskites is characterized by a polaron hopping mechanism. Electronic charge carriers are transported through the bond network form between B-site cations and oxygen ligands, involving aliovalent TM sites.^47–49 This process, known as small-polaron hopping mechanism (Fig. 4(a)),^50–52 requires that the concentration of multivalent TMs exceeds the bond percolation threshold—a critical region (15–29%) where conductivity markedly increases over a specific range of compositions.^53,54 Below this threshold, dispersed TM sites act as low-energy traps, resulting in a sharp decline in electronic conductivity while they establish low-energy conduction pathways above the percolation threshold.^50,55 The composition at which this critical region emerges can be reduced through ordering, aggregation, or preferential site occupancy.^54,56

Fig. 4 Charge conduction mechanism (red spheres: O²⁻; yellow spheres: e⁻; green spheres: H⁺). (a) e⁻ conduction mechanism: small-polaron hopping mechanism. O²⁻ conduction mechanisms: (b) vacancy diffusion mechanism and (c) oxygen interstitial diffusion mechanism. H⁺ uptake mechanism: (d) hydration mechanism, (e) hydrogeneration mechanism, and (f) new-type hydration mechanism; H⁺ transfer mechanism: (g) vehicle mechanism, (h) Grotthuss mechanism, and (i) H⁺ reorientation and hopping for Grotthuss mechanism.

In perovskite air electrodes for oxidizing atmospheres, O²⁻ diffusion is the primary mode of mass transport. This process is facilitated by O²⁻ vacancies formed through acceptor substitution which create channels for transport. As a result, the mechanism is conveniently referred to as vacancy diffusion (Fig. 4(b)).^57,58 The symbol, “˙˙”, represents a net double negative charge while “×” represents a neutral charge. On the surface, these capture adsorbed oxygen (O₂) into lattice oxygen () and release holes (h˙), which then combine with electrons from the fuel electrode to complete the electrical circuit (eqn (4)). In single-structure perovskites, oxygen ion transport via the vacancy diffusion mechanism whereas some layered and spinel perovskites utilize the oxygen interstitial diffusion mechanism.^46,59–61 Oxygen ions, in the latter process, migrate through the interstitial spaces between lattice sites (Fig. 4(c)).^55,62

Since protons are not native to oxide lattices, proton uptake and conduction in perovskites heavily relies on hydroxide defects.⁶³ Three mechanisms have been proposed for proton uptake: hydration mechanism, the hydrogeneration mechanism, and a novel hydration mechanism recently introduced in the literature.^20,64,65 The Kröger–Vink notation describes the primary hydration mechanism, where H₂O decompose into hydroxide ions (OH˙) and H⁺, with OH˙ filling and H⁺ bonding with (eqn (5) and Fig. 4(d)). It is predominant when hole concentration is low.^66–68 Hydrogenation mechanism, occurs when oxides have abundant h˙, involving a redox reaction that incorporates H⁺ and releases O₂, bypassing the need for (eqn (6) and Fig. 4(e)).^69–71 Recently proposed hydration mechanism based on La_0.7Sr_0.3MnO_2.95 (LSM) suggests that water absorption and desorption are linked to the oxidation of manganese cations and the occupation of oxygen holes by H⁺ (eqn (7) and Fig. 4(f)).⁷² It further suggests that oxides with significant oxygen vacancy carriers and mixed valence states, like LSM, are effective for proton conduction, making them suitable for applications requiring triple-conducting materials.⁷² This theory has since been applied to a wider range of perovskite materials, including those containing cobalt cations, iron oxide ions, and others. Since water adsorption is an exothermic reaction,^73,74 proton conduction becomes the dominant mechanism at lower temperatures, leading to a reduced operating temperature for P-SOCs compared to O-SOCs.

Proton transfer in perovskite can occur through the vehicle and Grotthuss mechanisms, both identified by Kreuer et al.^68,75,76 In the vehicle mechanism (Fig. 4(g)), protons move with the O²⁻, a process constrained by the diffusion rate of O²⁻, leading to higher activation energy because of the limited mobility of O²⁻ ions.^77,78 Conversely, the Grotthuss mechanism (Fig. 4(h)), which is more commonly accepted, exhibits lower activation energy (<0.4 eV) and involves protons hopping between O²⁻ ions through the alternation of OH˙ bonds. This process consists of two steps: (1) reorientation of the OH˙ group to reduce the energy barrier for H⁺ transfer, and (2) H⁺ hopping between O²⁻ ions, aided by protonic defects (Fig. 4(i)).^77,79 In P-SOCs air electrodes, the reaction kinetics are dictated by the balance between electronic and ionic conductivity. High electronic conductivity facilitates the rapid movement of electrons, ensuring that redox reactions occur efficiently at the electrode–electrolyte interface. However, excessive electronic conductivity can reduce the concentration of mobile ionic species, limiting the transport of O²⁻ or H⁺. Conversely, high ionic conductivity is essential for sustaining continuous ion diffusion within the electrode bulk, but it may come at the cost of reduced electronic charge transfer, leading to increased polarization resistance and sluggish electrode kinetics.^80,81 However, optimizing one property often compromises the other. For example, accelerate ionic transport but can deteriorate the electronic conductivity and weaken structural stability. Therefore, air electrode design is not solely about maximizing one type of conductivity but rather achieving an optimal balance between ionic and electronic transport properties to enhance overall performance.

2.3. Performance-influencing factors

Effective operation of air electrodes relies on several complex sub-steps including the adsorption/desorption of reactants, generation of active ions, and ion transport within the electrode. High-performance air electrodes should possess both superior ion transfer capabilities and surface catalytic activity to improve reaction kinetics.⁸² Key factors influencing performance include oxygen vacancies,⁸³ hydration abilities,⁸⁴ oxygen/proton diffusion capabilities,⁸⁵ and surface oxygen ORR and OER activities.^86–89 Optimal performance in P-SOCs requires a balance among these various factors which this section looks into, based on recent advances.

2.3.1. Triple-phase conductivity and catalytic activity. Various perovskite-based air electrodes have been reported for P-SOCs, including conventional electronic conductors,^90,91 mixed ionic–electronic conduction,^92–95 and H⁺/O²⁻/e⁻ composite conductors.^22,96–98 Despite these advancements, the unsatisfactory ORR and OER keep the technology unsustainable for commercial implementation. Given that ORR and OER in P-SOCs air electrodes primarily involve oxygen/water adsorption/desorption on the electrode surface, reduction of active oxygen/oxidization of water to corresponding active ions (O²⁻, H⁺/hydroxyl), surface and bulk transport of active ions, and generation of water/oxygen with the participation of electrons.⁹⁹ Highly efficient triple conduction (O²⁻, H⁺, e⁻) is required. Recent advances have exhibited triple-conducting oxides with abundant surface active sites as promising materials for the role, achieving reasonable to high conversions while maintaining structural integrity at elevated temperatures.^100,101 This is partially thanks to significantly enhanced surface and bulk ion generation and transport capabilities of triple-conducting oxides, extending the effective three-phase boundaries and accelerating the electrode reaction kinetics.¹⁰² OER and ORR can occur over the entire electrode surface (double-phase boundary) rather than being confined to the electrode–electrolyte–gas triple-phase boundary.

Catalytic activity of perovskite influences generation of active ions thus influencing the concentration of charges in system.^63,103–105 The electronic structure is considered crucial for catalytic activity involve the occupancy of the e_g orbitals in surface TMs and the energy levels of the O p-band center.^106,107 To enable rapid surface oxygen-exchange kinetics and unimpeded transport of dissociated oxygen species, to controlling the redox activity of lattice oxygen in perovskites is crucial.¹⁰⁸ This control can generally be achieved through several strategies: substituting A-site with lower-valence cations, inducing oxygen non-stoichiometry, and replacing B-site cations with more electronegative elements.^109–113 The concentration of affects the adsorption process of reactants in the proton-involved ORR and OER, thereby affecting the efficiency of the catalytic reaction. For instance, concentration connected to B-site cations influences proton uptake in oxides that rely on a hydration mechanism. The micro/nano-structure of perovskite, including its surface area and exposed lattice planes, directly influences surface reaction activity.¹⁰⁵ Lattice planes exhibit higher activity like (110) for PrO_1.8, high surface area, and large pores, for example, improve the activity through increasing active sites dispersion, reducing the distance for bulk ionic diffusion, etc.^114–116 To advance the industrialization of P-SOCs, it is, therefore, imperative to develop electrodes with excellent triple-phase conductivity and highly electrocatalytic activity for both ORR and OER.^117,118

2.3.2. Structural symmetry, cations ordering, and defects. High symmetry in crystal structures, such as in cubic perovskites, enhances the mobility of charge carriers (H⁺, e⁻, and O²⁻) by facilitating the rotation of corner-sharing oxygen octahedra, minimizing lattice distortion needed for vacancy diffusion, promoting isotropic electron transport and greater atomic orbital overlap, which reduces small-polaron localization and binding energy.^{4,32,52,119–121} However, deviations from cubic symmetry, such as transitioning from a cubic phase to orthorhombic, tetragonal, or other non-cubic structures caused by multiple cations at A- and B-sites, can increase O–O bond distances, thereby increasing the hopping distance and reducing mobility of protons. Furthermore, distortions in phase structure can create multiple inequivalent oxygen sites. These variations between adjacent sites may form isolated low-energy sites which can act as traps for protons and , thereby impeding proton conduction. Additionally, hydration enthalpies differ based on bond distances and local environments, influencing H⁺ absorption and conduction. Lower hydration enthalpies generally favor enhanced H⁺ conduction contrary to high enthalpies.

In addition to charge-carrier mobility which is influenced by structure of perovskites, structural stability including chemical and thermodynamic stability is significate parameters. Based on the findings of Anthony F. et al., the tolerance factor (t), calculated from the ionic radii of A-site (r_A), B-site (r_B), and O site ions (r_O) (eqn (8)), accurately estimates stability of perovskite structures.^122,123

Perovskite oxides exhibit structural stability within the range of 0.75 < t < 1 with deviations causing varying degrees of lattice distortion.^124,125 Besides radii, the ordering of cations is also important. A-site ordering improves oxygen transport and hydration, as evidenced by materials like GdBaMn₂O_6−δ and LaBaCo₂O_6−δ,^126–128 while B-site ordering and oxygen-vacancy ordering are detrimental to proton conductivity and hydration.^91,129,130

Structural defects such as vacancies, dislocations, stacking faults, and grain boundaries play a crucial role in determining the electrochemical performance of perovskite air electrodes. , h˙, and protonic defects are particularly important in enhancing ionic and electronic conductivity, as they facilitate O²⁻/H⁺ diffusion and e⁻ transfer through the electrode bulk. Dislocations, which occur due to lattice mismatches or strain during synthesis and operation, can alter charge carrier mobility and introduce localized stress points that may accelerate material degradation.¹³¹ As an effective strain carrier, stacking faults can directly introduce strain fields and significantly affect the centre position of the d-band, thereby changing the interaction of surface adsorbates. Grain boundaries serve as pathways for ion transport but can also act as barriers that increase resistance, depending on their structure and chemical composition. However, excessive defects formation can lead to crystal lattice distortion and phase instability, negatively impacting electrode durability. Therefore, an ideal perovskite structure features symmetry cubic with a moderate defect concentration, and ordered A-site occupancy, while maintaining disorder in the B-site and anion sublattices to maximize overall conductivity.¹³²

2.3.3. Electronegativity. The basicity of O²⁻ is influenced by the electronegativity of cations, which in turn affects proton uptake and transport in oxides.²⁰ Specifically, a reduction in cation electronegativity tends to enhance the basicity of O²⁻, leading to increased proton absorption and the equilibrium constant of the hydration reaction.⁶⁶ For instance, substituting 50% of Sr (electronegativity = 0.95) in Sr_0.5Ba_0.5FeO_2.5 with Ba (electronegativity = 0.85) notably improved the proton concentration (C_H value) to 0.019 mol unit⁻¹, significantly higher than 0.0002 mol unit⁻¹ in SrFeO_2.5.¹³³ Similarly, the C_H value of Ba_0.75La_0.25FeO_2.53 increased from 0.004 to 0.03 mol unit⁻¹ when La (electronegativity = 1.1) was replaced with Ba. In La_1−xSr_xCoO₃, as the value of x increased from 0.2 to 0.7, the C_H value rose from 0.005 to 0.028 mol unit⁻¹.⁹¹ These changes in hydration capacity are attributed to Ba's lower electronegativity (0.85) compared to Sr (0.95) and La (1.1), which reduces the basicity of O²⁻ and enhances its hydration ability. Additionally, the electronegativity of elements at B-site elements (such as Mn (1.55), Fe (1.83), Co (1.88), Ni (1.91), Zn (1.65), Zr (1.33), and Ce (1.12)) plays a crucial role in improving proton uptake.^134,135 For example, inclusion of Zn at B-site increases the C_H value of Ba_0.95La_0.05Fe_0.8Zn_0.2O_2.4 to 0.1 mol unit⁻¹, compared to just 0.03 mol unit⁻¹ without Zn.⁹¹ Similarity, the C_H value of Ba_0.5Sr_0.5Fe_0.8Zn_0.2O_2.4 which features B-site cations with lower electronegativity, is approximately ten times greater than that of Ba_0.5Sr_0.5Co_0.8Fe_0.2O_2.4.

2.3.4. Local environment. Carrier concentrations in electrochemical materials are influenced by the local environment, such as composition, temperature, and atmospheric conditions. At high temperatures, lattice dehydration decreases proton concentration, shifting conductivity towards hole or oxygen-ion conduction.^136,137 Oxidizing environments provide oxygen filling the vacancies, reducing both O²⁻ transport and H⁺ uptake while increasing hole concentration.¹³⁸ Conductivity can drop significantly under oxidizing conditions and lower temperatures, especially if p-type small polarons are involved, as they have low mobility in the lower temperature. For example, protons dominate as charge carriers in Ba_0.85K_0.15ZrO_3−δ in pure oxygen up to 600 °C, but at higher temperatures, holes become predominant.¹³⁹ The gas environment influences the adsorption and conduction processes in both ORR and OER, and it may also impact the material's durability. The presence of water can initially boost ORR/OER activity but may eventually inhibit oxygen adsorption, leading to degradation and material changes, particularly at lower temperatures where this issue is exacerbated.^140–142 Protons are the main carriers in hydrogen environment, as seen in BaCe_0.9−xZr_xY_0.1O_3−δ, whereas under oxidizing environment, both protons and oxygen ions are generated and the materials exhibits p-type triple conduction.¹⁴³ When exposed to reducing atmosphere, perovskites can experience oxygen release from the lattice, leading to an increase in lattice volume.¹⁴⁴ The reactivity of alkaline earth metals in environments containing water vapor must also be taken to consideration amid their high likelihood to react. Closely managing material properties relative to reaction conditions and environment, in the nutshell, is essential for a successful reaction.¹⁴⁵

2.4. Evaluation methods

The development of new technologies, such as P-SOCs, typically follows a conventional progression from lab-scale research to pilot plant trials and, eventually to industrial-scale production.¹⁶ The initial lab-scale phase is crucial, as the reliability and accuracy of the experimental data gathered at this stage heavily influence the subsequent stages of technology development. For successful scaling-up and industrialization of sustainable P-SOCs technology, it is essential to have all necessary information well documented and understood. This section gives an overview of electrochemical methods and characterizations based on recent developments, and the recommended research path for air electrodes is summarized in Fig. 5.

Fig. 5 Recommended research paths for air electrode materials.

2.4.1. Electrochemical evaluation. Catalytic activity of air electrodes in the laboratory is evaluated by determining the area-specific resistances (ASRs) from the electrochemical impedance spectra (EIS) of an air electrode|proton-conducting electrolyte|air electrode symmetrical cell configuration (Fig. 6(a)). In the Nyquist plot, the impedance spectra display distinct, separable depressed arcs associated with different electrode processes.^146,147 The Ohmic resistance (R_ohm) comprises the electrode Ohmic resistance, electrolyte resistance, and lead resistance, and R_ohm is usually ignored to focus on the performance of electrodes. The difference between intercepts on the real axis of the impedance plot is typically interpreted as electrode polarization resistance (R_p). ASR is calculated by multiplying R_p by the tested areas (A) according to eqn (9).^148,149

ASR = (A × R_p)/2 (9)

Fig. 6 Recent advances in evaluation methods of air electrode materials. (a) EIS curves, (b) DRT, and (c) k_H and D_H of PBSCF and W infiltrated PBSCF. Reproduced with permission.¹⁵⁰ Copyright 2024, Wiley-VCH. (d) I–V curves of FC and EC modes. Reproduced with permission.¹⁵¹ Copyright 2024, Wiley-VCH. (e) Scheme for the H₂ permeation test with the proton-conducting membrane. Reproduced with permission.⁸⁵ Copyright 2023, Wiley-VCH. (f) Schematic of ToF-SIMS measurement. Reproduced with permission.¹⁵² Copyright 2021, Wiley-VCH. (g) XPS spectra of O 1s. Reproduced with permission.¹⁵⁰ Copyright 2024, Wiley-VCH.

Distribution of relaxation time (DRT) analysis can be performed based on impedance spectra by deconvoluting the EIS response to distinguish overlapping electrode processes and activation energy (E_a) is then determined from the Arrhenius plots (Fig. 6(b)).^153,154 Catalytic activity test is typically performed in dry air or humidified air. Typically, electronic conductivity dominates because the mobility of e⁻ is several orders of magnitude greater than that of H⁺ and O²⁻.^155,156 The conductive mechanism of the electrode material can be inferred from its conductivity behavior (Fig. 6(c)). For instance, the conductivity of SSNCF material initially rises and then declines with increasing temperature, peaking at 375 °C, which suggests a transition in its conductive mechanism from p-type to n-type around this temperature. Electrical conductivity relaxation (ECR) measurements, utilizing a 4-probe DC conductivity method, are employed to investigate the chemical diffusivities and surface exchange of H⁺ and O²⁻, as well as electrical conductivity.^83,157 Chemical diffusion coefficient (D_chem) and surface exchange coefficient (k_chem) are derived by fitting the transient conductivity data to the solution of Fick's second law using a non-linear least-squares fitting approach, based on the Levenberg–Marquardt algorithm.¹⁵⁷D_chem is calculated from the transient conductivity behavior in response to changes in partial pressure. However, for samples with low proton concentrations, the accuracy of ECR measurements may be limited.³¹ To isolate the conductivity attributed solely to surface protons, a common approach involves calculating the difference between conductivities measured in wet and dry atmospheres.¹⁵⁸ It is crucial to ensure that the volume conductivity of the target materials remains unaffected by atmospheric conditions.¹⁵⁹

Electrochemical performance of air electrode materials in SOCs is typically accessed using single cell configuration featuring air electrode|BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ (BZCYYb)|Ni-BZCYYb. In P-SOFC, H₂ serves as the fuel on the anode side, while air acts as the oxidant on the cathode side. P-SOEC utilizes steam as the feedstock for H₂ production, on the contrary. The evaluation of electrochemical properties involves measuring polarization resistance (R_p) from EIS test, determining peak power densities (PPDs) and current densities from current density–voltage (I–V) curve, and assessing faradaic efficiency (FE) for hydrogen production (Fig. 6(d)). Additionally, air electrode stability tests are performed to monitor any performance degradation. Tests in different atmospheres such as CO₂ or steam are also conducted to assess the chemical stability of the electrode materials.¹⁶⁰ Perovskite air electrodes in P-SOCs are vulnerable to humidity, which can lead to surface hydroxylation, phase decomposition, and a reduction in oxygen ion conductivity, particularly in Ba-containing perovskites. Similarly, CO₂ exposure can cause carbonate formation (BaCO₃, SrCO₃), which blocks active sites, lowers oxygen vacancy concentration, and degrades ionic transport, especially at moderate to high temperatures (500–800 °C). Currently, to enhance overall catalytic efficiency for practical production, researchers are exploring large-area cells configurations, such as 12 × 12 cm² planar cells and tubular cells, in addition to traditional button cells.^161–163

2.4.2. Characterization. The characterization of air electrodes encompasses morphology, structure, and composition, often utilizing techniques such as X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and energy-dispersive X-ray spectroscopy (EDS). Alongside conventional assessments, attention to H⁺/e⁻/O²⁻ three-phase conductivity interactions is crucial for understanding their performance and durability in operational environments.¹⁶⁴

For protonic conductivity, Shao et al. pioneered measuring the H₂ permeation flux through the Pd|Dense Sr₂Sc_0.1Nb_0.1Co_1.5Fe_0.3O_6−δ (SSNCF, air electrode)|Pd membrane.¹⁶⁵ Based on this, hydrogen permeation membrane technique was developed. The SSNCF powder materials are directly compressed, followed by sintering in air into dense pellets, H₂ and N₂ are fed into one side of the chamber, respectively. During the experiments, the hydrogen permeation was detected with gas chromatography (GC) (Fig. 6(e)).⁸⁵ Time-of-flight secondary ion mass spectrometry (ToF-SIMS) is widely used to quantify proton kinetic properties by analyzing the distribution of deuterium/hydrogen (D/H) isotope concentrations in a quenched air electrode (Fig. 6(f)).¹⁵² For instance, Ren et al. combined ToF-SIMS and isotope exchange diffusion profile (IEDP) method to observe stronger distribution of D signal in BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_2.9−δF_0.1 (BCFZYF) sample than undoped.¹⁶⁶ However, its implementation can be challenging due to high costs and the limited availability of specialized equipment. Additionally, the presence of a hydroxide layer on the air electrode surface can lead to an uneven distribution of the measured ion signals. Protonic conductivity of P-SOCs system via vehicle mechanism at low-temperature can be evaluated with a “sandwich-type” membrane conductivity test designed in PEM design. Zhou et al. placed a layer of the oxide between two Nafion layers, each in contact with a gas diffusion layer (GDL) electrode that features a standard polymer electrolyte membrane fuel cells Pt/C catalyst-coated membrane under 75 °C.¹⁶⁷ Proton conduction in the Nafion membrane occurs through hydronium ions (H₃O⁺), and the conductivity can be calculated using the formula σ (S cm⁻¹) = (h (cm))/(R (ohm) × S (cm²)), where h represents the thickness of the oxide layer, R is the resistance, and S denotes the cross-sectional area through which H₃O⁺ migrate.^168–170 The thickness of oxide layer is determined using SEM analysis. This reflects proton conducting via vehicle mechanism, as the Grotthuss mechanism does not involve an H₂O vehicle.

Proton conduction relies on the introduction of protons into the crystal lattice via hydration reactions. Therefore, beyond directly measuring the protons conductivity of air electrode materials, it is important to investigate their hydration properties, including hydration amount, enthalpy, and entropy. The proton adsorption state at elevated temperatures can be preserved by quenching from high to room temperature of H₂O-temperature-programmed desorption (TPD).¹⁶⁷In situ H₂O-TPD has been employed to assess the hydration levels of materials like (Ba,Sr,La)(Fe,Co,Zn,Y)O_3−δ, (La,Sr)(Co,Mn,Ni)O₃, and PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ (PBSCF), in both dry and humidified condition, as a function of temperature and pH₂O. The hydration ability can also be accessed with Fourier transform infrared spectroscopy (FTIR), which sensitively detects adsorbed water and OH˙, OH˙ characteristic peak being detected between 3400 cm⁻¹ and 3800 cm⁻¹.^171,172 For water uptake with special characteristics such as the absorption of water with proton defect results in a noticeable increase in sample weight, thermogravimetric (TG) analysis is employed to measure the proton defect concentration following water uptake.⁷³

The oxygen conduction abilities of air electrode can be measured by oxygen permeation testing which involves measuring O₂ permeation flux detected by GC and ToF-SIMS.⁸⁵ Additionally, the oxygen adsorption behavior can be analysis via O₂-TPD. Chen et al.¹⁵¹ detected the oxygen desorption properties using O₂-TPD over a temperature range of 50 °C to 800 °C. As temperature increases, the active cations undergo thermal reduction, leading to a decrease in their oxidation state and the concurrent release of oxygen. The optimized air electrode materials exhibit a lower initial desorption temperature (303.9 °C), suggesting enhanced surface oxygen exchange kinetics and migration rates within the perovskite lattice. Additionally, charge transfer within the oxide bulk, along with the formation of H⁺ carriers, is often linked to modifications in electronic structure. X-ray photoelectron spectrum (XPS) is a valuable tool for analyzing surface chemical state while X-ray absorption spectroscopy (XAS) provides powerful, element-selective insights into the oxidation states of TMs.^173–175 For example, Zhou et al. conducted operando hard XAS to observe changes in both the geometric and electronic structures of the air electrodes.¹⁶⁷

Oxygen vacancies are essential in influencing the catalytic activity of air electrodes, as they function both as active sites for electrochemical reactions and as pathways for ionic transport.¹⁷⁶ Considering the importance of oxygen vacancies, a series of advanced methods have been employed to detect the existence of oxygen vacancies, concentration, and homogeneity.¹⁷⁷ TG analysis and Iodometric titration¹⁷⁸ is carried out to determine the oxygen nonstoichiometry (δ), while XPS is conducted to analyze the chemical state of oxygen in air electrode materials. In the study of Liu et al., O 1s XPS was fitted in four peaks: 529 eV, 530.6 eV, 531.5 eV, and 532.8 eV, corresponding to lattice oxygen (O_latt), superoxidative oxygen species (O₂²⁻/O⁻), adsorbed oxygen/hydroxide (O₂/OH⁻), and adsorbed water (H₂O), respectively (Fig. 6(g)).¹⁷⁹ The oxygen vacancy density is directly related to the amount of O₂²⁻/O⁻. Raman spectroscopy is used to investigate the vibrational modes which results in shifts in the Raman spectrum or appearance of new peak.^180,181 Additionally, positron annihilation lifetime spectroscopy (PALS) is carried out to detect oxygen vacancies concentrations based on the positron lifetime.¹⁸² And electron paramagnetic resonance (EPR) test can detect peak intensity for oxygen vacancies by comparing the g-factor.¹⁸³ Soft X-ray adsorption spectroscopy (sXAS) can be used to investigate the electronic structures of the O K-edge, which are highly sensitive to oxygen vacancy concentration, electron concentration and hybridized states.¹⁸⁴ Resonant inelastic X-ray scattering (RIXS) can complement XAS measurements by providing detailed information about electron occupancy in d-orbitals and the extent of metal–oxygen hybridization.¹⁸⁵ Scanning transmission electron microscopy (STEM) is employed to analyze the structure of nanomaterials, offering direct atomic-level imaging. Mao et al. utilized STEM to examine the atomic scale structure of RuO₂ and identified defects on the materials surface.¹⁸⁶ The stability of air electrode material is assessed using TG methods, along with comparative analyses of morphology, structure, and composition before and after operation. The thermal expansion coefficient (TEC) of air electrode is confirmed by determining lattice parameters through temperature-dependent XRD testing. Advanced in situ characterization techniques, such as in situ XRD, are employed to monitor the dynamic changes occurring during the operation of the air electrode.^187,188

3. BaMO_3−δ-based perovskite air electrode

Among the diverse family of perovskite oxides, Ba-based simple perovskite BaMO_3−δ (M = Co, Fe) stands out as a particularly attractive option for air electrode in P-SOCs. The larger ionic radii of Ba (1.61 Å) compared to Sr (1.44 Å) and La (1.36 Å), along with its lower electronegativity (Ba: 0.89; Sr: 0.95; Ca: 1.0; La: 1.1; Pr: 1.13, according to Pauling electronegativity), significantly lowers the activation energy required for O²⁻ migration within the crystal lattice.^91,189 This characteristic simultaneously enhances the overall basicity of the oxide ions, promoting adequate conductivity for both oxygen ions and protons. Thus, placing Ba at the A site effectively ensures the transport of carriers. Additionally, its abundant reserves and significantly low cost of Ba (∼1029 USD mt⁻¹) compared to La (∼2621 USD mt⁻¹), Sr (∼8830 USD mt⁻¹), Ca (∼3353 USD mt⁻¹), and Pr (∼73 779 USD mt⁻¹) place it among best candidate for A-site in sustainable industrialization P-SOCs (Prices as of March 2025).^91,190–192 Redox-active TMs like Co and Fe are commonly employed at its B-site for adequate electronic conductivity and enhanced catalytic activity. This chapter gives a comprehensive overview of typical BaMO_3−δ-type perovskites used as high-performance air electrode materials (Fig. 7, performance summarized in Table 1).

Fig. 7 Schematic of recent advances in BaMO_3−δ-type air electrode materials, including (a) heteroatom substitution, (b) phase segregation, and (c) interface regulation.

Table 1 Summary of advanced perovskite air electrode materials

Air electrode	ASR (Ω cm^2) at 3% H2O in symmetrical cell						Single cell					Ref.
	700 °C	650 °C	600 °C	550 °C	500 °C	450 °C	Electrolyte thickness (μm)	Temperature (°C)	PPD (mW cm^−2)	Current density at 1.3 V (mA cm^−2)	Stability (h)
BCO–PBC–LSCF: BaCoO3−δ–Pr1−xBaxCoO3−δ–(La0.6Sr0.4)0.95Co0.2Fe0.8O3−δ; PBCC–BCO: PrBa0.8Ca0.2Co2O5+δ–BaCoO3−δ; BCO–LSCF: BaCoO3−δ–La0.6Sr0.4Co0.2Fe0.8O3−δ; BCO–BPCFY: BaCoO3−δ–Ba0.9Pr0.1Co0.7Fe0.2Y0.1O3−δ; BCT: BaCo0.8Ta0.2O3−δ; BCFT: BaCo0.7Fe0.2Ta0.1O3−δ; BSTC: BaSc0.1Ta0.1Co0.8O3−δ; BCFZnY: BaCo0.4Fe0.4Zn0.1Y0.1O3−δ; PBCF: Pr0.5Ba0.5Co0.7Fe0.3O3−δ; HE-FMCNC: Fe0.6Mn0.6Co0.6Ni0.6Cr0.6O4; HE-PLNBSCC: Pr1/6La1/6Nd1/6Ba1/6Sr1/6Ca1/6CoO3−δ; HE-PBSLCC: Pr0.2Ba0.2Sr0.2La0.2Ca0.2CoO3−δ; HE-BCFZSP: BaCo0.2Fe0.2Zr0.2Sn0.2Pr0.2O3−δ; HE-BCZGZrY: BaCo0.2Zn0.2Ga0.2Zr0.2Y0.2O3−δ; BCCY: BaCo0.7(Ce0.8Y0.2)0.3O3−δ; BC1.5MN: Ba2Co1.5Mo0.25Nb0.25O6−δ; HE-CBSLCC: Ce0.2Ba0.2Sr0.2La0.2Ca0.2CoO3−δ; BSCF–BZCY: Ba0.5Sr0.5Co0.8Fe0.2O3−δ–BaZr0.65Ce0.2Y0.15O3; BSCF–BZCYYb: Ba0.5Sr0.5Co0.8Fe0.2O3−δ–BaZr0.1Ce0.7Y0.1Yb0.1O3−δ; BSCFS: Ba0.5Sr0.5Co0.8Fe0.1Sc0.1O3−δ; LBSCF: La0.1Ba0.4Sr0.5Co0.8Fe0.2O3−δ; BSCFP: Ba0.5Sr0.5(Co0.8Fe0.2)0.95P0.05O3−δ; BSCFF: Ba0.5Sr0.5Co0.8Fe0.2O2.9−δF0.1; BSCFF: Ba0.5Sr0.5Co0.8Fe0.2O2.9−δCl0.1; BSCFE–Er2O3: Ba0.5Sr0.5Co0.72Fe0.18Er0.09O3−δ–Er2O3; C/H-BSCF: Ba0.5Sr0.5Co0.8Fe0.2O3−δ–Ba4Sr4(Co0.8Fe0.2)4O16−δ; N-BCFZYNF: Ba(Co0.4Fe0.4Zr0.1Y0.1)0.95Ni0.05F0.1O2.9−δ; BCFZYN: Ba0.95(Co0.4Fe0.4Zr0.1Y0.1)0.95Ni0.05O3−δ; Ag-BCFZY: Ba0.95Ag0.05Co0.4Fe0.4Zr0.1Y0.1O3−δ; BCFZY: BaCo0.4Fe0.4Zr0.1Y0.1O3−δ; BCF–30Pr: BaCe0.2Fe0.8O3−δ–Pr6O11; BCF36: BaCe0.36Fe0.64O3−δ; S-BCFZY: self-recoverable BaCo0.4Fe0.4Zr0.1Y0.1O3−δ; L-BCFZY: laser-engineering BaCo0.4Fe0.4Zr0.1Y0.1O3−δ; SCM–BCFZY: Sm0.85Cu0.15MnO3−δ–offset BaCo0.4Fe0.4Zr0.1Y0.1O3−δ; BCYF: BaCe0.16Y0.04Fe0.8O3−δ; BFZ: Ba0.875Fe0.875Zr0.125O3−δ; BLFZ0.95: Ba0.95La0.05(Fe0.8Zn0.2)0.95O3−δ; BFCS: BaFe0.6Ce0.2Sc0.2O3−δ; 2Ru–BCF: Ru–anchor BaCe0.125Fe0.875O3−δ; PBSCF: PrBa0.5Sr0.5Co1.5Fe0.5O5+δ; PBCsC: PrBa0.9Cs0.1Co2O5+δ; 16F-PBSCF: PrBa0.5Sr0.5Co1.5Fe0.5O5.84−δF0.16; PBCCHf0.1: PrBa0.8Ca0.2Co1.9Hf0.1O5+δ; PBCFN: PrBaCo1.6Fe0.2Nb0.2O5+δ@PrBaCo1.6Fe0.2Nb0.2−xO5+δ; BZC@PBCZ: BaZrO3@PrBaCo1.92Zr0.08O5+δ; DP-BGCO: BaGdCo2O6−δ; BSCFW@PBSCF: (Ba/Sr)(Co/Fe/W)O3−δ@PrBa0.5Sr0.5Co1.5Fe0.5O5+δ; PCO@PBC: Pr0.1Ce0.9O2+δ@PrBaCo2O5+δ; PFC@PBSCF: Pr0.9Fe0.7Co0.3O3@PrBa0.5Sr0.5Co1.5Fe0.5O5+δ; PCNCFO: Pr0.2Ce0.2Ni0.2Co0.2Fe0.2Ox; PLD-PBSCF: pulsed laser deposition-PrBa0.5Sr0.5Co1.5Fe0.5O5+δ; PBCHf10: PrBaCo1.9Hf0.1O5+δ; PNBFZ: Pr0.8Nd0.2BaFe1.9Zn0.1O5+δ; LSN: La1.2Sr0.8NiO4+δ; PNCu: Pr2Ni0.8Cu0.2O4+δ; NSTF: NaySrzTiuFe1−uO3−δ; NSTF0.3@SC: Na0.3Sr0.7Ti0.1Fe0.9O3−δ@SrCoO3−δ; SCN–PSCN: SrCo0.5Nb0.5O3−δ@PrSrCo1.8Nb0.2O6−δ; SCFN: Sr0.9Ce0.1Fe0.8Ni0.2O3−δ; PNC: PrNi0.5Co0.5O3−δ; PNC–LSCF: Pr2Ni0.5Co0.5O4−δ–La0.6Sr0.4Co0.2Fe0.8O3−δ.
BaCoO3−δ-based air electrode in P-SOCs
BCO–PBC–LSCF	0.06	0.15	0.25	0.45	0.95		10	550	620	850	300 h (600 °C at −1 A cm^−2)	193
PBCC–BCO							10	550	660	690	1500 h (600 °C at −1 A cm^−2)	194
BCO–LSCF	0.04	0.07	0.16	0.35	0.84		10	500	410	500	1150 h (550 °C at −1 A cm^−2)	195
BCO–BPCFY							70	700	750		85 h (700 °C at 0.7 V)	196
Ba0.95CoO3							15	700	1155			197
BCT		0.12	0.23	0.39	0.65	1.2	10	450	450	250	357 h (550 °C at −1 A cm^−2)	132
BCFT							24	700	1272		120 h (700 °C at 0.84 V)	198
BSTC		0.08	0.15	0.27	0.5	1.3	6.5					199
BCFZnY450							11.8	450	320			200
PBCF							6	600	1120	1790	100 h (600 °C at 0.7 V)	201
HE-FMCNC	0.057		0.03				10	600	453		120 h (600 °C at 0.2 A cm^−2)	202
HE-PLNBSCC							6.5	500	660	1100	200 h (500 °C at −0.5 A cm^−2, >10% H2O)	203
HE-PBSLCC	0.06	0.12	0.26	0.75	2.13		10	500	400	280	500 h (600 °C at −0.5 A cm^−2)	204
HE-BCFZSP			0.15				8	600	680	920	120 h (600 °C at −1.3 A cm^−2)	205
HE-BCZGZrY							6.5	600	287		100 h (600 °C)	206
BCCY		0.051	0.11	0.20	0.50	1.49	16.1	450	187		800 h (550 °C)	207
BC1.5MN	0.07	0.13	0.24	0.4	1	3	10	550	540	680	1100 h (550 °C at 0.4 A cm^−2)	208
HE-CBSLCC	0.039	0.088	0.23	0.6	1.65		6	550	1140	840	200 h (600 °C)	209
BaCoO3−δ derivatives in P-SOCs
BSCF–BZCY	0.16							700	418			210
BSCFS600							15	600	566		100 h (600 °C at 0.2 A cm^−2)	211
LBSCF	0.04		0.42				20	600	350	870		211
BSCFP							20	500	425	401	240 h (550 °C)	212
BSCFF		0.06	0.12	0.25	0.54	1.41	16	500	302	190	70 h (550 °C at −0.295 A cm^−2)	213
BSCFC		0.10	0.23	0.53	1.37	4.37	16	650	770	485		213
BSCFE–Er2O3							7	500	525	564	300 h (550 °C at 0.2 A cm^−2)	179
C/H-BSCF		0.11	0.26	0.68	2.13		8	500	1000	1240	500 h (600 °C at −2 A cm^−2)	214
N-BCFZYNF		0.076	0.165	0.407	1.338	4.783	11	500	374	345	80 h (550 °C at −0.5 A cm^−2)	151
BCFZYN							20	550	540		400 h (550 °C at 0.2 A cm^−2)	215
Ag-BCFZY		0.06	0.107	0.242	0.5		10	500	398		600 h (500 °C at 0.75 V)	216
KBCFZY			0.43	0.75			7	550	770		600 h (600 °C at −4 A cm^−2 and 0.5 A cm^−2)	217
BCFZY			0.71				20	500	131			218
S-BCFZY							90	600	350		60 h (600 °C at 0.45 A cm^−2)	218
L-BCFZY		0.1	0.17	0.28			20	550	900		100 h (600 °C at 0.2 A cm^−2)	219
SCM–BCFZY			0.35				14	600	655			220
BaFeO3−δ-based air electrode in P-SOCs
BCF–30Pr	0.039						16	700	1406		100 h (700 °C at 0.3 A cm^−2)	221
BCF36							20	600	550		100 h (600 °C at 0.6 V)	222
BCYF							15.8	450	151		450 h (550 °C at 0.2 A cm^−2)	160
BFZ	0.06	0.25	0.5	1.2	3		6	500	670		230 h (600 °C at 1.2 A cm^−2)	34
BLFZ0.95	0.04	0.07	0.1				10	600	620		250 h (600 °C at 0.9 V)	223
BFCS							8	550	330	610	120 h (600 °C at 0.8 V)	224
2Ru–BCF	0.033	0.2	0.48	1			7	700	1780		200 h (55 °C at 0.7 V)	225
LnBaM2O5+δ-based air electrode in P-SOCs
PBSCF							15.7	500	420		200 h (600 °C at 0.6 V)	152
PBCsC		0.162	0.275	0.531	1.526		8	500	410	310	200 h (650 °C at 0.5 A cm^−2)	226
16F-PBSCF		0.2	0.49	1.2			8	550	300			227
PBCIn	0.06	0.13	0.24	0.52	1.47		10	500	480	310	200 h (re P-SOC)	228
PBCCHf0.1	0.068	0.107	0.269	0.546			7	550	673	836	200 h (600 °C at ±0.5 A cm^−2)	229
PBCFN	0.08	0.24	0.8				10	600	723	1036	200 h (650 °C at ±0.5 A cm^−2)	230
BZC@PBCZ							12	600	551		140 h (600 °C at 0.55 A cm^−2)	231
DP–BGCO	0.13	0.22	0.4	0.1			20	550	386.4		600 h (600 °C at 0.5 V)	232
BSCFW@PBSCF		0.24	0.36	0.80			10	550	570		240 h (650 °C at 0.7 V)	150
PCO@PBC	0.06	0.1	0.3				10	600	870	2090	100 h (650 °C at ±0.5 A cm^−2)	233
PFC@PBSCF		0.25	0.81	1.88			10	550	640		100 h (650 °C at 0.5 A cm^−2)	234
PCNCFO	0.04	0.09	0.18				10	600	910		100 h (650 °C at 0.5 A cm^−2)	235
PLD–PBSCF							7	350	500	250	300 h (500 °C at 1 A cm^−2 and −1 A cm^−2)	236
PBCHf10			0.187	0.374	0.933		7	450	600	400	500 h (600 °C at −0.5 A cm^−2)	237
PNBFZ							16	550	401		24 h (600 °C at 0.7 V)	238
Others
PNCu							26	650	280			239
NSTF							20	600	807			167
NSTF0.3@SC							22.5	600	960			167
SCN–PSCN							8	550	600	750	120 h (600 °C at 0.85 V)	240
SCFN			0.2	0.57	2.05		26	500	225	179	290 h (550 °C at 0.8 V)	171
PNC							16	450	640	1600	200 h (600 °C at 0.75 V)	241
PNC–LSCF							30	550	662		200 h (700 °C at 0.8 A cm^−2)	242

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3.1. BaCoO_3−δ-based series

Co, as a B-site cation, is particularly effective for reactions involving electron exchange, such as OER and ORR. This is due to the strong hybridization between the metal 3d and O 2p orbitals,^32,243 as well as the close values of their intra-atomic exchange energy and crystal electric field splitting energy, which leads to small energy differences between their spin states and oxidation states.^244–247 The unique properties combined with the aforementioned of Ba collectively make BaCoO_3−δ (BCO) an exceptional choice for P-SOCs.^193,194,248

However, BCO's hexagonal structure of BCO restricts charge carrier mobility within the lattice,^249–251 which limit its effectiveness as an independent air electrode.^132,252 It is, consequently, utilized as a reinforcement material, functioning as a modified layer to enhance catalytic activities as well as air electrode durability as in reported in situ studies using PrBa_0.8Ca_0.2Co₂O_5+δ, Ba_0.9Pr_0.1Co_0.7Fe_0.2Y_0.1O_3−δ, and La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ (LSCF).^193–196 The BCO coatings additionally mitigate toxic effects of Cr from interconnectors owing to their ability to reduce Cr accumulation on the electrode surface.²⁵³ Recent research efforts have increasingly focused on the bulk structural regulation of BCO for its ultimate autonomy. Atomic substitution techniques have been widely employed to modify their air electrodes phase structure, evolving from single-element substitutions (like Pr, La, Nb, Zn, Sc, etc.)^254–262 to double and then multi-element modifications (Fig. 7(a)).^{197,259,263,264} Notably, heteroatomic doping with high-valence cations like Nb⁵⁺, Ta⁵⁺, or Mo⁶⁺ at the B-site has proven particularly effective. For instance, Kim et al. developed cubic BaCo_0.8Ta_0.2O_3−δ (BCT) by incorporating Ta⁵⁺, which has highly corrosion-resistance, stable valence state, and a lower electronegativity (Ta: 1.5; Nb: 1.6; Mo: 2.16).¹³² The inclusion of Ta⁵⁺ stabilized the high-symmetry cubic perovskite structure, significantly enhancing phase stability, conductivity, and catalytic activity for both OER and ORR (Fig. 8(a)). Thus, BCT air electrode achieved remarkable PPD of 2.26 W cm⁻² at 650 °C and electrolysis current density of 1.1 A cm⁻² at 1.3 V and 550 °C in P-SOC. Combining two dopants on BCO-type electrodes, like Fe–Ta,^198,264,265 Sm–Zr,²⁶³ and Co–Zr,²⁶⁶ creates synergistic effects that enhance electrical, ionic, and catalytic properties to levels superior to single-element doping. Inspired by dual-doping strategy observed in SrCoO₃ for O-SOCs,²⁶⁷ Kim et al. developed Sc–Ta co-doped BaSc_0.1Ta_0.1Co_0.8O_3−δ (BSTC), which stabilizes the cubic perovskite structure.¹⁹⁹ Dual-doping strategy can lead to reduced TEC, improved water uptake, hydration properties, proton transport capabilities, and higher resistance to CO₂ (Fig. 8(b)). The BSTC air electrodes in P-SOCs achieved PPD of 3.15 W cm⁻² and 2.25 W cm⁻² at 650 °C and 600 °C in FC mode, and a current density of 4.21 A cm⁻² at 1.3 V and 650 °C in EC mode. Additionally, Sr–Fe co-doped Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF) inspired from O-SOCs has seen widespread practical implementation in P-SOCs, which is discussed in detail in Section 3.2. Multi-element doped materials,^200,201 represented by BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZY) inspired by electrolyte materials, demonstrate exceptional electrochemical potential and serve as the foundation for numerous research studies. A detailed exploration of BCFZY is presented in Section 3.2. In the design of multi-element substitution strategy, high-entropy design has emerged as a significant topic.^{202,268–270} High-entropy perovskite oxides (HEPOs) that contain five or more elements in nearly equimolar ratios exhibit unique properties due to complexity of their localized environments and expected broad adjustability afforded by the “cocktail” effect.²⁷¹ Unlike traditional enthalpy-driven phase separation during multi-doping processes,^272,273 HEPOs are stabilized by configurational entropy. In addition, they display a range of unique properties that differ from those of individual oxides including unique crystal structures, enhanced thermal and electronic characteristics, superior catalytic activity, improved electrochemical performance, efficient ionic transport, and distinct thermal expansion behaviors.^{203,274–278} He et al. developed A-site HEPO Pr_0.2Ba_0.2Sr_0.2La_0.2Ca_0.2CoO_3−δ (HE-PBSLCC), that effectively combined high conductivity and low TEC of Pr, larger ionic radius of La³⁺, and chemical stability of Ca²⁺. This composition led to reduced TEC (23.8 × 10⁻⁶ K⁻¹) and effective phase segregation suppress.¹⁸⁷ HE-PBSLCC achieved a PPD of 1.51 W cm⁻² and a current density of 2.68 A cm⁻² at 1.3 V at 650 °C, while operational lifetimes exceeding 270-h FC and 500-h EC opeartion.¹⁸⁷ Sun et al. constructed B-site HEPO BaCo_0.2Fe_0.2Zr_0.2Sn_0.2Pr_0.2O_3−δ, demonstrating integrated triple-phase conduction, along with exceptional structural stability in high concentration steam.^205,206 In addition to bulk structure modification, atomic substitution can also facilitate the formation of multiple phases within a single bulk grain which is called surface segregation or exsolution. Improper t of certain elements in perovskites can lead to phase separation, resulting in the formation of two or more distinct phases with varying elemental compositions. The self-assembly strategy for creating hybrid electrodes ensures a more homogeneous phases distribution and improved phase contact compared to straightforward physical mixing, thereby reducing energy barriers for charges conduction.²⁰⁷ Song et al. introduced a novel self-assembled BaCo_0.7(Ce_0.8Y_0.2)_0.3O_3−δ electrode for P-SOFC, incorporating Ce and Y cations into the B-site of the BCO base.²⁰⁷ This material undergoes self-assembly into nanocomposite consisting of mixed H⁺/e⁻ phase, mixed O²⁻/e⁻ phase, and the BCO phase during calcination. Enhanced electrochemical activity and ionic transport were achieved due to the synergistic effects of the distinct phases. The composite electrode Ba₂Co_1.5Mo_0.25Nb_0.25O_6−δ (BC1.5MN), as reported by He et al., undergoes elegant decomposition into BCO and double-perovskite Ba_2−xCo_1.5−xMo_0.5Nb_0.5O_6−δ (DP-BCMN).²⁰⁸ This transformation enables the electrode to achieve a PPD of 1.17 W cm⁻² and a current density of 2.04 A cm⁻² at 1.3 V and 650 °C. A-Site HEPO Ce_0.2Ba_0.2Sr_0.2La_0.2Ca_0.2CoO_3−δ (HE-CBSLCC), developed by He et al.,²⁰⁹ self-assembles into a three-phase heterostructure under specific processing conditions. This structure includes a deficient Ce_0.2−yBa_0.2Sr_0.2−xLa_0.2−xCa_0.2CoO_3−δ (CD-CBSLCC), along with CeO₂, and La_0.5Sr_0.5CoO_3−δ (LSC) nano-catalysts. Oxygen reduction occurs across the entire air electrode surface, with water formation occurring predominantly at or near CD-CBSLCC. CeO₂ phase can either donate or accept H⁺, thereby facilitating the ORR and OER kinetics in reversible P-SOCs. The electrode achieves a PPD of 1.66 W cm⁻² and a current density of 1.76 A cm⁻² at 1.3 V and 600 °C. They also exhibit excellent operational stability, with 200 h in FC mode, 200 h in EC mode, and 548 h in reversible cycling at 550 °C.

Fig. 8 Recent advances in BaMO_3−δ-type air electrode materials. (a) XRD patterns of BCT. Reproduced with permission.¹³² Copyright 2023, Royal Society of Chemistry. (b) Tolerance factors of bimetal-doped BMM’C. Reproduced with permission.¹⁹⁹ Copyright 2024, Wiley-VCH. (c) O-site substitution with F/Cl of BSCFF. Reproduced with permission.²¹³ Copyright 2023, Elsevier. (d) Schematic of phase content-controlled hybrid electrode. Reproduced with permission.²¹⁴ Copyright 2024, Springer Nature. (e) Reaction mechanism on the surface of ABCFZY electrode in humidified air. Reproduced with permission. Copyright 2024, Royal Society of Chemistry.²¹⁶ (f) Schematic of the laser ablating process and SEM micrographs of holes and S-shape. Reproduced with permission.²¹⁹ Copyright 2024, American Chemical Society. (g) Mechanism of the SCM limited reactive bonding “glue”. Reproduced with permission.²²⁰ Copyright 2024, Wiley-VCH.

3.2. Eye-catching BaCoO_3−δ derivatives

Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF) is the center of attention as air electrode material for intermediate temperature O-SOCs.^279,280 In 2004, BSCF made its debut as a cathode for O-SOFCs, achieving remarkable PPDs of 1.01 W cm⁻² at 600 °C and 402 mW cm⁻² at 500 °C.^281,282 With mixed conductivity of oxygen ions and a reported proton exchange k value of 3.85 × 10⁻⁶ cm s⁻¹ at 500 °C in 0.21 atm pO₂, the cubic perovskite BSCF has been identified as air electrode candidate for P-SOCs, too.^158,281 To extend three-phase boundaries and minimize the TEC mismatch between the electrode and electrolyte, BSCF has been composited with various electrolyte materials including BaCe_0.7Zr_0.1Y_0.2O_3−δ,²⁸³ BaCe_0.9Y_0.1O₃,²⁸⁴ BaCe_0.8Zr_0.1Y_0.1O_3−δ,²⁸⁵ BaCe_0.4Zr_0.4Y_0.2O_3−δ,²⁸⁶ BaZr_0.65Ce_0.20Y_0.15O₃,²¹⁰ and BZCYYb.²⁸⁷ However, this kind of composite design may compromise the conductivity of air electrode.

Due to the differences in the working principles of P-SOCs and O-SOCs—specifically, that in O-SOCs, water vapor is generated and consumed on the fuel electrode side^288,289—the electrode materials for the two types of cells are not fully interchangeable.²⁹⁰ BSCF as P-SOCs air electrode, for example, degrades in high-concentration steam environments.²⁹¹ To tackle this problem, atomic substitution (like Sc, La, P) has been attempted to BSCF to stabilize its structure.^211,212,292 Besides cation regulation, anion (O-site) substitution has recently emerged as an effective method for tailoring perovskite air electrode (Fig. 7(a)).^293–295 Anion doping involves substituting lattice oxygen or occupying oxygen vacancies with anions, typically at a stoichiometric ratio of 0.1 to 0.2, thereby modifying the basicity, valence balance, electronegativity, and band structure. To date, halogens (F⁻ and Cl⁻) have been used as anionic dopants.^{293,296–298} Halogen doping can reduce the basicity of perovskite oxides without compromising their proton conductivity, thus enhancing their chemical stability against acidic gases like CO₂.^293,294,299 Halogen doping can alleviate F⁻ diffusion from the electrolyte.³⁰⁰ By introducing halogen, the oxygen vacancy concentration is decreased to maintain electron neutrality, with halogen potentially occupying these vacancies, further lowering vacancy concentration.^22,293,297 The fluorinated Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCFF) air electrode, for instance at 650 °C, achieved a PPD of 977 mW cm⁻² and a current density of 950 mA cm⁻² at 1.3 V.²¹³ Chen et al. further compared the effects of F⁻ and Cl⁻ (Fig. 8(c)), noting that although both have similar valence states, they differ significantly in electronegativity (χ_F: 3.98 > χ_O: 3.44 > χ_Cl: 3.16) and ionic radii (F⁻: 1.33 Å < O²⁻: 1.40 Å < Cl⁻: 1.81 Å).²¹³ The higher electronegativity of F⁻ enhances its electron-attracting ability, thereby increasing the polarity of the M–O bonds. In P-SOCs, the fluorinated BSCFF air electrode outperformed chlorine Ba_0.5Sr_0.5Co_0.8Fe_0.2O_2.9−δCl_0.1 (BSCFC) air electrode in terms of both D_chem and k_chem. The fluorinated BSCFF also exhibited superior electrochemical performance in both FC (977 mW cm⁻²vs. 770 mW cm⁻² for chlorine BSCFC at 650 °C) and EC (950 mA cm⁻²vs. 485 mA cm⁻² for chlorine BSCFC at 1.3 V and 650 °C).

Recent research indicates that exceeding material tolerance limits through atomic substitution can form self-assembled multi-phase nanocomposites within a single grain. By carefully controlling concentration of high-valence cations (like Mo⁶⁺ and W⁶⁺) or large-radius cations (like Er, Mg, and Ag), it is possible to customize electrical conductivity, catalytic activity, and structural stability within composite air electrode (Fig. 7(a)).^301,302 Liu et al., for instance, introduced Er into BSCF, which lead to the formation of a nanocomposite featuring predominant cubic perovskite phase (Ba_0.5Sr_0.5Co_0.72Fe_0.18Er_0.09O_3−δ), along with a minor Er₂O₃ component.¹⁷⁹ The Er₂O₃ phase acts as a rapid H⁺ transport channel, thereby facilitating kinetics of both ORR and OER. While the self-assembly strategy effectively creates strong interactions between hybrid phases, the dynamically changing composition and inconsistent elemental composition render concerns about durability. Liu et al. efficiently controlled the phase contents of hybrid material by adjusting the stoichiometric ratio of A-site and B-site elements, while ensuring the consistency of elemental composition (Fig. 8(d)). The resulting Ba_1.5Sr_1.5Co_1.6Fe_0.4O_7−δ consisted of 57.26 wt.% cubic Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF) and 42.74 wt.% hexagonal Ba₄Sr₄(Co_0.8Fe_0.2)₄O_16−δ (H-BSCF).²¹⁴ This hybrid perovskite integrated the superior oxygen activation and conductivity of cubic phase and strong hydration reaction and abundant of hexagonal phase. The strengthened interaction between the two phases enhanced the structural and chemical stability.

Since 2013, BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZY) has garnered significant attention from researchers, particularly after O’Hayre's group proposed its application in P-SOFC.^103,266 Numerous studies have since been published exploring the chemistry and electrochemistry of BCFZY based on typical compositing (BCFZY–BaZr_0.5Y_0.5O₃ and BCFZY–NiO)^303,304 and substitution (Zn,³⁰⁵ Mg,^200,306 Sc,³⁰⁷ Y³⁺,³⁰⁸ F,^166,213 and A-site-deficient) strategies. Recent investigations on BCFZY increasingly advocate for a dual-functional strategy that combines both bulk and interfacial modifications instead of focusing in one aspect.^{195,215,309–311} Chen et al. involved Ni, F co-doped Ba(Co_0.4Fe_0.4Zr_0.1Y_0.1)_0.95Ni_0.05F_0.1O_2.9−δ (N-BCFZYNE), resulting in precipitation of NiO as a surface catalyst.¹⁵¹ The higher electronegativity of F⁻ compared to O²⁻ enhances the polarity of bonds, facilitating H₂O adsorption on and leading to the formation of additional protonic defects. The F⁻ introduction decreases the negative charge on lattice oxygen, which in turn weakens the O⋯H interaction, thereby lowering the barrier for proton diffusion. The in situ generated metal oxide catalyst demonstrated enhanced surface reaction activity. This approach achieved improvement both internally (via highly reactive metal–oxygen bonds) and externally (via enabled nanoscale catalysis). Secondary phase nanoparticles (NPs) exsolution can occur under various conditions including water vapor, reducing/oxidizing atmospheres, high temperatures, and electric fields.^215,312 Kim et al. designed steam-induced Ag metal nanocatalysts that functions solely as an electronic conductor, without engaging in proton uptake and diffusion.²¹⁶ Building on this, through steam-introduced strategy, Park et al. introduced exsolved BaO_x species in alkali metal ions (Li⁺, Na⁺, K⁺)-doped BCFZY, which served as active sites for ORR/OER (Fig. 8(e)).²¹⁷ The driving force for Ba dissolution arises from lattice strain caused by differences in ionic radius and defect interactions, while BaO_x is thermodynamically favored due to its negative Gibbs free energy. In addition to water-induced exsolution, treating Ba_0.95Ag_0.5Co_0.8Nb_0.1Ta_0.1O_3−δ in a reductive atmosphere³¹³ and employing voltage-driven exsolution in Ln_0.2Ba_0.8Co_0.7Fe_0.3O_3−δ^314,315 were also effective in generating NPs catalyst on the surfaces.

Exsolved NPs like Co, Fe, and Ni may slowly dissolve back into the perovskite lattice at elevated temperatures and oxidizing environment.^316,317 Inspired by this, Wang et al. developed a self-recoverable BCFZY (S-BCFZY) electrode for P-SOFCs, featuring a reversible exsolution/dissolution mechanism based on the incorporation of NPs.²¹⁸ Under cathodic conditions, the weak Co–O and Fe–O bonds effectively lower the energy barrier for oxygen release, promoting the generation of and H⁺.²¹⁸ In anodic conditions, conversely, these weakened TM–O bonds facilitate the segregation of TM atoms (Co⁰ and Fe⁰) to the perovskite surface, leading to NPs formation and enhanced electrocatalytic activity.^318,319 This multifunctional modification, spanning from the bulk to the surface, positively influences the catalytic process. However, significant differences in the elemental composition of various phases can create thermal strains that are difficult to mitigate and can impair both catalytic activity and stability. Additionally, elemental migration may lead to confusion regarding the actual composition of the hybrid electrode, particularly under high-temperature operational conditions. The exsolution process may inadvertently induce structural modifications in the parent oxide, potentially resulting in the development of RP phases or unwanted cation precipitation.³²⁰ Therefore, the long-term effectiveness of this approach must be thoroughly assessed.

Modifying air electrodes involves not only the electrode materials themselves but also the electrode/electrolyte interface (Fig. 7(c)).^321,322 The blunt surface of perovskite-type proton-conducting electrolytes contributes to interfacial resistances, leading to significant discrepancies in the ASR compared to theoretical values.^323,324 To alleviate interface problem, Zhou et al. developed a novel picosecond laser ablation technique that precisely modifies the electrolyte surface (Fig. 8(f)).²¹⁹ This method enables uniform removal of the electrolyte's top layer and creates various patterns, such as cross, S-shape, and hole designs, which increase the effective surface area and establish strong bonding anchors between the BCFZY electrode and BZCYYb electrolyte (Fig. 8(f)). Another interface-controlled approach is to improve the compatibility between electrode materials and electrolytes, especially thermomechanical compatibility.^170,325 Yu et al. employed negative thermal expansion oxide Sm_0.85Cu_0.15MnO_3−δ (SCM with TEC of −5.93 × 10⁻⁶ K⁻¹) to offset the thermal expansion of BCFZY.²²⁰ SCM can also strengthen electrode structure and adhesion, and provide acceptable oxygen–reduction–reaction activity (Fig. 8(g)). Thus, this strategy achieves a PPD of 1.455 W cm⁻² at 700 °C.

3.3. BaFeO_3−δ-based perovskite

Given the significant differences in TEC between cobalt-based electrodes and electrolytes (e.g., BSCF: 24.7 × 10⁻⁶ K⁻¹; BCFZY: 22.9 × 10⁻⁶ K⁻¹; BZCYYb: 10.8 × 10⁻⁶ K⁻¹; BZCY: 9.3 × 10⁻⁶ K⁻¹),^199,207,326 along with the easy evaporation of Co at elevated temperatures and economic considerations (Co: 66.5 USD kg⁻¹; Fe: 6.7 USD kg⁻¹),¹⁹⁰ Fe emerges as a viable alternative.^{160,309,327–329} Fe shares comparable reducibility with Co and, unlike Co-based perovskites, Fe-based perovskites are more easily hydrated and demonstrate improved chemical and thermal stability. The TEC of Fe-based air electrode materials shows enhanced compatibility with proton-conducting electrolytes.³³⁰ Among all Fe-based perovskite, BaFeO₃ (BFO) exhibit superior H⁺ conductivity compared to placing Sr or La at A-site.³² However, undoped BFO crystallizes in a hexagonal system with low symmetry, poor oxygen ion conductivity, limited tolerance to CO₂ and H₂O, thereby hindering its effectiveness as a high-performance air electrode.^331–333 To mitigate these challenges, doping (e.g., La, Zn, Sm, Cl, and F) and compositing (e.g., Sm_0.2Ce_0.8O_2−δ–Ba_0.5Sr_0.5Fe_0.8Sb_0.2O_3−δ) strategies are frequently employed.^{332,334–343} Recently, Lu et al. combined the two modifications to develop a BaCe_0.2Fe_0.8O_3−δ–Pr₆O₁₁ (BCF–30Pr) composite cathode for P-SOFC through impregnation technique.²²¹ The incorporation of Pr₆O₁₁ increased the specific surface area of electrode, reduced the Fe³⁺ content, and enhanced the concentration of on the BCF surface. These changes improved both the ORR activity and H⁺ conduction, leading to a PPD of 1406 mW cm⁻² at 700 °C.

Inspired by the design of BCFZY, Ce, Zr and Y are also favored dopants in BFO phase.^{223,344–346} The Ce-doped BaCe_0.36Fe_0.64O_3−δ (BCF36) developed by Tong et al. features a cubic BFO phase with high O²⁻/e⁻ conductivity and an orthorhombic BaCeO₃ phase with H⁺ transfer, thus achieving a PPD of 1525 mW cm⁻² at 700 °C.²²² By tailoring the Ce/Y ratio in bimetal-doped Ba(Ce_0.8Y_0.2)_xFe_1−xO_3−δ (BCYF), a O²⁻/e⁻ conducting single-phase transforms into a triple-conducting multi-phase composite.³⁴⁷ The optimized BCYF nanocomposite displayed improved e⁻ conductivity, ORR activity, thermo-mechanical compatibility, and CO₂ tolerance, owing to the strong interaction between O²⁻/e⁻ and H⁺/e⁻ conducting phases, as well as the optimized dual-phase composition. The BaCe_0.16Y_0.04Fe_0.8O_3−δ cathode demonstrates 450 h durable operation at 550 °C for P-SOFC. Based on BFO, Wang et al. combined A- and B-site co-substitution to construct Ba_0.875Fe_0.875Zr_0.125O_3−δ (BFZ) with the formation of and while retaining structural, which can catch up power outputs of state-of-the-art Co-based BCFZY with a PPD of 0.67 W cm⁻² at 500 °C and operational stability of 230 h at 600 °C.^22,34 In the context of BFO-based perovskites used in reversible P-SOCs, Wang et al. developed BaFe_0.6Ce_0.2Sc_0.2O_3−δ (BFCS) perovskite characterized by optimized Fe 3d-e_g orbital occupancy, a higher concentration of , and enhanced Fe⁴⁺–O²⁻ interactions which promotes the activation and mobility of lattice oxygen.²²⁴ The BFCS air electrode for P-SOC achieves encouraging output performance in both EC (1.55 W cm⁻²) and FC (−2.96 A cm⁻² at 1.3 V) modes at 700 °C. Single-atom catalysis, an emerging technology, has recently been utilized in Fe-based air electrodes. Zhao et al. developed a novel air electrode (2Ru-BCF) by integrating single-atom Ru rivets onto BaCe_0.125Fe_0.875O_3−δ (BCF).²²⁵ The Ru atoms in this structure adopt a distinctive 4-coordinate Ru–O–Fe configuration, which not only promotes reverse hydrogen spillover but also functions as an active site for P-ORR. This optimized 2Ru–BCF cathode (2 wt% Ru) delivers a remarkable PPD of 1.78 W cm⁻² at 700 °C and demonstrates stable operation over 200 h.

4. LnBaM₂O_5+δ-based perovskite

LnBaM₂O_5+δ-based double perovskites feature A-sites occupied by 4f lanthanide and alkaline earth cation Ba. The B-sites host abundant TM cations, including Mn, Fe, Co, and Ni. These materials exhibit enhanced electrochemical performance due to their higher D_chem and k_chem (performance summarized in Table 1).^348,349 The atomic level of the separation of cations with an alternative layer structure happens due to the large difference in the A or B sites for the two different cations with equal amounts. This chapter offers an in-depth review of typical LnBaM₂O_5+δ-type perovskites employed as high-performance materials for air electrodes (summarized as Fig. 9).

Fig. 9 Schematic of recent advances in LnBaM₂O_5+δ-type air electrode materials, including (a) heteroatom substitution, (b) phase segregation, (c) phase infiltration, (d) atomic trapping, and (e) interface regulation.

4.1. LnBaCo₂O_5+δ-based perovskite

The presence of Co on the B-sites has proven advantageous for both OER and ORR processes. LnBaCo₂O_5+δ-type double perovskite oxides including NdBaFe_2−xCo_xO_5+δ,³⁵⁰ NdBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ,^43,351 and Ba_1−xGd_0.8La_0.2+xCo₂O_6−δ^352,353 with triple-conducting capability, have been demonstrated as effective air electrodes for P-SOCs. Among cobalt-based double perovskites, incorporation of Pr²⁺/Pr³⁺ in the A site lead to favorable TECs, high conductivity, and strong catalytic activity for OER and ORR reactions.^354,355 Thus, PrBaCo₂O_5+δ (PBC) system is widely studied as air electrode material,^{84,231,335,356–366}e.g., famous PrBa_0.5Sr_0.5Co_2−xFe_xO_5+δ developed by Choi et al., which can achieve PPD exceed 500 mW cm⁻² at 500 °C and long-term durability under CO₂.^84,152 Doping with aliovalent cations including Zn, Ca, Sr, Fe, Ca, and Pd in PBC has been explored to enhance oxygen transport, match TECs with electrolyte, and improve stability.^{335,367–371} Xu et al. introduced a low-Lewis-acid-strength Cs⁺-doped PrBa_0.9Cs_0.1Co₂O_5+δ (PBCsC) electrode (Fig. 9(a)).²²⁶ A-Site dopants with lower acidity facilitate easier H⁺ uptake (Fig. 10(a)). The polarization of ionic Lewis acid strength at the A site shifts electrons pairs, leading to increased and surface exchange kinetics (Fig. 10(b)).³⁷² PBCsC electrode demonstrates excellent performance in both FC mode (1.66 W cm⁻²) and EC mode (−2.85 A cm⁻² at 1.3 V) at 650 °C. Regarding anion substitution, the high electronegativity of F weakens the bond strength between transition-metal 3d orbitals and O 2p states (metal–oxygen bonds) in perovskites, thereby enhancing oxygen surface exchange and diffusion processes.^227,343 The PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5.84−δF_0.16 electrode for P-SOFC exhibits an improved PPD of 510 mW cm⁻² at 600 °C.²²⁷ The long-term durability of perovskite electrodes is significantly influenced by the alkaline earth cations (e.g., Ba and Sr) surface segregation,³⁷³ which stems from composition, temperature, acidity or alkalinity, oxygen/steam partial pressure, and polarization current (Fig. 9(c)).³⁷⁴ The resultant insulating impurities impede the transfer of electrons and ions in proton-involved ORR and OER processes.^374–376 Researchers explored various strategies to mitigate the negative impacts of segregation, represented by phase structure engineering.^18,26,313 Du et al. utilized In³⁺, which has a large ionic radius (In³⁺ = 0.8 Å, Co³⁺ = 0.545 Å, and Co⁴⁺ = 0.53 Å) and is recognized for enhancing sinterability and stability in proton-conducting electrolytes,^377–379 to develop a PrBaCo_1.9In_0.1O_5+δ (PBCIn) air electrode.²²⁸ The composition-driven formation of composite phases in the PBCIn electrode consists of a primary phase of deficient PrBa_0.95Co_1.85In_0.09O_5+δ and a secondary cubic perovskite phase BaCo_0.85In_0.15O_3−δ.²²⁸ The synergistic effect of these two phases enables PBIn to show 2.25 W cm⁻² in FC mode, −4.41 A cm⁻² at 1.3 V in EC mode at 700 °C and 210 h durability under 600 °C. High-valence cation Hf⁴⁺-doped PrBa_0.8Ca_0.2Co_1.9Hf_0.1O_5+δ (PBCCHf_0.1) naturally is regarded as double perovskite backbone, PrBa_0.8−xCa_0.2Co_1.9Hf_0.1−xO_5+δ (PBCCHf_0.1−x), with nano-sized BaHfO₃ forming on the surface (Fig. 10(c)).^229,380,381 The in situ exsolution of BaHfO₃ NPs induces additional oxygen vacancies on the surface of Ba- and Hf-deficient PBCCHf_0.1−x. Xu et al. demonstrated that Fe and Nb co-doped PBC form a heterostructure under steam-driven conditions (Fig. 10(d)), featuring in situ exsolved Nb-deficient PrBaCo_1.6Fe_0.2Nb_0.2−xO_5+δ (Nb-deficient PBCFN) on a double perovskite backbone, PrBaCo_1.6Fe_0.2Nb_0.2O_5+δ (Nb-rich PBCFN).²³⁰ The combination of in situ formed Nb-deficient NPs and the Nb-rich parent perovskite significantly faster surface exchange process, enhancing both catalytic activity (PPD of 1.059 W cm⁻² and current density of 2.148 A cm⁻² at 1.3 V at 650 °C) and durability (200 h). In situ exsolution of BaZrO₃ NPs from the host PrBaCo_1.92Zr_0.08O_5+δ (BZO@PBCZ) occurs under oxidizing conditions, facilitating the liberation Ba and Zr cations and forming proton transfer channels (Fig. 10(e)),²³¹ Thereby accelerating the oxygen reduction catalytic activity. Cation-nonstoichiometric Ba_1+xGd_1−xCo₂O_6−δ spontaneously grows into the double perovskite BaGdCo₂O_6−δ (BGCO), anchored by BCO NPs under thermal-driven conditions (Fig. 10(f)).²³² The synergy between the two components leads to exceptional performance, with the mixed O²⁻/e⁻-conducting perovskite-derived oxide exhibiting superior catalytic activity for the ORR, while the double perovskite structure enhances bulk H⁺ conductivity, thus increasing the number of available reaction sites.

Fig. 10 Recent advances in LnBaCo₂O_5+δ-type air electrode materials. (a) Structure of PBCsC, (b) comparisons of different oxygen species contents of PBS and PBCsC. Reproduced with permission.²²⁶ Copyright 2023, American Chemical Society. (c) Schematic of Hf-doped PBCCHf_0.1. Reproduced with permission.²²⁹ Copyright 2024, Elsevier. (d) Schematic of PBCFN. Reproduced with permission.²³⁰ Copyright 2022, Wiley-VCH. (e) Schematic of BZO@PBC. Reproduced with permission.²³¹ Copyright 2022, Wiley-VCH. (f) Schematic of BGCO–BCO hybrid catalyst. Reproduced with permission.²³² Copyright 2024, Wiley-VCH. (g) Schematic of PFC@PBSCF in the presence of contaminants (Cr and steam). Reproduced with permission.²³⁴ Copyright 2022, Wiley-VCH. (h) Schematic for the fabrication procedures of PCNCFO–PBC. Reproduced with permission.²³⁵ Copyright 2024, Wiley-VCH. (i) Structural evolution during reverse atom capture process. Reproduced with permission.¹⁵⁰ Copyright 2024, Wiley-VCH. (j) Mechanism underlying the thermo-mechanical enhancement of TEC gradient formation. Reproduced with permission.³⁸² Copyright 2024, Wiley-VCH. (k) Cross-sectional SEM image of PBSCF/BZCYYb composite interlayer. Reproduced with permission.²³⁶ Copyright 2024, Wiley-VCH.

Surface decoration of nano-catalysts, typically on the order of tens of nanometers, on cathodes through a cost-effective infiltration process has garnered considerable attention (Fig. 9(c)).^383,384 The porous skeleton facilitates significant electronic and ionic conductivity and the infiltrated coating enhances both catalytic activity and durability, e.g., Gd_0.1Ce_0.9O_2−δ-infiltrated PBC.³⁸⁵ Pei et al. incorporated a fluorite-based Pr_0.1Ce_0.9O_2+δ (PCeO) catalyst coating on PBC skeleton, demonstrating a much-reduced polarization resistance (reduced 58%), improved performance (PPD of 1.21 W cm⁻² at 650 °C), and Ba segregation inhibition.²³³ The infiltration method promotes a well-mixed combination of raw materials at the atomic level, allowing for lower synthesis temperatures, and extends the length of the three-phase boundaries,^386–388 where the electron, ion, and gas phases converge. For instance, at 700 °C, the Pr_0.9Fe_0.7Co_0.3O₃ (PFC) coating formed on PBSCF base after infiltration resulted in substantial performance improvements and markedly poisoning (including Cr and steam) tolerance (Fig. 10(g)).^234,389,390 Besides creating a surface coating, the infiltrated solution may interact with the perovskite matrix. Gao et al. designed a multi-cationic oxide nano-catalyst, Pr_0.2Ce_0.2Ni_0.2Co_0.2Fe_0.2O_x (PCNCFO), anchored on the surface of the PBC electrode via infiltration (Fig. 10(h)).²³⁵ The cerium oxide in the PCNCFO coating reacts dynamically with Ba segregated from the PBC, forming BaCeO₃, which stabilizes the PBC phase structure, enhances proton conduction and transfer, and accelerate oxygen surface exchange. Consequently, the PCNCFO-coated PBC cathode achieved a high PPD of 1.31 W cm⁻² at 650 °C.

Recently, an innovative method involving atomic trapping at high temperatures has been effective in modifying the surface chemistry of perovskites to tackle segregation.^391–393 In the (La_0.6Sr_0.4)_0.95Co_0.2Fe_0.8O_3−δ perovskite, strontium atoms can be selectively extracted using an acidic MoO₄ trap, forming inactive SrMoO₄, which enhances durability.³⁹⁴ Building on this atomic capture approach, Zhao et al. introduced liquid-phase dispersible (NH4)₁₀W₁₂O₄₁ onto the PBSCF surface, allowing for the capture of segregated barium and strontium cations from the PBSCF matrix (Fig. 9(d)).¹⁵⁰ This process results in the formation of a heterostructure (Ba/Sr)(Co/Fe/W)O_3−δ (BSCFW)@PBSCF (Fig. 10(i)). The novel P-SOFC cathode achieved remarkable performance, attaining a PPD of 1.32 W cm⁻² at 650 °C, alongside impressive long-term durability for 240 h.

Several methods exist for modifying the bulk structure of PBC-type air electrode materials including morphology engineering,³⁹⁵ heterostructure construction,^{234,396–398} compositing,^399–402 and microstructure design.^203,271,403 However, the large TECs of cobalt-based PBC materials (10.8 × 10⁻⁶ K⁻¹ for BZCYYb, 23.7 × 10⁻⁶ K⁻¹ for PBSCF) restrict their practical application by leading to mechanical incompatibility with electrolyte and poor electrolyte/electrode interface adhesion.^400,404 Negative thermal expansion materials represented by Y₂W₃O₁₂ and Sm_0.85Cu_0.15MnO_3−δ are added to the air electrode to offset the thermal expansion.^170,220,325 However, inadequate mechanical mixing and poor physical contact between powders can result in delamination due to the significant variation in the thermal expansion coefficient of the electrode material and added powders. Additionally, the use of negative thermal expansion materials is constrained by phase transitions triggered by temperature or pressure changes, as well as its low stiffness and hygroscopic nature. To address these challenges, Gao et al. proposed composite electrodes with a TEC gradient.³⁸² The transition phase, generated in situ due to topological atomic effects, helps bridge the thermal behavior gap and maintains thermo-mechanical stability. This approach also ensures chemical stability by preventing water vapor infiltration into the electrode. In this recently published study, the formation of the transitional phases, BaWO₄ and Y₁₀W₈O₂₁, at the interface between PrBa(Co_0.7Fe_0.3)₂O_5−δ (PBCF) and Y₂W₃O₁₂ contributes to the development of TEC gradient electrode (Fig. 10(j)). And the transition phase captures the A-site element Ba in the electrolyte layer, thereby optimizing the electrolyte–electrode interface.

The abrupt transition from a dense electrolyte to a porous electrode creates discontinuous reaction pathways and oxygen vacancies, impeding charge transfer and the transport of gas and ionic phases, leading to a high polarization resistance at electrolyte/electrode interface. This weak adhesion, structural and chemical discontinuities, and limited contact points, exacerbates stagnation at the interface and restricts effectiveness of electrode at low operating temperatures.^{84,132,405,406} Therefore, interface engineering is essential to address the challenges of efficient operation in P-SOCs at intermediate to low temperatures by ensuring a seamless transition between the electrolyte and electrode. Choi et al. introduced a mono-grain PBSCF/BZCYYb composite interlayer, which enabled an even lower operating temperature of 350 °C.^236,407 This composite interlayer was fabricated using pulsed laser deposition (PLD), with triple-conducting oxide PBSCF and proton-conducting oxide BZCYYb co-deposited onto a dense BZCYYb electrolyte (Fig. 10(i)). The resulting mono-grain composite interlayer features a quasi-2D thin layer with an alternately arranged, phase-separated structure, which extends the electrode/electrolyte interface by approximately 13.7 times compared to a reference sample with a conventional interface between a porous electrode and a dense electrolyte (Fig. 10(k)). The composite also exhibits a higher concentration, about twice that of the PBSCF bulk powder electrode, due to charge modulation reactions at the heterointerface. This unique structure facilitates reaction pathways through vertically aligned oxygen vacancies, significantly reducing both ohmic and polarization resistances. This mono-grain composite interlayer in a fuel electrode-supported configuration achieved a PPD of 0.50 W cm⁻² and a current density of 0.25 A cm⁻² at 1.3 V at 350 °C.

4.2. LnBaFe₂O_5+δ-based perovskite

Despite their advantageous properties for electrochemical water splitting and power generation, cobalt-based perovskites exhibit drawbacks such as cobalt toxicity, high cost, and large TEC. Recent studies have highlighted several Fe-based double perovskites that show promising performance,^408–410 including LaBaFe₂O_5+δ,⁴¹¹ PrBaFe₂O_5+δ,⁴¹² GdBaFe₂O_5+δ,⁴¹³ NdBaFe₂O_5+δ,⁴¹⁴ and SmBaFe₂O_5+δ.⁴¹¹ Rational substitution strategies have been employed to enhance oxygen vacancy concentrations in LnBaFe₂O_5+δ, for example, doping the A-site with Sr⁴¹⁵ and Ca,⁴¹⁶ and substituting the B-site with W,⁴¹⁷ Ta,⁴¹⁸ Nb,⁴¹⁹ and Mo.⁴²⁰ Pr_0.8Nd_0.2BaFe_1.9Zn_0.1O_5+δ (PNBFZ)²³⁸ developed by Teketel et al., demonstrates improved electrical conductivities, lower TEC (13.9 × 10⁻⁶ K⁻¹), and higher PPD of 401 mW cm⁻² at 550 °C. BCO have been incorporated to construct composites.⁴²¹ However, compared to cobalt-based materials, layered iron-based perovskite air electrode materials exhibit slightly lower performance, and their potential requires further investigation.

5. Other representative perovskite

In addition to barium-containing perovskite materials, various other perovskite materials have seen significant advancements in recent years. This chapter focuses on lanthanide-rich perovskite structures and strontium-based perovskite materials, inspired by O-SOCs, to highlight their recent developments in intermediate-temperature P-SOCs air electrodes (summarized in Table 1).

5.1. Ln₂NiO_4+δ series

The nickel-based perovskite represented by RP-type oxide Ln₂NiO_4+δ (where Ln = La, Pr, and Nd) possesses a unique K₂NiF₄ structure, characterized by the alternating arrangement of perovskite and Ln–O rock salt layers along the c-axis. The presence of excess interstitial oxygen in the Ln–O layer serves as an ionic carrier, resulting additional electron holes,⁴²² which contribute to enhanced electrocatalytic capacity and enhanced surface oxygen-exchange kinetics.^423–425 Grimaud et al. were pioneers in investigating the hydration and electrochemical behaviors of La₂NiO_4+δ (LNO), Pr_2−xSr_xNiO_4+δ, and other related compounds.⁶³ Their work established a direct connection between the layered perovskite structure and triple-conducting behavior, making the RP phase Ln₂NiO_4+δ promising candidates for P-SOCs.^426–428

La³⁺ in the A site with the larger ionic radius (La³⁺: 1.36 Å; Ca²⁺: 1.00 Å; Pr³⁺: 0.99 Å) can effectively lower B–O bond strength for easier O²⁻ migration in the crystal lattice, which is beneficial for ensuring sufficient O²⁻ conductivity at reduced temperatures.^429–432 Unlike typical perovskite structures, where protons reside and migrate on regular O²⁻, layered perovskite (RP), exampled by LNO, may contain oxygen interstitials that allow protons to locate and migrate. However, the performance of LNO-based electrodes is limited by their low electronic conductivity and excessive alkaline earth metals. To improve this, various doping and compositing strategies have been employed.^433,434 Sr-substituted La_1.2Sr_0.8NiO_4+δ (LSN), for instance, achieved resistance of 0.15 Ω cm² in symmetric cell and a PPD of 460 mW cm⁻² at 700 °C.⁴³⁵ Zhong et al. reported that the water content in both pristine LNO and La_2−xA_xNiO_4+δ (where A = Ca, Sr, Ba) rises with the level and basicity of the dopant increase (Fig. 11(a) and (b)).^427,436 Wang et al. improved electronic conductivity by partially substituted Ni with Cu in La_1.5Sr_0.5NiO_4+δ, increasing it from 74 S cm⁻¹ to 171 S cm⁻¹ at 800 °C. This enhancement was primarily attributed to the contraction of Ni–O bonds and the oxidation of Ni²⁺ to Ni³⁺. Cu, Zn co-doped LSN demonstrates a high electrolysis current density of 1.5 A cm⁻² (1.3 V) at 600 °C under high humidity (60%) in air. A self-assembled composite of La_0.6Sr_0.4FeO_3−δ–LSN, developed by Yang et al., exhibited low polarization resistances of 0.055 Ω cm² at 700 °C in O₂.⁴³⁷

Fig. 11 Recent advances in Ln₂NiO_4+δ, SrMO_3−δ, and PrMO_3−δ-type air electrode materials. Proton concentrations per formula unit (pfu) of (a) Ln₂NiO_4+δ, and (b) La₂NiO_4+δ and La_1.6A_0.4Ni_4+δ (A = Ca, Sr, Ba). Reproduced with permission.⁴²⁷ Copyright 2022, Wiley-VCH. (c) Schematic of H₃O⁺ diffusion in NSTF. Reproduced with permission.¹⁶⁷ Copyright 2024, Wiley-VCH. (d) Schematic fabrication process. (e) Peeling strength of cathode-electrolyte interface. (f) AFM for electrolyte surfaces. Reproduced with permission.²⁴¹ Copyright 2022, Springer Nature.

The Pr₂NiO_4+δ (PrN)-based phases have been attempted in P-SOCs systems. Grimaud et al. found that PrN can function as a triple-conducting conducting oxide with low R_p when combined with a proton conductor like BaCe_0.9Y_0.1O_3−δ (BCY10).⁶³ In their subsequent work, they developed the BCY10-PrN composite to reduce TEC and optimize electrochemical properties through microstructure tuning.⁴³⁸ To improve insufficient adhesion of mixed systems, Tarutin et al. developed Cu-doped Pr₂Ni_0.8Cu_0.2O_4+δ, which demonstrates a PPD of ∼340 mW cm⁻² at 750 °C.²³⁹ Nevertheless, the phase stability of PrN requires further improvement due to the decomposition of PrN into PrNiO₃ and PrO₂, leading to mixed ionic oxides formation at the electrode/electrolyte interface.⁴³⁹ The electrocatalytic,⁴²⁴ oxygen transport,⁴⁴⁰ and electrochemical properties^441,442 of Nd₂NiO_4+δ (NNO) have been extensively documented in the literature.⁴⁴³ Luebber et al. utilized NNO to fabricate micro-tubular SOFC. Ba-doped Nd_1.9Ba_0.1NiO_4+δ demonstrates lower polarization resistances, reaching as low as 1.7 Ω cm² at 700 °C.⁴⁴⁴ Additionally, the incorporation of Sr in Nd_2−xSr_xNiO_4+δ alters the oxidation state of Ni from Ni²⁺ to Ni³⁺, thereby influencing its magnetic properties.⁴⁴⁵ However, NNO undergoes a tetragonal to orthorhombic phase transition at 610 °C, which necessitates careful consideration of its application in intermedium-temperature catalysis.⁴⁴⁶

5.2. SrMO_3−δ-based perovskite

SrMO_3−δ with alkaline earth element Sr at A site is used as catalytic activity enhancement materials. SrCoO_3−δ (SCO), for example, was infiltrated on Na_0.3Sr_0.7Ti_0.1Fe_0.9O_3−δ (NSCF0.3@SC) thus achieving an improved PPD of 966 mW cm⁻² at 600 °C.¹⁶⁷ Zhu et al. decorated SrCo_0.5Nb_0.5O_3−δ (SCN) NPs on the surface of PrSrCo_1.8Nb_0.2O_6−δ (PSCN).²⁴⁰ This SCN NPs decoration significantly reduces the energy barriers associated with surface oxygen and vapor dissociation. Thus, SCN-PSCN composite electrode manifests a PPD of 1.30 W cm⁻² and a current density of 1.91 A cm⁻² at 1.3 V at 650 °C. However, Sr-containing electrodes often degrade due to Sr segregation, which results from the reactivity of Sr with steam. This segregation problem is particularly pronounced in EC mode, where higher steam concentrations exacerbate the issue.^194,447 To mitigate it, various methods have been proposed, including doping with Ca,⁴⁴⁸ however, these have shown limited effectiveness. A more promising approach is transforming unstable Sr-containing single phases into stable hybrid materials by leveraging the segregation phenomenon through careful structure regulation.^449–451 Song et al. synthesized a Sr_0.9Ce_0.1Fe_0.8Ni_0.2O_3−δ (SCFN)-based nanocomposite electrode, featuring tetragonal and RP phases with surface-enriched CeO₂ and NiO NPs.¹⁷¹ The RP phase in SCFN enhances hydration and H⁺ conduction, while the nanoscale NiO and CeO₂ phases facilitate oxygen surface exchange and transfer of O²⁻ ions from the RP or NiO surfaces to the primary phase. Zhou et al. developed a self-assembling two-phase Na-doped SrTi_0.1Fe_0.9O_3−δ cathode for P-SOFCs, featuring a Na_ySr_zTi_uFe_1−uO_3−δ (NSTF) perovskite main phase and nanosized β-NaFeO₂ (NF) through thermal-induced phase exsolution (Fig. 11(c)).¹⁶⁷ Unlike most air electrode materials that exhibit a negative dependence on pH₂O, the NSTF/NF nanocomposite cathode demonstrates a positive pH₂O dependence in SOFC applications. This is attributed to enhanced H⁺ uptake and the potential development of a H₃O⁺ transport pathway at the NSTF and NF phases interface, resulting in a new quadruple-conducting (H₃O⁺/H⁺/O²⁻/e⁻) cathode that yields a PPD of 0.807 W cm⁻² at 600 °C.

Extrinsic Cr and S impurities at the surface represent another source of degradation in strontium-based perovskites.⁴⁵² The accumulation of electronically and ionically insulating contaminants—such as Cr₂O₃, SO₂, and secondary compounds like SrCrO₄ and SrSO₄—on the perovskite surface during SOFCs operation, primarily from vaporized Cr originated from metal interconnects and S-based species in the gas flow, ultimately degrades catalytic performance.^453–458 Several strategies have been explored to mitigate the detrimental effects of Cr, including: (i) the use of Cr getters,⁴⁵⁹ (ii) coating of the interconnects,⁴⁶⁰ and (iii) surface modifications that demonstrate high resistance to Cr poisoning.^383,461,462 However, the issue of Cr and S poisoning of Sr-based materials in P-SOCs requires further investigation.

5.3. PrMO_3−δ-based perovskite

PrNi_0.5Co_0.5O_3−δ (PNC), belonging to the PrMO₃ family, is recognized as an effective triple-conducting air electrode material for P-SOCs due to its high proton conductivity. PNC is inspired by PrCoO₃ (PCO) and lanthanide nickelates, combined the advantages of segregation mitigation and excellent tolerance to high steam vapor and varying oxygen partial pressures.^463–466 Ding et al. demonstrated that appropriate nickel substitution at the B-sites of the PNC perovskite can lower the migration barrier for proton conduction, and easily introduce proton defects through hydration reactions.¹⁷² This enhancement in proton conduction, combined with the triple-conducting capabilities of PNC, facilitates both WOR and ORR, thereby improving the electrochemical performance during self-sustained and reversible operations at reduced temperatures. To broaden the application of PNC in low-temperature P-SOCs, addressing the issue of poor contact at the electrode/electrolyte interface is crucial. A straightforward acid treatment developed by Bian et al. can effectively rejuvenate the surface of high-temperature annealed electrolytes (Fig. 11(d)).²⁴¹ This process fosters reactive bonding between the air electrode PNC and the electrolyte BZCYYb, leading to enhanced electrochemical performance and stability (Fig. 11(e) and (f)). Exceptional performance in P-SOFC is achieved at temperatures as low as 350 °C, with PPDs reaching 1.6 W cm⁻² at 600 °C, 650 mW cm⁻² at 450 °C, and 300 mW cm⁻² at 350 °C. However, Pr is classified as a critical material by the European Commission and its large-scale use both unsustainable and costly.⁴⁶⁷ Additionally, When PNC is combined with the proton-conducting electrolyte BZCY, it demonstrates low tolerance to H₂O and CO₂.⁴⁶⁸ Therefore, more research is needed to improve their compatibility. Yao et al., for example, impregnated Pr₂Ni_0.5Co_0.5O_4−δ, which consists of a PNC perovskite phase and a PrO₂ phase, onto LSCF base.²⁴² The PNC perovskite phase facilitates rapid proton transport, while the PrO₂ nanoparticles, characterized by a high concentration of oxygen vacancies, enhance both oxygen adsorption and transfer. The modified electrode achieving a PPD of 1857 mW cm⁻² at 700 °C, and exhibit durability for 200 h.

In addition to alkaline earth and rare earth element-based air electrodes, alkali metal-based materials are also being investigated for solid oxide cell development. For example, LiCoO₂, a widely used cathode material in lithium-ion batteries, has achieved a PPD of 0.86 W cm⁻² at 800 °C in O-SOFCs, benefiting from coordinated lithium volatilization and anion doping.⁴⁶⁹ Given the extensive adoption of lithium-ion technology, integrating recycled lithium-ion battery materials into solid oxide cell research presents a cost-effective approach. Moreover, the high-temperature operating environment efficiently removes volatile impurities, enhancing the feasibility of repurposing recycled materials for solid oxide cell applications.

6. Theoretical insights and computational design

Theoretical calculations are a powerful research tool in the field of SOCs including synthesis of perovskite and their application as air electrode materials. The complex compositions of perovskites render the exploration and development of efficient air electrode materials both time-consuming and challenging, which inevitably drove a search for more efficient tools. Theoretical calculations—especially high-throughput screening and machine learning—inherently became leading tools for the role thanks to their reliability and adaptability. The operational temperature of approximately 600 °C complicates direct observation of the air electrode's functioning SOCs which reduce their reliability there and beyond. Theoretical calculations effectively simulate operating conditions of the electrode and elucidate the underlying microscopic principles at play (summarized in Fig. 12).

Fig. 12 Application of theoretical calculations in the development of perovskite air electrodes.

6.1. Computational approaches for high-performance design

Traditional approaches of discovering and developing new materials typically depend on empirical experiments and trial-and-error techniques, which can be time-consuming, costly, and often constrained in their scope. In contrast, simulations can integrate computational chemistry, molecular orbital theory, and experimental methods to effectively guide the discovery and optimization of air electrodes.^117,470

By leveraging theoretical calculations, researchers can systematically explore material properties and behaviors, thereby streamlining the design process and identifying optimal candidates more efficiently. In the study conducted by Mao et al., the optimization of bimetal-doped SCO was guided by DFT calculations focusing on two key parameters: the formation energy and the migration barrier of oxygen vacancies.²⁶⁷ A lot of 13 different TMs was evaluated as potential dopants with working resulting in 91 possible combinations for bimetal-doped SCO to identify the most promising electrode candidate. Ciucci and colleagues have utilized ab initio simulations, molecular orbital insights to design A- and B-site co-substituted BFO materials.³⁴ By combining calculated lattice constants, substitutional defect formation energy E_form, oxygen vacancy formation energy and hydration energy E_hydr, they concluded a potential Ba_0.875Fe_0.875Zr_0.125O_3−δ candidate (Fig. 13(a)).³⁴ And as validated in experiments, the designed cathode can achieve impressive PPDs of 0.67 W cm⁻², 1.28 W cm⁻², and 2.04 W cm⁻² at 500 °C, 600 °C, and 700 °C, respectively, and outstanding stability exceeding 200 h even in high-steam environments. Simulation-driven strategies enhance the potential for innovative material development while reducing time, manpower, and economic expenditures. Researchers can streamline the experimental process, focus resources on the most promising candidates of high effective catalyst, and avoid redundant or less effective experiments. The high-throughput computational methods employed by Luo et al. facilitated the screening of 932 materials based on key parameters such as formation energy, hydration energy, and adsorption energy, thereby expediting the discovery and optimization of proton conductors.⁴⁷¹ Very recently, Hu et al. used high-throughput calculations combined with data-driven decomposition analysis to forecast essential properties for 4455 unique perovskite oxides, focusing on their thermodynamic stability and decomposition tendencies crucial for various applications (Fig. 13(b)).²³⁷

Fig. 13 (a) Plot of computed E_form against E_vac and E_hydr for BFO and its derivative materials (dots). Reproduced with permission.³⁴ Copyright 2022, Springer Nature. (b) Promising perovskite candidates with exceptional thermodynamic stability. Reproduced with permission.²³⁷ Copyright 2024, Royal Society of Chemistry. (c) Workflow of oxygen vacancies prediction, and (d) and (e) the measured and predicted ASR. Reproduced with permission.⁸¹ Copyright 2024, Springer Nature.

Since the advent of artificial intelligence (AI) and the increasing amount of experimental data in material synthesis, machine learning (ML) techniques have become the corner stone of new material discoveries, accelerating the design, prediction of property, and optimization.³⁷² ML involves the use of computational algorithms and statistical models to analyze extensive datasets, enabling the extraction of meaningful patterns and facilitating predictions or optimizations based on these insights.^472–474 Consequently, it is able to speed up the discoveries and guides experimental work towards the most promising candidates materials at relatively short time and low budge.^475–479 ML models to date have been used to identify the critical factors affecting H⁺/O²⁻/e⁻ conductivities through feature selection and engineering techniques. In 2021, an XGBoost model was constructed to efficiently and accurately predict the conductivities of 97 625 perovskite oxides under wet (3% H₂O) hydrogen and wet air environments.⁴⁸⁰ In hydration abilities analysis, Wang et al. employed a random forest (RF) model to identify efficient mixed proton–electron conductive oxides by predicting the hydrated proton concentration (HPC) of 3200 A_1−xA1_xB_1−yB1_yO₃ perovskite oxides.⁴⁸¹ Through this RF approach, they successfully predicted the HPC of these perovskite oxides, enabling the selection of highly efficient mixed proton–electron conductive materials. Recent studies also indicate that ML models can effectively predict the catalytic activity of air electrode materials in solid oxide cells, particularly regarding resistance to oxygen reduction (R_p), power density, energy barriers, oxygen p-band centers, and vacancy formation energies, often calculated via DFT.^482–485 Integrating ML with DFT has been found enhancing efficiency and accuracy of these calculations, facilitating faster material innovation.⁴⁸⁶ Ni et al. introduced ionic Lewis acid strength (ISA) as an effective physical descriptor to accelerate the research of ORR activity.³⁷² Their model successfully predicted the R_p of 6871 oxides, demonstrating the best fit among eight different regression models.

The important score assigned to each descriptor revealed that A/B-site ionic Lewis acid (AISA/BISA) significantly influences model predictions, suggesting a strong correlation with intrinsic ORR activity. An active learning strategy can be used to predict H₂ production in high-entropy oxides, enabling a more rational design of high-entropy materials rather than relying solely on experience and intuition.⁴⁸⁷ A very recent study by Li et al. shows ML techniques have been used by to quantify the relationship between oxygen vacancy concentrations and catalytic activity, based on 235 cobalt-based and 200 iron-based perovskite catalysts (Fig. 13(c)).⁸¹ The temperature-dependent variation in oxygen vacancy concentrations in both Co-based and Fe-based perovskites can be accurately predicted from well-established elemental properties. The oxygen electrocatalytic activity of perovskite oxides exhibit a “volcano” dependence on oxygen vacancies concentrations, where the optimal oxygen vacancy concentration for electrocatalysis changes with the operating temperatures (Fig. 13(d) and (e)). This study offers an efficient approach with oxygen vacancy-activity relations to efficiently predict perovskite for electrocatalysis operated across a broad temperature range.

Additionally, ML models have been developed to predict the thermodynamic phase stability of perovskite oxides by estimating their energy above the convex hull (E_hull) under realistic conditions.⁴⁸⁸ The reliability of these models was confirmed through comparisons with DFT results and validation against newly generated compounds.⁴⁸⁹ Furthermore, EIS data has been used to train neural networks, which successfully forecast long-term performance degradation. The challenges, however remain as shown by Hagen et al., who found weak correlations in degradation assessments based on a comprehensive dataset of SOC tests, highlighting the need for more extensive data to enhance statistical significance in ML applications for durability predictions.⁴⁹⁰ ML holds great promise for developing high-efficiency perovskite air electrode materials for P-SOCs. However, challenges such as data availability, complexity of materials data, and the need for effective feature engineering hinder its application. Balancing multiple target properties like catalytic activity and stability is also difficult due to trade-offs between performance and degradation. Despite these obstacles, leveraging big data and advanced computing can enhance efficiency and accuracy of ML in materials research. Overcoming these challenges requires collaboration, better data standards, and improved model interpretability.

6.2. Mechanistic understanding from simulations

Due to the challenges of direct observation, theoretical calculations have emerged as a powerful tool for elucidating the microscopic mechanisms of efficient air electrodes. These mechanisms primarily involve reactions related to protons and oxygen ions. By analyzing the adsorption, transfer, and reaction processes of charge carriers, researchers can establish a connection between structure and performance, thereby guiding the optimization of air electrode materials. Hydration reactions involve proton incorporation and interactions with lattice oxygen, which is regarded as water adsorption and proton transport. Li et al. evaluated the impact of doped boron (B) on surface chemical properties by calculating water adsorption energies.⁴⁹¹ They found that the B-modified Pr₄Ni₃O_10+δ (0.5B-PN) surface had a water adsorption energy of −1.4 eV, compared to −0.9 eV for the PN surface, reflecting a 55.6% improvement in hydration ability due to the surface modification with dispersed B. Lu et al. calculated hydration energies for different cells by inserting H₂O into oxygen vacancies of extracted cells in parent perovskite structures (Fig. 14(a)).¹⁸³ BaCo_0.4Fe_0.4Nb_0.2O_3−δ (BCFN) showed the highest hydration energy which is indicative of strong affinity for H₂O molecules, the performance partly attributed to Nb⁵⁺ doping. For F-doped BCFZY examined by Ren et al., the reduction in lattice size was confirmed through cell structure analysis, and hydration properties were illustrated by the adsorption energies of H₂O on Fe–V_O–Co sites and integrated crystal-orbital Hamilton populations (iCOHPs, Fig. 14(b)).¹⁶⁶ By simulating the electronic structure and potential energy surfaces of perovskites, proton conduction mechanisms can be provided. Liu et al. used climbing image nudged elastic band (CINEB) calculations to show higher energy barrier (E_b) for proton migration for hopping pathway than rotation pathway in perovskites, identifying hopping as the rate-limiting step (Fig. 14(c)).¹⁷⁹ They found Er-doping in BSCF lowering E_b for the hopping pathway by 0.13 eV. Er₂O₃, with its simpler proton hopping process, achieves an even lower E_b of 0.35 eV and a greater hopping distance compared to Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF) and Ba_0.5Sr_0.5(Co_0.8Fe_0.2)_0.9Er_0.1O_3−δ (BSCFE), indicating that Er₂O₃ markedly improves proton migration in composite electrodes. Beyond E_b, factors such as the O⋯H bonds formation energy and distance are also crucial for predicting proton conduction. Chen et al. reported higher positive formation energy for O⋯H bonds in Ba_0.5Sr_0.5Co_0.8Fe_0.2O_2.9−δF_0.1 (BSCFF) than Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (BSCF), suggesting more robust bonds formation with oxygen by the proton in BSCF (Fig. 14(d)).²¹³ This stronger bonding results in a more challenging hopping process for protons, following the Grotthus mechanism. Moreover, fluorinated oxides like BSCFF exhibit longer O⋯H distances, which implies easier disruption of these bonds and a reduced energy barrier for proton displacement. To analyze proton diffusion pathways and conductivity,⁴⁹² first-principles molecular dynamics (FPMD) simulations at 1000 K were employed by Kim et al.¹⁹⁹ These simulations showed more rapidly protons diffusion in BaSc_0.1Ta_0.1Co_0.8O_3−δ (BSTC) compared to BaCoO_3−δ (BCO) as evidenced by a higher mean squared displacement (MSD) for BSTC (Fig. 14(e)). This enhanced diffusion is linked to hydroxide rotation, proton transfer, and trapping phenomena. In terms of diffusion pathways, BSTC establishes a continuous 3D proton transport network within 100 ps, whereas BCO develops a more limited 2D network (Fig. 14(f)). The proton conductivity of BSTC measured at 4.94 × 10⁻⁴ S cm⁻¹ using the Nernst–Einstein relation, indicating its suitability for P-SOCs air electrodes. Furthermore, negative crystal orbital Hamiltonian populations (–COHPs) reveal that O⋯H bonds in BSTC are weaker than those in BCO with integrated –COHP values of 3.87 eV bond⁻¹ for BSTC and 3.95 eV bond⁻¹ for BCO. This suggests that BSTC facilitates more efficient proton transfer due to its weaker O⋯H bonds and decomposition in the O 2p and H 1s orbitals.

Fig. 14 Recent advances in theoretical explanation for efficient air electrode materials. (a) Hydration energies of the BCFN, BCFNS, and BCFS (inset: optimized structural models for the hydrated samples of BCFN). Reproduced with permission.¹⁶⁶ Copyright 2023, Elsevier. (b) M–O iCOHP value comparison for BCFZY(F). Reproduced with permission.¹⁶⁴ Copyright 2022, Elsevier. (c) Energy profile for proton hopping and rotation in BSCF, BSCFE, and Er₂O₃. Reproduced with permission.¹⁷⁹ Copyright 2023, Wiley-VCH. (d) O⋯H formation energy and distance for BSCF(F). Reproduced with permission.²¹³ Copyright 2023, Elsevier. (e) MSD of a hydrogen atom and (f) time-dependent evolution of H⁺ diffusion pathways in BSTC and BCO. Reproduced with permission.¹⁹⁹ Copyright 2024, Wiley-VCH. (g) PDOS of the O 2p and Co 3d orbitals of PBCC and PBSLCC. Reproduced with permission.⁴⁹³ Copyright 2023, Elsevier. (h) Oxygen vacancy formation energy in electrodes SBCF-2L and SBCF-5L. Reproduced with permission.⁴⁹⁴ Copyright 2024, Wiley-VCH. (i) Energy barrier of cubic C-BSCF (100) and hexagonal H-BSCF (0001) surfaces. Reproduced with permission.²¹⁴ Copyright 2024, Springer Nature.

Proton-involved oxygen reactions dominate the kinetics of air electrode. He et al. utilized DFT calculations to investigate how A-site entropy tuning affects the performance of perovskite materials for OER and ORR.^212,493 The study compared high-entropy Pr_0.2Ba_0.2Sr_0.2La_0.2Ca_0.2CoO_3−δ (PBSLCC) with binary PrBaCo₂O_5+δ (PBC) and ternary Pr_0.8Ba_0.8Ca_0.4Co₂O_5+δ (PBCC), focusing on key descriptors such as the O p-band center,^36,107,495 formation energies, and oxygen adsorption. The projected density of states (PDOS) of PBCC and PBSLCC was displayed to verify the strong hybridization of cations in the B-site Co ion with oxygen (Fig. 14(g)).⁴⁹³ The O p-band center PBSLCC shifted closer to the Fermi level (−3.29 eV) compared to PBC and PBCC. The oxygen vacancy formation energies (E_vac) in PBSLCC (0.97 eV) is lower compared to PBC (1.35 eV) and PBCC (1.31 eV).⁴⁹⁶ Oxygen adsorption energies (ΔE_ads)studies revealed that PBSLCC exhibits weaker oxygen bonds and better kinetics, in line with Sabatier's principle.^497,498 Overall, the research confirms that A-site entropy tuning significantly enhances OER and ORR performance by optimizing the activity of oxygen reactions.

Oxygen vacancies play a crucial role in affecting ion transfer and hydration properties of perovskite oxides. To assess the hydration and oxygen ion transfer capabilities of air electrode materials, researchers examine the formation energy of oxygen vacancies in different materials and lattice sites. Liu et al. showed that the Er-doping notably reduces E_vac and hydration energy (E_hdr) by 0.12 eV and 0.22 eV, respectively.¹⁷⁹ This doping strategy thus enhances hydration ability of Ba_0.5Sr_0.5Co_0.72Fe_0.18Er_0.09O_3−δ–Er₂O₃ composite air electrode. In a separate study, Sun et al. used DFT calculations to evaluate the oxygen vacancies formation energy in Sm_0.8Ba_1.2Co_0.8Fe_1.2O_5+δ (SBCF-5L) and SmBaCo_0.8Fe_1.2O_5+δ (SBCF-2L).⁴⁹⁴ he vacancies for SBCF-5L are more easily formed in the –[SmO_1−x]– layer, with formation energies of 0.289 eV and 0.314 eV for the –[SmO_1−x]– and perovskite layers, respectively. In contrast, SBCF-2L shows values of 0.217 eV and 0.291 eV (Fig. 14(h)). When vacancies are present in the –[SmO_1−x]– layer of SBCF-5L, the formation energy in the perovskite layer drops, facilitating enhanced three-dimensional oxygen ion conduction. This results in a higher initial concentration of oxygen vacancies in SBCF-5L compared to SBCF-2L, making SBCF-5L more efficient for oxygen ion transport. In Chen et al.'s work, the oxygen vacancy formation and PDOS of Fe-substituted Nd_0.8Sr_1.2Ni_1−xFe_xO_4±δ was predicted by DFT.⁴⁹⁹ The oxygen vacancy formation energy was calculated as 4.34 eV for Nd_0.8Sr_1.2NiO_4±δ (NSN) and 4.19 eV for Nd_0.8Sr_1.2Ni_0.7Fe_0.3O_4±δ (NSNF), showing that Fe doping reduces this energy compared to other perovskite oxides. Fe 3d electrons are closer to the Fermi level compared to Ni 3d, indicating improved catalytic activity with Fe doping. Lower oxygen vacancy formation energy, coupled with accelerated charge transfer through Fe–O–Ni bridges, suggests enhanced catalytic activity for ORR and OER in Fe-doped NSN.^500,501

In perovskite oxides, electronic conduction often occurs via polaron hopping. To understand factors influencing overall electronic conductivity, density-of-states (DOS) and band alignment simulations are employed. Yu et al. analyzed DOS plots for Sr₃Fe_1.8Nb_0.2O_7−δ (SFN) and Sr_2.8Fe_1.8Nb_0.2O_7−δ (D-SFN) to understand the effects of Nb doping and A-site cation defects.⁵⁰² These plots show that the conduction band near the Fermi level is mainly influenced by Fe 3d orbitals while the valence band comes from O 2p orbitals. In D-SFN, there are more unoccupied states near the valence band which suggests better charge transfer. The addition of Sr defects increases Fe valence states, enhancing the overlap between Fe 3d and O 2p orbitals. This stronger Fe–O interaction improves electron conduction in D-SFN. Additionally, DOS analysis was performed by Zhou et al. to analyze different active sites of PrBa_0.8Ca_0.2Co₂O_5+δ–BaCoO_3−δ (PBCC–BCO) composite air electrode.¹⁹⁴ Their findings showed that Co majority spin states at PBCC (010) lose more electrons compared to BaCoO₃ (110), which is linked to differences in coordination. Specifically, BaCoO₃ (110) benefits from electron-donating Ba, which creates an electron-rich region at the Co sites, enhancing interactions with oxygen-containing intermediates. This increased interaction facilitates water deprotonation but limits oxygen desorption on BaCoO₃ (110).

The OER process on the air electrode surface involves several steps: H₂O adsorption (*H₂O, * indicates the adsorbed state), oxygen evolution (*H₂O → *OH + *H → *O + *H → *O → *O₂), and *O₂ desorption, whereas the ORR process is essentially the reverse of OER process.^30,194,503 Combining calculating adsorption energies, reaction intermediates, and activation barriers, OER and ORR pathways can be simulated at the atomic level. This approach helps identify the rate-determining steps and optimize the catalytic activity of air electrodes. Liu et al. investigated the energy barriers for OER and ORR on cubic Ba_0.5Sr_0.5Co_0.8Fe_0.2O_3−δ (C-BSCF) and hexagonal Ba₄Sr₄(Co_0.8Fe_0.2)₄O_16−δ (H-BSCF) surfaces to understand their mechanisms in P-SOCs air electrodes.²¹⁴ They determined the most stable surface configurations by evaluating various A- and B-site arrangements and surface terminations. The results revealed that C-BSCF has a higher energy barrier of 3.25 eV for OER mainly due to limited proton conduction whereas H-BSCF requires a lower barrier of 1.15 eV, highlighting its superior proton migration capability. Conversely, H-BSCF faces a 2.03 eV barrier for oxygen desorption while C-BSCF performs better in this regard. The hybrid Ba_1.5Sr_1.5Co_1.6Fe_0.4O_7−δ (C/H-BSCF) electrode combines the advantages of both phases, with initial water decomposition occurring on the cubic phase and further decomposition and oxygen desorption on the hexagonal phase. For ORR, C-BSCF shows a more favorable oxygen reduction process with a significant advantage over H-BSCF which has lower energy barriers for oxygen adsorption and desorption (Fig. 14(i)). Overall, the C/H-BSCF hybrid electrode benefits from enhanced ORR and OER activities due to the synergistic effects of combining cubic and hexagonal phases. When examining complex hybrid air electrodes, it is often highly effective to investigate the microscopic mechanisms of oxygen reactions with theoretical calculations. In composite air electrode PBCC–BCO constructed by Zhou et al., the H₂O adsorption and oxygen evolution are favored on BCO with ΔG values of 0.099 eV (H₂O adsorption) and 0.62 eV (oxygen evolution) for BCO, compared to 1.794 eV (H₂O adsorption) and 1.055 eV (oxygen evolution) for PBCC at 600 °C.¹⁹⁴ The PBCC surface, on the other hand, is more favorable for O₂ formation and desorption (1.307 eV for BCO vs. −0.823 eV for PBCC at 600 °C). Combined with DOS, the proposed OER sequence for the PBCC–BCO air electrode involves initial H₂O absorption on the BCO (110) surface, where it dissociated into *OH and *H. The *OH intermediate is then further dissociated into *O and *H on BCO (110), and finally, O₂ is formed and desorbed on PBCC (010). This sequence highlights the combined effects of rapid H₂O dissociation on BCO NPs and fast oxygen desorption on PBCC in enhancing OER activity.

7. Pre-commercialization considerations

Renewable energy sources now contribute more than 29% of global electricity generation, with countries like Australia reaching levels as high as 35%.⁵⁰⁴ Several large-scale hydrogen plants have recently been launched worldwide by corporations including European Energy in Denmark,⁵⁰⁵ TotalEnergies in French,⁵⁰⁶ and German Aerospace Center in Germany.⁵⁰⁷ By integrating electrolysis with renewable energy sources, fuels and chemicals production can move away from fossil fuels dependence, achieving the development of a fully renewable energy system. P-SOCs offer a reliable and affordable solution to incorporate green energy into the current energy infrastructure. This chapter outlines key factors to address before P-SOCs technology can be fully commercialized.

7.1. Scaling up from lab to industry

Lab-scale P-SOCs are generally designed with compact sizes, for instance, O’Hayre, Shao, and Duan et al. report active areas of just around 1 cm² for their P-SOCs.^9,81 This small scale is suitable for laboratory research, where studies focus on optimizing electrode materials, cell structure, and performance parameters under controlled conditions. However, for practical applications, P-SOCs must be significantly larger to meet the energy demands of real-world systems (10² W for electronic devices and 10⁴ W for electronic vehicles). Application-oriented P-SOCs typically adopt a planar or tubular design to increase surface area and enhance efficiency. Currently, numerous studies have explored larger-scale P-SOCs (as shown in Fig. 15(a) and Table 2) including 12 × 12 cm² planar designs and 8 cm² tubular configurations.^161,162 However, despite these increased sizes, the total output power remains insufficient to operate electronic devices. The production capacity of a P-SOC plant is directly proportional to the total active area. Expanding the active area at the individual cell level and assembling multiple cells into stacks are the two most straightforward methods to increase output.^508,509 Scaling up requires addressing challenges such as maintaining mechanical integrity, achieving uniform gas distribution, and managing thermal stresses across a larger area. In particular, the design must ensure durability and stable performance under operating conditions that differ from the controlled lab environment (Fig. 15(b)), making factors like material selection, sealing methods, and thermal management crucial to the development of commercial-scale P-SOCs systems.^510,511

Fig. 15 Large-area P-SOCs: (a) PPDs at 600 °C and (b) durability (data detailed in Table 2).

Table 2 Application-oriented P-SOC air electrodes design

Air electrode	Area (cm^2)	PPD at 600 °C (mW cm^−2)	Current density at 1.3 V (mA cm^−2)	Stability (h)	Ref.
Planar
BaCe0.2Zr0.7Y0.1O3−δ	135			400 h (ASR)	512
BaZr0.1Ce0.7Y0.2O3−δ	144			90 h (650 °C and 1330 mA cm^−2)	162
Ba0.5Sr0.5Co0.8Fe0.2O3−δ–BaCe0.8Zr0.1Y0.1O3−δ (BSCF–BCZY)	20	165		3000 h (600 °C, 170 mA cm^−2 and −100 mA cm^−2)	285
Ba0.5Sr0.5Co0.8Fe0.2O3−δ (BSCF-1)	25	1302		80 h (550 °C and 500 mA cm^−2)	513
BaCo0.4Fe0.4Zr0.1Y0.1O3−δ (BCFZY-1)	81	518			514
BCFZY-1	25	1400		500 h (550 °C and 267 mA cm^−2)	514
BaCo0.4Fe0.4Zr0.1Y0.1O3−δ (BCFZY-2)	5	690		1400 h (550 °C and 100 mA cm^−2)	515
Nd2NiO4+δ (NNO)	2.01	60		1000 h (600 °C at 60 mA cm^−2)	516
Sm0.5Sr0.5CoO3−δ (SSC)	3.33	747			517
Tubular
BaZr0.1Ce0.7Y0.1Yb0.1O3−δ–La0.6Sr0.4Co0.2Fe0.8O3−δ (BZCYYb–LSCF)	2	331			518
La0.6Sr0.4Co0.2Fe0.8O3−δ–Ba0.5Sr0.5Ce0.6Zr0.2Gd0.1Y0.1O3−δ (LSCF–BSCZGY)	1.79	150		96 h (700 °C and 0.7 V)	519
BCFZY-3	8	517	915	1000 h (550 °C and 0.062 A cm^−2)	161
La0.6Sr0.4Co0.2Fe0.8O3−δ–Sm0.2Ce0.8O2−δ (LSCF–Sm0.2Ce0.8O2−δ)	2	580			520
La0.6Sr0.4Co0.2Fe0.8O3−δ–BaCe0.7Zr0.1Y0.2O3−δ (LSCF–BCZY)	2.3	465		70 h (600 °C and 0.7 V)	521
Ba0.5Sr0.5Co0.8Fe0.2O3−δ (BSCF-2)	1.29	534			359
PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (PBSCF)	4.4	776		20 h (650 °C and 0.5 A cm^−2)	522

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7.2. Key challenges for commercialization

As an air electrode material progresses toward practical applications, the criteria for evaluating its performance shift accordingly. Several practical considerations must be addressed, spanning technical, economic, and market-related factors. A primary concern is production scalability, which requires the development of cost-effective synthesis methods and dependable supply chains for essential raw materials such as Ba and Co. The traditional approaches for fabricating perovskite electrode materials include molten salt synthesis, hydrothermal reaction, sol–gel synthesis, and solid-state method. Each method plays a crucial role in determining phase purity, microstructure, defect concentration, and ionic/electronic transport properties, which directly impact electrode performance in P-SOCs. The hydrothermal method, performed under low-temperature (∼150–300 °C) and high-pressure conditions, enables fine morphology control and low-defect-density perovskites, which enhance proton and oxygen ion transport, but scalability remains a challenge due to high-pressure requirements.^523,524 Molten salt synthesis is an emerging technique that allows for the formation of highly crystalline, defect-engineered perovskites at around 800 °C by utilizing molten salts as a reaction medium to promote controlled grain growth and enhanced ionic diffusion, which improves oxygen vacancy mobility and catalytic activity.⁵²⁵ The sol–gel method, operating at a relatively low calcination temperature (∼600–900 °C), facilitates the synthesis of nanostructured perovskites with a high surface area and uniform dopant distribution while preserving porosity and enhancing gas diffusion. These characteristics contribute to improved catalytic activity and ionic conductivity. However, potential issues such as organic solvent impurities, phase separation require careful management, unavailability and high price of the corresponding salts.⁵²⁶ Conversely, the solid-state reaction method, the most cost-effective and scalable approach for large-scale perovskite production, involves high-temperature calcination (∼1000–1400 °C) of mixed solid precursors, producing highly crystalline and thermodynamically stable perovskites with superior bulk electronic conductivity. However, this approach often results in large grain sizes with a low surface area, leading to non-uniform dopant distribution that may necessitate additional grinding or doping strategies for optimization.⁵²⁷ Additionally, advanced electrode structuring techniques such as freeze-casting, sacrificial templating, and 3D printing can help create highly porous, hierarchical structures that enhance both electronic and ionic transport without compromising mechanical integrity. Moreover, roll-to-roll processing and screen-printing techniques offer promising routes for integrating perovskite electrodes into commercial-scale P-SOCs with high reproducibility and cost efficiency.^528,529 By implementing these strategies, the transition from research to commercialization can be accelerated. Achieving competitive costs is crucial, necessitating reductions in manufacturing expenses through process optimization, material efficiency, and economies of scale.

Performance improvement also remains essential, particularly in enhancing outputs, efficiency, and lifespan, to ensure that P-SOCs can meet the rigorous demands of high-performance applications like electric vehicles and energy storage systems. In addition to technical and economic considerations, practical deployment of P-SOCs requires attention to delivery logistics, safety, and recycling strategies.^530–532 Efficient delivery systems ensure that both raw materials and finished products reach their destinations securely and cost-effectively, with protective packaging, optimized shipping routes, and local distribution centers to minimize damage and costs. Safety is also critical, especially given the high operating temperatures and reactivity of materials in P-SOCs; ensuring thermal stability, reliable sealing, and resistance to environmental stress are essential for operational integrity. Finally, recycling strategies are needed to recover valuable materials like Co from spent cells, reducing reliance on raw material extraction and supporting sustainable production. Together, these considerations enable P-SOC technology to be deployed safely, sustainably, and at scale.

8. Challenges and prospects

In the intermediate-to-low temperature range, state-of-the-art P-SOCs continue to exhibit suboptimal performance compared to conventional O-SOCs.⁵³³ Air electrodes for P-SOCs have garnered significant interest due to their high polarization resistance, yet no truly successful air electrode material has emerged, unlike the established LSM and LSCF for O-SOCs. A review of the latest advancements in P-SOC air electrodes shows that the BCO series achieves the lowest ASR (Fig. 16(a) and (b)), with LnBaM₂O_6+δ double perovskite materials following closely (Fig. 16(c)). In contrast, Fe-based materials typically demonstrate less remarkable performance (Fig. 16(d)). To minimize the resistance of P-SOCs device and enhance both power output and hydrogen production efficiency, the BCO series may emerge as a key focus in future research efforts (Fig. 16).

Fig. 16 Comparison of ASR in symmetrical cell configuration of advanced air electrodes for P-SOCs at 3% H₂O: (a) BaCoO_3−δ-based perovskites, (b) BaCoO_3−δ-derivated perovskites, (c) LnBaM₂O_6+δ-based perovskites, and (d) Fe-based perovskites. Summary of electrochemical properties of advanced perovskite air electrode materials (detailed data summarized in Table 1): (e) ASR of lowest working temperature in symmetrical cells configuration, (f) PPDs in fuel cell (FC) mode, and (g) current densities in electrolysis cell (EC) mode.

Lowering operational temperatures is crucial for sustainability, as it reduces costs, enhances market feasibility, and opens the door for portable and transportable applications. Although P-SOCs operate at lower temperatures than O-SOCs—down to 350 °C—their performance remains insufficient for practical applications (Fig. 16(e)–(g)).²³⁶ At low temperatures, the primary challenge is typically the slow kinetics of ORR and OER. Current research often focuses on modifying existing air electrodes to enhance proton conduction, resulting in numerous variations and studies but limited genuine innovation. As operating temperatures continue to decrease, combining insights from low-temperature catalytic systems, such as PEM and anion exchange membranes, could offer valuable inspiration. For instance, low-temperature catalysis often achieves performance optimization through controlled crystal orientation and complex surface interface design. Furthermore, the reduced operating temperature enables direct observation of microevolution within the P-SOCs system. Advanced characterization techniques, including in situ TEM, can provide insights into the underlying mechanisms, thereby aiding and expediting the development of advanced air electrode materials.

Ensuring stable operation is a key aspect of optimizing both power generation and hydrogen production in P-SOCs, directly impacting efficiency, reliability, and system lifespan. Current treatment methods for air electrode materials like segregation, can raise concerns about durability. The failure of P-SOCs during operation may also arise from interface issues stemming from the thermal expansion differences between the electrode and the electrolyte. Improving the electrolyte/electrode interface is essential for enhancing air electrode. Given the complexity of the composition and elements in perovskite air electrode materials, employing AI-driven design can significantly speed up the material development process. Moreover, the P-SOC system is strongly influenced by physical factors like temperature and pressure. Finite element analysis can be applied to assess and forecast the performance of P-SOC devices in real-world operating environments.⁵³⁴

Lab-scale experiments are crucial for developing new technologies as they provide essential data for subsequent pilot and industrial-scale investigations. The reliability of this experimental data is vital as it guides future research directions, especially in the development of SOC technology for electrochemical energy storage and conversion which is key for achieving a net-zero carbon society. However, the complexity of lab-scale testing involving multiple parameters such as cell configuration, performance measurement strategies, and temperature distribution, can obscure true cell performance and affect power output. Discrepancies in results often arise from the omission of important parameter information in reports, leading to significant variances in literature. To facilitate reliable comparisons in future research, it is recommended that detailed information on various parameters including cell size, surface area ratios, sealing methods, and flow rates, be consistently provided.¹⁶³ Additionally, the extra heat produced due to irreducible internal resistance during testing may lead to actual operating temperatures significantly exceeding the recorded values. This discrepancy underscores the need for more precisely defined experimental conditions.

Despite extensive research, commercial implementation of P-SOCs systems still has not materialized because of insufficient power outputs and high cost.³⁴ Various considerations must be addressed during the laboratory stage to effectively align with commercial goals including the costs and availability of raw materials, expanded cells area and electrolyte sealing, safety and lifespan during storage and transportation, as well as the recyclability and environmental impact of waste P-SOCs.^535,536 Consequently, it is advisable, in the early phases of research and development, to choose air electrode materials that are low-cost, safe, and environmentally friendly. This approach ensures that the materials are not only functional but also viable for future commercial applications.

Author contributions

C. P. and X. X. contributed to the central idea and conducted the draft. S. M. collected resources. P. K., Y. Z., and X. X. contributed to the writing and editing of the document.

Data availability

No primary research results, software or code have been included and no new data were generated or analysed as part of this review.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

C. P. gratefully acknowledge financial support provided by University of Adelaide research scholarship. Open Access publishing facilitated by The University of Adelaide – The University of Adelaide agreement via the Council of Australian University Librarians. Y. Z., X. X. acknowledges funding from the Australian Research Council (LP210301397).

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