Recent advances, practical challenges, and perspectives of intermediate temperature solid oxide fuel cell cathodes

Abstract

As a highly efficient clean power generation technology, intermediate temperature (600–800 °C) solid oxide fuel cells (IT-SOFCs) have gained much interest due to their rapid start-up and shut-down capability, longer life-time and lower cost compared to the conventional SOFCs. However, the sluggish oxygen reduction reaction (ORR) at the cathode at lower temperatures, chromium (Cr) poisoning of cathodes when exposed to Cr-based interconnects, material degradation under CO₂ and humid atmospheres, and compatibility of Co-containing cathodes with existing IT-SOFC electrolytes still affect their large-scale development. This work aims to present an overview on the latest achievements in developing IT-SOFC cathodes based on perovskite-type and other crystal structures, and composites. The utilisation of distribution of relaxation times for analysing the impedance spectra of SOFC cathodes has been discussed. Furthermore, this article presents summary towards the rational design of the cathode materials and structures, to realize cost-effective and high-performance IT-SOFCs.

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Amanda Ndubuisi

Amanda Ndubuisi received her M.Sc. degree in Physics from Covenant University, Nigeria, specializing in Renewable Energy and Materials Science. She is a PhD student at the Department of Chemistry, University of Calgary, Canada, supervised by Professor Venkataraman Thangadurai. Her research focuses on developing mixed ionic and electronic catalysts for oxygen reduction in intermediate temperature solid oxide fuel cells and CO2 reduction in solid oxide electrolysis cells.

Sara Abouali

Dr Sara Abouali is a materials scientist in the field of electrochemical energy storage/conversion devices. She obtained her PhD in Mechanical Engineering (Energy Concentration) from the Hong Kong University of Science and Technology (HKUST) in 2016. Since then, she has been working as a researcher/postdoctoral associate in international institutes in Hong Kong, Italy and Canada. She has published several scientific papers and patents on designing advanced electrodes and electrolytes for batteries, supercapacitors and fuel cells.

Kalpana Singh

Dr Singh received her PhD in 2016 from the University of Calgary (UofC), Canada, with specialisation in Physical Chemistry. She was a graduate student in the laboratory of Venkataraman Thangadurai at Calgary. Dr Singh's research interests include fundamentals of solid state electrolytes and electrodes for solid oxide fuel cells. During her PhD research at Calgary, she extended her research direction towards mixed conducting metal oxides for solid oxide electrolyser cells. Dr Singh has authored 18 peer-reviewed publications, 3 conference proceedings, and 1 patent application as a result of goal driven PhD and post-doctoral research, independent research projects, and collaborations with research institutes in USA, China, Australia, and Israel.

Venkataraman Thangadurai

Dr Venkataraman Thangadurai is a full professor of chemistry, University of Calgary. He received his PhD in 1999 from Indian Institute of Science, Bangalore, India. He conducted his postdoctoral research at the University of Kiel, Germany, with a prestigious fellowship from the Alexander von Humboldt Foundation, Bonn. In 2004, he received the Habilitation degree from the University of Kiel. He is a Fellow of the Royal Society of Chemistry, United Kingdom, Fellow of the International Association of Advanced Materials, Sweden, and Fellow of The Electrochemical Society. He received the prestigious Keith Laidler Award and Research Excellence in Materials Chemistry Award from the Canadian Society of Chemistry. His research focuses on the design, synthesis and characterization of novel solid-state electrolytes and elemental and mixed conducting electrodes for next generation solid-state batteries, solid oxide fuel cells, electrolysis cells, catalysis, and gas sensors. He has 220+ publications (h-index 58), which have been cited over 15 500 times.

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Introduction

Solid oxide fuel cells are electrochemical energy conversion devices attractive for clean generation of electricity from a variety of fuels with high efficiency. The chemical energy of reactants in SOFCs is directly converted into electrical energy through the oxidation of the fuel without any intermediate step of combustion. Thus, unlike combustion engines, this technology is not limited by the Carnot cycle efficiency. Due to their high efficiency and energy densities, SOFCs have the potential for large-scale stationary power generation. Current SOFCs are operating at high temperatures (800–1000 °C) which eliminates the need for expensive noble metal catalysts by increasing the electrochemical kinetics of reactions. Waste heat generated from high temperature SOFCs (HT-SOFCs) can be fed into combined heat and power (CHP) systems to increase the total efficiency of the system. Integrating SOFC units in central heating systems of households will not only provide electricity but heat and hot water with natural gas as fuel. Furthermore, employing the SOFC technology will lower carbon footprint as it generates high quality CO₂ which can be sequestered and used as fuel in reverse mode SOFCs (electrolyser cells) to yield syngas, suitable for zero carbon economy.

Despite the advantages of HT-SOFCs, operating at high temperature presents significant drawbacks which hinder their full implementation in energy systems. These challenges include high cost of interconnects and sealants, accelerated degradation of the components, and subsequent degeneration in the performance of the cell as a result of elevated working temperatures. The state-of-the-art materials for SOFCs consist of Ni-yttria-stabilized zirconia (YSZ) composite anodes, oxide ion conducting YSZ electrolyte, and La_1−xSr_xMnO_3−δ (LSM) cathode. Ni–YSZ composite anode is susceptible to redox cycling instability, and H₂S and coke poisoning.^1,2 LSM is typically mixed with YSZ electrolyte in a composite to extend the triple phase boundary which is an active site for oxygen reduction and increases the ionic conductivity.

Lowering the operating temperature of SOFCs to the intermediate temperature (IT) range (600–800 °C) has been reported to mitigate the challenges associated with HT-SOFCs thus offering technical and economic advantages.^3,4 On the other hand, decreasing the operating temperature to intermediate levels, will generate some other challenges towards the oxygen reduction reaction (ORR) activity of the cathode. The reduction of oxygen at the cathode is a thermally activated process and the kinetics of the reaction is decelerated, leading to significant electrical losses and a drop in the electrochemical performance of the cell at lower temperatures. Therefore, designing an advanced cathode with a high catalytic activity is essential to enhance the electrochemical performance of the cell at intermediate temperatures. Several cathode designs with novel compositions and engineered micro/nanostructures have been proposed. Also, advanced techniques have been used to further shed light on the electrochemical properties of the cathode and anode. One of these relatively new approaches is using the distribution of relaxation times (DRT) to analyse the impedance spectra to measure the polarization resistance of the anode and cathode.⁵ Herein, we present an overview on advances in the development of IT-SOFC cathodes, mainly focusing on the development of novel cathode compositions and structures based on single perovskites, double perovskites, Ruddlesden–Popper layered perovskite-type oxides, swedenborgite-type, garnet-type and composite cathodes. The challenges associated with each group have been discussed and some of the proposed solutions have been reviewed. Afterwards, a brief discussion on the ORR mechanism has been presented and the application of DRT for analysing the degradation mechanisms of IT-SOFC cathodes is discussed.

Perovskite-type structure SOFC cathodes

As an important family of oxides, perovskite-type structures have shown great potential to be used as electrocatalysts for the ORR in an SOFC cathode. A variety of perovskite compositions have been developed with unique characteristics while specific concerns have been identified for each group. Overview on the structure, synthesis, compositions and electrocatalytic behaviour of perovskites along with the latest progress and strategies on mitigating the practical challenges of this class of complex oxides when used as IT-SOFC cathodes is presented.

Crystal structure

The first perovskite oxide was discovered by Gustav Rose in 1839 with a chemical composition of CaTiO₃ and was named after Lev Aleksevich von Perovski, a Russian mineralogist.^6,7 The general formula of a perovskite can be written as ABO₃, where A and B are cations with a coordination number of 12 and 6, respectively.⁸ Alkaline, alkaline-earth and rare-earth metals can occupy larger A-sites while transition metals with smaller size can occupy B-sites.^8,9Fig. 1a schematically shows an ideal cubic perovskite structure, however, due to the variation of ionic radius after cation substitutions, structural distortion changes the cubic lattice to tetragonal, orthorhombic, and rhombohedral, as demonstrated in Fig. 1b–d.^10–17 The deviation of the lattice from an ideal cubic structure can be calculated using the tolerance factor (t) as:where r_A and r_B are the ionic radii of A and B cations, respectively, and r_O is the ionic radii of oxygen. The perfect cubic symmetry is obtained when t = 1, while t < 1 and t > 1 lead to the formation of less symmetric crystals with rhombohedral/orthorhombic (Fig. 1b and c) and hexagonal structures (Fig. 1d), respectively. It should be noted that the physico-chemical properties of perovskites are highly governed by their structure which in turn can be tuned by partial or complete substitution of A and/or B cations.⁹

Fig. 1 Schematic illustration of a perovskite-type (ABO₃) crystal with (a) cubic, (b) rhombohedral. Reprinted with permission from ref. 11. Copyright 2017 American Chemical Society. (c) Orthorhombic¹² and (d) hexagonal structures. Reprinted from ref. 13. Copyright 2015, with permission from Elsevier.

Synthesis methods of perovskite structures

Table 1 summarizes the most common techniques to synthesize perovskite-type SOFC cathodes and compares their advantages and disadvantages.^14–17 The solid-state synthesis method is a simple and mature technique, however, less homogeneity of the composition, grain growth due to sintering at high temperatures and lower surface area limiting their catalytic behaviour are some of the disadvantages of this method.^14–17 In contrast, the sol–gel technique benefits from a good homogeneity of the chemical composition, better control over grain size and processing at lower temperatures.^15–17 Another common technique is the co-precipitation method which produces single-phase perovskite oxides, but the synthesis time is longer due to a slow precipitation rate.¹⁷ Combustion method is a cost-effective, low-temperature technique to make perovskites with controlled particle size and morphology and enables a higher concentration of dopants.^17–19 The selection of the synthesis route should be based on the cost/scalability, compositional homogeneity and particle-size/morphology requirements which affect the catalytic behaviour of the final product.

Table 1 Comparison of common synthesizing techniques for perovskite-type structure cathodes^14–17

Technique	Advantages	Disadvantages
Solid-state method	Simple and scalable, mature technology	High sintering temperatures and long processing times, inhomogeneity of the composition, grain growth, lower surface area
Sol–gel	Lower annealing temperatures and uniform morphologies	Works in a liquid phase, less-scalability, needs a soluble salt (precursor)
Co-precipitation	Uniform chemical composition with less impurity	Slow rate of precipitation
Solution combustion	Low-temperature process, good control over composition and particle size, allows for higher dopant concentration, fast rate of production	Low surface area

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Perovskites for IT-SOFC cathodes: composition–structure–property relationship

Efforts to design a high-performance IT-SOFC cathode have led to the development of a variety of perovskites with different chemical compositions and structures.¹⁷ Mixed ionic-electronic conductors (MIECs) have gained an increasing interest due to their improved ORR kinetics at IT range benefitting from both ionic and electronic conduction.¹⁸ Elemental doping has been a strong tool for tuning different properties of a perovskite material such as ORR activity, electronic/ionic conductivity, thermal expansion coefficient (TEC), and structural stability. Moreover, surface functionalization, designing novel electrode structures such as core–shell structures¹⁹ and electrospun fibers,²⁰ nanostructuring of the electrode,^16,21–24 and composite cathodes²⁵ are other important techniques that have been used to enhance the electrochemical performance and stability of the electrode under operational conditions.

The A-site in a perovskite structure is occupied by larger size cations compared to B-site cations. Lanthanide elements such as La, Pr, Nd, Sm and Gd are common occupants of the A-site, while common B-site cations include transition metals such as Mn, Co, Fe, Cu and Ni.¹⁷ Single doping or co-doping at A and/or B-sites creates a series of perovskites with different properties. Valence and ionic radii of the dopants are two critical parameters to determine the conduction behaviour of the material. With a similar valence of the dopant and lattice element, the change of electronic conductivity is attributed to the change of structural parameters due to the size effect.⁹ In the case of aliovalent doping in the A-site, the electrical neutrality of the system can be compensated by changing the oxidation state of multivalent cations at B-sites or by formation of lattice oxygen vacancies.⁹ A general formula of can be used to describe the chemical composition of the doped perovskite in which δ (0−1) indicates the lattice oxygen vacancies or oxygen non-stoichiometry.¹⁸

One of the most investigated high-temperature SOFC cathodes is lanthanum strontium manganite (La_1−xSr_xMnO_3−δ, LSM) where Sr²⁺ is doped at the La³⁺ site to introduce oxide ion vacancies due to the charge compensation mechanism in the parent structure of lanthanum manganite (LaMnO₃).^26,27 Sr and other alkaline-earth elements such as Ca²⁺ and Ba²⁺ have been commonly used for La substitution to enhance electrical conductivity.⁹ However, these elements can chemically react with CO₂ and the electrolyte with increasing reactivity from Ca to Ba.⁹ Several studies have investigated the effects of Sr concentration on different properties of LSM perovskites.^26–28 The x ≅ 0.5 composition has shown a high conductivity and high catalytic activity towards the ORR, high thermal and microstructural stability, and good compatibility of TEC with common SOFC electrolytes such as YSZ (TEC_YSZ: 10.5 × 10⁻⁶ K⁻¹ in air at 800 °C).^26,27,29 However, this composition is not an optimum cathode candidate in IT-SOFCs due to the inferior performance at lower operational temperature.^8,26

It is known that replacing Mn with Fe (La_1−xSr_xFeO_3−δ, LSF) or Co (La_1−xSr_xCoO_3−δ, LSC) generates oxygen vacancies that promotes the ORR kinetics at lower temperatures.²⁷ LSC shows high electronic conductivity with a metallic behaviour attributed to the partially filled conduction band with delocalized conduction electrons.³⁰ LSF shows lower electronic conduction compared to LSC, with the hopping mechanism of localized electrons/holes responsible for the electronic conduction.³⁰ While LSC compounds show much improved ionic conduction (0.22 S cm⁻¹)¹⁵ compared to LSM, high contents of Co increase the TEC leading to a mismatch with conventional electrolytes.²⁷ Moreover, the high cost of Co is another limiting factor, motivating the partial or full substitution of Co with other elements.³¹ One of the most investigated MIECs is the La_1−xSr_xCo_1−yFe_yO_3−δ (LSCF) family³⁰ possessing high electronic and ionic conductivity of ∼10² and ∼10⁻² S cm⁻² at 800 °C, respectively, and a TEC value of 14.8–21.4 × 10⁻⁶ K⁻¹ at 500–900 °C.^{15,30,32–34} Different properties of the LSCF can be tuned by changing its chemical composition.

In general, the electronic conductivity of the LSCF is more controlled by the concentration of Fe and Co, while Sr content has a higher impact on controlling the ionic conductivity.²⁷ Also, it is reported that a high concentration of Sr and Co will increase the TEC value.²⁷ Investigations on the stability behaviour of the LSCF perovskites under the operational conditions have revealed some concerns including the chemical instability and TEC mismatch with YSZ electrolyte, surface segregation of Sr and Co, and reactivity of the cathode with contaminants including gas contaminants, i.e., water vapor, CO₂ and SO₂ or volatile species coming from other cell components such as sealants or interconnects, i.e., Cr, B and Si.^17,30

The reactivity of LSCF towards YSZ leads to the formation of secondary phases such as SrZrO₃ at temperatures ≥800 °C.³⁰ To overcome the reactivity issue and to decrease the TEC mismatch, interlayers such as Gd-doped ceria (GDC) or Sm-doped ceria (SDC) with intermediate TEC values have been used.^{14,30,35–37} Segregation of alkaline earth elements, mainly Sr in LSCF materials, is a well-known phenomenon that has been extensively studied^38,39 and an example is shown in Fig. 2a and b. The mechanism of this phenomenon has been under investigation and two major driving forces have been proposed:

Fig. 2 (a) SEM image of a freshly prepared LSCF showing a dense, pinhole-free structure; (b) SEM image of LSCF heat treated at 800 °C, 96 h in the absence of Cr₂O₃ in air. Segregation of Co/Sr-rich particles can be observed. Reproduced from ref. 47 with permission from the Royal Society of Chemistry; (c) Sr surface region concentration vs. oxygen partial pressure. Reprinted from ref. 41, with the permission of AIP Publishing; (d) Arrhenius plots of atomic fractions of surface Sr in pristine and coated samples. Reproduced from ref. 40 with permission from the Royal Society of Chemistry; (e) schematic illustration of LSM-infiltrated LSCF cathode; (f) TEM image of an individual LSCF particle with LSM coating after long-term operation; (g) Fourier-filtered image of the LSCF grain after operation preserving the perovskite structure; (h) convergent beam electron diffraction (CBED) pattern of the shell in (f) showing the loss of crystallinity. Reproduced from ref. 52 with permission from the Royal Society of Chemistry.

(i) The electrostatic interaction between the negatively charged and the positively charged oxygen vacancy attracts the Sr dopant to the surface specifically when there is a high concentration of oxygen vacancies.^40–43 The effect of oxygen partial pressure on Sr segregation behaviour has been investigated by Fister et al.⁴¹ in an epitaxial La_0.7Sr_0.3MnO₃ thin film confirming as shown in Fig. 2c where Sr concentration increases on the surface when oxygen partial pressure is lower; and

(ii) Another important driving force for Sr segregation originates from the difference in Sr and La ionic size introducing elastic forces to the structure. Hence, Sr segregation occurs to minimize this elastic strain energy.^40–42,44 Wen et al.⁴⁰ studied the effect of temperature on Sr surface concentration in a La_1−xSr_xCo_3−δ epitaxial film revealing a weak thermally activated Sr diffusion process following an Arrhenius law with a small activation energy (Fig. 2d). They could effectively suppress the Sr segregation using a ZrO₂ coating (Fig. 2d) because of the reduced surface oxygen vacancies due to the cation exchange between Zr and Co.⁴⁰

Surface segregation of Sr under SOFC operating temperatures leads to the formation of insulating SrO/Sr(OH)₂/SrCO₃ that decreases the kinetics of oxygen surface diffusion, increases the area specific resistance (ASR) and subsequently increases the degradation of the cathode.^30,45–49 Experimental and theoretical studies have shown that Sr-segregation can be remarkably suppressed via surface modifications, doping higher valence cations in the B-site,⁵⁰ and doping larger elements to generate compressive strains.^30,51 In addition, designing novel structures such as nano-architectures with the infiltration/wet impregnation technique, and core–shell structures have been proposed to enhance the performance and decrease the degradation of the LSCF cathodes.^52,53 For example, a nanosized surface layer of La_xSr_1−xMnO_3−δ (x = 0.8 and 0.85) has been fabricated via the infiltration technique on La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ (Fig. 2e–h) leading to inhibition of Sr segregation and enhancement of surface electrocatalytic activity.^52,53

Undesirable reactions of the cathode with contaminants (contaminant poisoning) can cause serious degradation of SOFCs. Degradation mechanisms of cathodes have been reviewed in previous studies^54,55 and a surface functionalization strategy has been proposed to make contaminant-tolerant LSCF cathodes. For example, BaO infiltrated La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ cathode demonstrated Cr poisoning resistance due to the formation of BaCrO₄ instead of SrCrO₄.⁵⁶ Similarly, infiltration has been used to make the BaCeO₃–La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ architecture with enhanced tolerance towards S poisoning by the formation of BaSO₄ instead of SrSO₄.⁵⁷

In addition to LSCF perovskites, several other La-containing compositions have been developed using A or B-site dopants such as Ba,^58,59 Cu,^60–67 Ni,^68,69 Mo,⁷⁰ and Ca^71,72 demonstrating a range of properties summarized in Table 2. A series of La_1−xBa_xCo_0.2Fe_0.8O_3−δ (LBCF) compositions have been investigated for application in IT-SOFC cathodes showing lower electrical conductivities than their LSCF counterparts,^58,59 with a maximum value of 100 S cm⁻¹ in the temperature range of RT-1000 °C for La_0.6Ba_0.4Co_0.2Fe_0.8O_3−δ.⁵⁸ However, the main advantage of this family is their high resistance to Cr poisoning and good polarization performance stability compared to conventional LSM and LSCF cathodes.⁷³ Cu-doped compositions such as La_0.6Sr_0.4Co_1−yCu_yO_3−δ,⁶⁰ La_1−ySr_yMn_1−xCu_xO_3−δ,^61,64 and La_xSr_1−xFe_1−yCu_yO_3−δ^{62,63,65,66,74} have been investigated. Specifically, Co-free compositions with Cu have attracted much attention due to the decreased cost and TEC and sufficient catalytic activity.^31,62,74,75 The Co-free composition of La_0.7Sr_0.3Ti_0.1Fe_0.6Ni_0.3O_3−δ was prepared using Ni and Ti as B-site dopants demonstrating high electrical conductivity (318 S cm⁻¹ at 700 °C), low polarization resistance and good stability in both oxidizing and reducing environments, which are desirable for application in symmetrical SOFCs.⁶⁸ Using Ni and Fe in B-sites and La in A-sites, TEC value decreased to 11.4 × 10⁻⁶ K⁻¹ in LaNi_0.6Fe_0.4O₃ in the temperature range of 30–1000 °C and an electrical conductivity of 580 S cm⁻¹ at 800 °C was achieved.⁶⁹ The Mo-doped composition of La_0.5Sr_0.5Fe_0.9Mo_0.1O_3−δ also showed a lower TEC value of 13.2 × 10⁻⁶ K⁻¹ at 300–600 °C with good structural stability in both oxidizing and reducing atmospheres.⁷⁰ The Ca-doped Sr/Co-free composition of La_0.65Ca_0.35FeO_3−δ has been prepared showing a high oxygen permeation flux and electrical conductivity reaching ∼100 S cm⁻¹ at 600–800 °C.^71,72

Table 2 Summary of the most important properties of IT-SOFC perovskites^{15,28,59,61–72,74,76–78,80–82,84,86,92–97,99–104,106–108,122,124,125,127,129–131,136–143a}

Composition	Conductivity (S cm^−1)		TEC (× 10^−6 K^−1)	Cell performance
	σ e	σ i		Electrolyte	R p (Ω cm^2)	ASR (Ω cm^2)
a SDC: samarium-doped ceria, LSGM: lanthanum strontium gallium magnesium oxide, GDC/CGO: gadolinium-doped ceria.
La0.8Sr0.2MnO3−δ (ref. 15, 28 and 136–138)	180 (700 °C)	5.93 × 10^−7 (900 °C)	11–13 (30–1000 °C)	—	—	—
	300 (900 °C)
La1−xSrxFeO3−δ (ref. 15 and 28)	129–369 (500–900 °C)	0.205–5.6 × 10^−3 (500–900 °C)	12.2 (30–1000 °C)	—	—	—
La1−xSrxCoO3−δ (ref. 15, 28 and 139)	1200–1360 (500–900 °C)	0.22 (500–900 °C)	18–20 (300–750 °C)	SDC	1.0825 (x = 0.3, 0.4)	—
La0.6Sr0.4CoO3−δ (ref. 76)	∼1990–2820 (300–750 °C)	—	21.3 (50–800 °C)	LSGM	—	—
La0.6Sr0.4Co0.2Fe0.8O3−δ (ref. 59, 140 and 141)	∼252–330 (300–900 °C)	0.23 (900 °C)	15.3 (100–600 °C)	GDC	0.04 (800 °C)	—
La0.4Sr0.6Co0.2Fe0.8O3−δ (ref. 59 and 140)	219 (900 °C)	0.4 (900 °C)	16.8 (100–400 C)	—	—	—
La0.2Sr0.8Co0.2Fe0.8O3−δ (ref. 59)	120 (900 °C)	0.62 (900 °C)	—	—	—	—
La0.2Sr0.8Co0.8Fe0.2O3−δ (ref. 59)	310 (900 °C)	0.87 (900 °C)	—	—	—	—
La0.8Sr0.2Co0.8Fe0.2O3−δ (ref. 142)	∼825–1000 (300–800 °C)	—	20.7 (100–900 °C)	—	—	—
La0.8Sr0.2Co0.2Fe0.8O3−δ (ref. 142)	∼100–150 (300–800 °C)	—	15.4 (100–800 °C)	—	—	—
La0.9Sr0.1Co0.2Fe0.8O3−δ (ref. 140)	∼15–60 (300–800 °C)	—	16 (300–900 °C)	—	—	—
La0.6Ba0.4Co0.2Fe0.8O3−δ (ref. 58, 59 and 143)	100 (750 °C)	0.01 (900 °C)	—	—	—	—
	123 (900 °C)
La0.6Ba0.4Co0.2Fe0.8O3−δ (ref. 59)	123 (900 °C)	0.01 (900 °C)	—	—	—	—
La0.4Ba0.6Co0.2Fe0.8O3−δ (ref. 59)	57 (900 °C)	0.33 (900 °C)	—	—	—	—
La0.2Ba0.8Co0.2Fe0.8O3−δ (ref. 59)	19 (900 °C)	0.37 (900 °C)	—	—	—	—
La0.54Sr0.46Fe0.8Cu0.2O3−δ (ref. 62)	9.029 (600 °C)	—	—	SDC	—	—
La0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 66)	300 (500 °C)	—	15.8(9) (30–850 °C)	GDC	—	1.15 (600 °C)
La0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 65)	—	—	17.7 (25–900 °C)	SDC	0.4 (700 °C)	—
La0.6Sr0.4Fe0.8Cu0.2O3−δ (ref. 63)	190–238 (600–800 °C)	—	14.6 (RT-850 °C)	SDC	0.138 (750 °C)
La0.54Sr0.46Fe0.8Cu0.2O3−δ (ref. 62)	9.029 (600 °C)	—	—	SDC	—	—
La0.7Sr0.3Mn0.8Cu0.2O3−δ (ref. 61)	208.4 (750 °C)	—	—	—	—	—
La0.8Sr0.2Mn0.8Cu0.2O3−δ (ref. 64)	190 (750 °C)	—	—	YSZ	—	4.3 (750 °C)
La0.8Sr0.2Fe0.8Cu0.2O3−δ (ref. 67)	184–150 (550–750 °C)	—	—	LSGM	—	0.25 (750 °C)
La0.7Sr0.3Ti0.1Fe0.6Ni0.3O3−δ (ref. 68)	318 (700 °C)	—	—	LSGM	0.185 (700 °C)	—
LaNi0.6Fe0.4O3 (ref. 69)	580 (800 °C)	—	11.4 (30–1000 °C)	—	—	—
La0.5Sr0.5Fe0.9Mo0.1O3−δ (ref. 70)	73–70 (600–800 °C)		14 (300–900 °C)	SDC	0.25 (700 °C)	—
La0.65Ca0.35FeO3−δ (ref. 72)	∼100 (600–800 °C)	—	—	—	—	—
La0.65Ca0.35FeO3−δ (ref. 71 and 72)	100–200 (600 °C)	—	—	YSZ/GDC	0.255 (700 °C)	—
La0.4Ca0.6Co0.2Fe0.8O3−δ (ref. 59)	52 (900 °C)	0.03 (900 °C)	—	—	—	—
La0.4Ca0.6Co0.8Fe0.2O3−δ (ref. 59)	296 (900 °C)	0.01 (900 °C)	—	—	—	—
Pr0.5Sr0.5Fe0.9Mo0.1O3−δ (ref. 70)	59–53 (600–800 °C)	—	13.6 (300–900 °C)	SDC	0.5 (700 °C)	—
Pr0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 74)	∼60–140 (300–750 °C)	—	17.3 (RT-800 °C)	SDC	0.036 (800 °C)	—
Pr0.6Sr0.4CoO3−δ (ref. 76)	∼2400–1780 (300–750 °C)	—	19.5 (50–800 °C)	LSGM	—	—
Pr0.8Sr0.2Fe0.8Co0.2O3−δ (ref. 78)	75.8 (800 °C)	1.54 × 10^−3 (800 °C)	12.8 (30–1000 °C)	—	—	—
Pr0.8Sr0.2Mn0.8Co0.2O3−δ (ref. 78)	83.17 (800 °C)	3 × 10^−5 (800 °C)	10.9 (30–1000 °C)	—	—	—
Pr0.65Sr0.3MnO3−δ (ref. 78)	208.92 (800 °C)	3.4 × 10^−4 (800 °C)	11.6 (30–1000 °C)	—	—	—
Nd0.9Sr0.1Fe0.1Co0.9O3−δ (ref. 80)	0.137 (600 °C)	—	—	—	—	—
Nd0.5Sr0.5Fe0.9Mo0.1O3−δ (ref. 70)	63–65 (600–800 °C)	—	13.3 (300–900 °C)	SDC	0.44 (700 °C)	—
Nd0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 74)	∼40–100 (300–750 °C)	—	17.19 (RT-800 °C)	SDC	0.089 (800 °C)	—
Nd0.75Sr0.25Fe0.2Co0.8O3−δ (ref. 81)	30 (700 °C)	—	—	LSGM	—	0.1 (800 °C)
Nd0.6Sr0.4CoO3−δ (ref. 76)	∼2240–1400 (300–750 °C)	—	18.7 (50–800 °C)	LSGM	—	—
Nd0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 82)	124 (700 °C)	—	14.7 (25–800 °C)	SDC	0.071 (700 °C)	—
Sm0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 74)	∼40–80 (300–750 °C)	—	16.72 (RT-800 °C)	SDC	0.097 (800 °C)	—
Sm0.6Sr0.4CoO3−δ (ref. 76)	∼1950–1320 (300–750 °C)	—	18 (50–800 °C)	LSGM	—	—
Sm0.3Sr0.7Nb0.08Co0.92O3−δ (ref. 84)	315 (350 °C)	—	—	SDC	—	0.062 (600 °C)
Sm0.5Sr0.5Fe0.8Cr0.2O3−δ (ref. 86)	7.32 (766 °C)	—	4.7 (100–800 °C)	—	—	—
Gd0.5Sr0.5Fe0.8Cu0.2O3−δ (ref. 74)	∼25–35 (300–750 °C)	—	12.89 (RT-800 °C)	SDC	0.16 (800 °C)	—
Gd0.6Sr0.4CoO3−δ (ref. 76)	∼1260–890 (300–750 °C)	—	17.1 (50–800 °C)	LSGM	—	—
Gd0.5Sr0.5CoO3 (ref. 77)	—	—	—	CGO	—	0.1–0.2 (650 °C)
Ba0.9Co0.7Fe0.2Ni0.1O3−δ (ref. 92)	—	—	—	GDC	0.046 (600 °C)
Ba0.9Co0.7Fe0.2Nb0.1O3−δ (ref. 93)	13.9 (700 °C)	—	13.2 (600 °C)	LSGM	0.07 (700 °C)	—
Ba0.5Sr0.5Zn0.2Fe0.8O3−δ (ref. 94)	9.4 (590 °C)	—	—	SDC	0.48 (650 °C)	—
Ba0.9Sr0.1Co0.9In0.1O3−δ (ref. 95)	8.7–13.6 (600–800 °C)	—	17.04 (30–1000 °C)	SDC	—	0.079 (700 °C)
SrNb0.1Co0.9O3−δ (ref. 96)	∼135–75 (350–700 °C)	—	—	SDC	—	0.040 (650 °C)
						0.094 (600 °C)
SrNb0.1Co0.9O3−δ (ref. 97)	461–145 (300–800 °C)	—	24.2 (30–1000 °C)	LSGM	—	0.21 (650 °C)
Sr0.95Nb0.1Co0.9O3−δ (ref. 102)	276–129 (450–750 °C)	—	95.8 (25–800 °C)	SDC	—	0.052 (500 °C)
SrNb0.1Co0.9O3−δ (ref. 99)	—	0.34 (650 °C)	—	—	—	—
SrNb0.2Co0.8O3−δ (ref. 101)	∼155–105 (400–700 °C)	0.28 (700 °C)	—	SDC	—	0.21–0.24 (550 °C)
SrTa0.2Co0.8O3−δ (ref. 101)	∼150–100 (400–700 °C)	0.31 (700 °C)	—	SDC	—	0.092–0.097 (550 °C)
SrTa0.05Co0.95O3−δ (ref. 103)	590–210 (400–700 °C)	—	—	SDC	—	0.11–0.089 (550 °C)
SrTa0.1Co0.9O3−δ (ref. 99)	—	0.17 (650 °C)	—	—	—	—
SrTa0.1Co0.9O3−δ (ref. 104)	∼325–175 (550–700 °C)	—	23.6 (500–900 °C)	—	—	—
SrCo0.95Ti0.05O3−δ (ref. 107)	268–160 (600–800 °C)	—	21.2 (30–1000 °C)	LSGM	—	0.17 (700 °C)
SrCo0.95Ti0.05O3−δ (ref. 106)	>150 (600 °C)	—	25.28 (400–850 °C)	LSGM	—	0.016 (850 °C)
SrCo0.97V0.03O3−δ (ref. 106)	>4 (600 °C)	—	13.40 (400–850 °C)	LSGM	—	0.025 (850 °C)
Sr0.7Y0.3CoO2.65−δ (ref. 108)	735 (650 °C)	—	19.6 (25–800 °C)	LSGM	—	0.11 (800 °C)
Bi0.5Sr0.5Fe0.8Co0.2O3−δ (ref. 124)	19.5–25.1 (600–800 °C)	—	12.2–14.7 (300–800 °C)	SDC	—	0.086 (700 °C)
Bi0.5Sr0.5Fe0.85Ti0.15O3−δ (ref. 125)	0.6–2.4 (300–800 °C)	—	13.4 (50–800 °C)	CGO	0.085 (700 °C)	—
Bi0.5Sr0.5Fe0.9Sn0.1O3−δ (ref. 127)	—	—	12.9 (50–800 °C)	CGO	0.09 (700 °C)	—
Bi0.5Sr0.5Fe0.95P0.05O3−δ (ref. 122)	—	—	13.5 (50–800 °C)	GDC	0.18 (700 °C)	—
SrFe0.9Si0.1O3−δ + 50 wt% Gd0.1Ce0.9O1.95 (ref. 129)	∼25–50 (300–800 °C)	—	—	—	—	0.08 (800 °C)
La0.6Sr0.4Co0.76Fe0.19B0.05O3−δ + 50 wt% Gd0.1Ce0.9O1.95 (ref. 130)	1253–1096 (600–800 °C)	—	—	—	—	0.08 (800 °C)
La0.6Sr0.4Co0.78Fe0.195Si0.025O3−δ + 50 wt% Gd0.1Ce0.9O1.95 (ref. 130)	853–682 (600–800 °C)	—	—	GDC	—	0.11 (800 °C)
Ba0.95La0.05Fe0.95P0.05O3−δ (ref. 131)	∼14–7 (500–750 °C)	—	25.48 (100–800 °C)	SDC	—	0.023 (700 C)

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By replacing La with other lanthanides including Pr,^{70,74,76–79} Nd,^{70,74,76,79–82} Sm^{74,76,79,83–89} and Gd,^74,76,77,79 a variety of perovskite compositions can be generated in which the A site is occupied partly by a Ln element together with Sr or Ca, and the B site is occupied with transition metals such as Co, Fe, Mn, Cu and Mo. In the lanthanide family, the ionic size decreases moving from La to Pr, Nd, Sm and Gd leading to a lower ionicity of Ln–O bonds.⁷⁶ This change is beneficial to decrease the thermal mismatch of the cathode with common electrolytes by decreasing the TEC of the cathode. However, the electronic conductivity and electrochemical performance will be sacrificed (Fig. 3).^{70,76,81,90,91}

Fig. 3 Variation of DC electrical conductivity (σ) and thermal expansion coefficient (TEC) in Ln_0.6Sr_0.4CoO_3−δ (Ln = La, Pr, Nd, Sm and Gd) in air. Values are extracted from ref. 70, 76, 81, 90 and 91.

Another important family of single perovskites for IT-SOFC cathodes are alkaline-earth based compositions including Sr- and/or Ba-based^92–95 compounds. Several compositions based on SrCoO_3−δ doped with Nb such as SrNb_0.1Co_0.9O_3−δ,^96–100 SrNb_0.2Co_0.8O_3−δ ¹⁰¹ and A-site deficient Sr_0.95Nb_0.1Co_0.9O_3−δ ¹⁰²) have been investigated. These perovskites showed high electrical conductivity and good phase structure stability.^96,97 A nanoscale layer of SrNb_0.1Co_0.9O_3−δ was coated on a (La_0.6Sr_0.4)_0.95(Co_0.2Fe_0.8)O_3−δ cathode and demonstrated an improved ORR activity compared to the non-coated electrode.¹⁰⁰ Ta-doped compositions, SrCo_1−xTa_xO_3−δ, are another subgroup in this class of perovskites.^{99,101,103,104} It was found that doping small amounts of Ta (20 mol%) stabilizes the crystal structure and enhances the ORR activity.¹⁰³ These changes were attributed to the increased oxygen surface exchange originating from the effects of Ta⁵⁺ on the oxidation states of Co ions.¹⁰³ Wang et al.¹⁰⁴ compared different properties of a SrTa_0.1Co_0.9O_3−δ cathode with a SrNb_0.1Co_0.9O_3−δ material and proved that the Ta-doped counterpart shows better thermal and electrochemical stability due to the stronger Ta–O bonds.¹⁰⁴ In another interesting study, a nanoscale SrTa_0.1Co_0.9O_3−δ layer was used as the capping layer on a commercial (La_0.6Sr_0.4)_0.95(Co_0.2Fe_0.8)O_3−δ-GDC composite.¹⁰⁵ This cathode showed excellent Cr-resistant properties with a good ORR activity. SrO-free surface of this cathode showed a much lower polarization resistance and degradation rate compared to the uncoated cathode.¹⁰⁵ SrCoO_3−δ doped with other elements such as Ti/V^106,107 and Y¹⁰⁸ has also been studied. Ti doping stabilized the crystal structure, improved the electrical conductivity, and decreased the polarization resistance^106,107 while V-doped compositions showed much lower TEC values.¹⁰⁸

Co-free Sr-based compounds have also been investigated, mainly based on a strontium ferrite (SrFeO_3−δ) composition with dopants such as Nb,^109–112 Ti,¹¹³ Zr¹¹⁴ and Cu¹¹⁵ on B-sites.¹⁵ Other Sr-based compositions such as Sr_1−xCe_xMnO₃ ^15,116 and SrZr_1−xNi_xO₃ have also been investigated.^15,78,117 Ba-based compositions are another group of single perovskite cathodes. Ba_1−xSr_xCo_1−yFe_yO_3−δ (BSCF) compositions have shown good electrochemical performance at low temperatures,¹¹⁸ however, they suffer from poor stability towards CO₂ originating from the susceptibility of alkaline-earth elements in reaction with CO₂.¹¹⁹ The B-site doping strategy has been used to enhance either the electrochemical performance or structural stability of the perovskite using Nb,^92,93 Ln,¹²⁰ Ni,⁹² and In.⁹⁵ Co-free compositions have also been prepared such as Ba_1−xSr_xFe_1−yMo_yO_3−δ ^31,121 and Ba_1−xSr_xFe_1−yZn_yO_3−δ ⁹⁴) demonstrating promising performance as an IT-SOFC cathode.

Bismuth-based perovskites are novel compositions that have been recently investigated by several researchers showing promising performance as IT-SOFC cathodes.^122–127 Replacing Ba²⁺ with Bi³⁺ increases the structural stability at the operating temperatures.^124,126,128 Moreover, owing to its 6s lone electron pair, Bi³⁺ demonstrates high polarizability, enhancing the mobility of oxygen vacancies.¹²⁴ Bismuth strontium ferrites have been designed as a new family of Co-free perovskites with low area-specific resistance, high oxygen flux density and improved kinetics of surface exchange reactions.¹²⁷ However, the reported conductivity values are still not sufficient for an IT-SOFC cathode. Another novel concept that has been introduced recently is using non-metal dopants such as Si,^129,130 P^122,130,131 and B.¹³⁰ Slater et al.¹³² reported the successful incorporation of Si into the structures of SrCoO_3−δ and SrMnO_3−δ. They reported a higher conductivity for the Si-doped perovskites due to the transformation from a hexagonal structure into a cubic perovskite.¹³² A similar observation was reported for a phosphate/sulfate-doped SrCoO_3−δ and the improvement in the conductivity was attributed to the change from a 2H- to 3C-perovskite.¹³³ The latter phase was found to be metastable when annealing at intermediate temperatures, however, co-doping with Fe was found to improve the stability.¹³³ The same group reported the synthesis of a Si-doped SrFeO_3−δ with an enhanced conductivity compared to the undoped structure attributed to the change of the crystal structure from a tetragonal symmetry to a cubic perovskite with disordered oxygen vacancies.¹²⁹ They also showed that increasing the Si level to higher than 10% decreases the conductivity due to the blocking effect of Si on electronic conduction pathways.¹²⁹ Doping silicate, borate, and phosphate into La_0.6Sr_0.4Co_0.8Fe_0.2O_3−δ and Sr_0.9Y_0.1CoO_3−δ compositions have been reported.¹³⁰ Introduction of oxide ion vacancies by oxyanion doping was found to be responsible for the improvement of the electronic conductivity in doped La_0.6Sr_0.4Co_0.8Fe_0.2O_3−δ material. However, in the case of Sr_0.9Y_0.1CoO_3−δ, oxyanion doping decreased the electronic conductivity due to the disruption of conduction pathways.¹³⁰ Interestingly, oxyanion doping of both compositions improved the stability towards CO₂ and the observation was attributed to the decrease in the basicity of the system by introduction of acidic dopants.¹³⁰

Stabilization of the cubic perovskite structure has also been observed when phosphate and borate were incorporated into the Ba_1−xSr_xCo_0.8Fe_0.2O_3−δ material with a small improvement of the electronic conductivity for low levels of dopants.¹³⁴ However, borate-doped La_1−xSr_xMnO_3−δ resulted in a lower electronic conductivity in comparison with the undoped material because of the lower concentration of Mn⁴⁺ in the doped-sample.¹³⁵ P-doped Ba_0.95La_0.05Fe_0.95P_0.05O_3−δ resulted in an enhancement in the electrical conductivity as well as a better electrocatalytic activity.¹³¹ DFT studies predicted a lower formation energy of oxygen vacancies and migration barrier by introduction of P into the structure. Experimental results further confirmed the DFT findings and showed that P-doing into the Fe sites increases the surface exchange rate and the diffusion coefficient in a symmetrical cell leading to an improved ORR performance.¹³¹ Similar observations have been reported for a P-doped (Bi, Sr)FeO_3−δ cathode.¹²²Table 2 summarizes the properties of different LSCF compositions along with other perovskite-type SOFC cathodes.^{15,28,59,61–72,74,76–78,80–82,84,86,92–97,99–104,106–108,122,124,125,127,129–131,136–143}

Double perovskite-type structure SOFC cathodes

Oxygen deficient double perovskite oxides are gaining prominent attention as oxygen electrodes in intermediate temperature SOFCs due to their enhanced catalytic properties, oxygen transport, and stability over single perovskite oxides.^144,145 Their compositional formula is generally represented by AA′BB′O_5+δ, where the A-site ion is a rare-earth element, A′ is an alkali or alkaline earth metal, B and B′ are transition metals and O is oxygen. They consist of alternating layers of [AO_δ]–[BO]–[A′O]–[BO], where oxygen vacancies, δ, are majorly localised in the [AO_δ] planes. The concentration of oxygen vacancies within their sublattices facilitates the rapid diffusion of oxygen ions, giving rise to superior electrochemical performances.

Double perovskite crystal structure

Double perovskite oxides possess the same crystal structure as single perovskites. Their unit cell size depends on the ordering of the A, A′, B and/or B′ cations in the sublattices. On investigating the relationship between A- and B-site cation ordering, Knapp and Woodward observed that while A-site cation ordering was inclined to occur in a layered manner, B-site cation ordering usually assumes a rock salt structure.¹⁴⁶ In the layered structure, A and A′ cations alternately occupy the same crystallographic site. This alternate arrangement is favoured by a substantial difference between the ionic radii and/or charge of the elements which occupy the A and A′ sites.¹⁴⁷ To attain the layered structure, large atoms such as Ba are typically doped on the A′ site such as in the REBaCo₂O_5+δ (RE = Rare Earth) perovskites. The layered ordering of A-site cations in the AA′B₂O_5+δ perovskite leads to a tetragonal P4/mmm space group symmetry (a_p × a_p × 2a) where the unit cell is doubled along the c-axis.^148–150 In the case of an insignificant difference in the ionic radii, the perovskite oxide will lose the long-range ordering between the A and A′ cations resulting in randomly distributed A cations with a cubic structure (Pm3̄m) as observed in LaBaCo₂O_5+δ.^151,152 At a given temperature, the unit cell volume of REBaCo₂O_5+δ decreases with decreasing Ln³⁺ ionic radius.

The double perovskite can accommodate a large concentration of vacancies in the oxygen sublattice without structural collapse.¹⁵³ This occurs as a result of the distortion of the [BO] framework induced by the large size or charge difference between the A and A′ cations. To attain stability and/or electroneutrality, vacancies are created in the oxygen sublattice which cause the double perovskite stoichiometry to depart from AA′BB′O₆ to an oxygen-deficient double perovskite oxide, AA′BB′O_5+δ where 0 < δ ≤ 1. These vacancies when ordered create a pathway for fast oxygen diffusion. Thus, this family of perovskite oxides provides mixed ionic and electronic conduction required for IT-SOFC cathodes. Taskin et al. discovered that transforming Gd_0.5Ba_0.5MnO_3−δ with randomly distributed Gd and Ba cations into a layered A site GdBaMn₂O_5+δ substantially enhanced the oxygen diffusivity.¹⁵⁴ In addition to cation ordering, oxygen vacancies directly influence the crystal system and lattice parameters of double perovskite oxides. Anderson et al. reported structural transition phases in REBaCo₂O_6−δ (RE = Pr³⁺ (1.13 Å), Nd³⁺ (1.11 Å), Sm³⁺ (1.08 Å), Eu³⁺ (1.07 Å), Gd³⁺ (1.05 Å), Tb³⁺ (1.04 Å), Dy³⁺ (1.027 Å), Ho³⁺ (1.02 Å)) perovskite oxide which corresponded with the ionic radius- and temperature-dependent oxygen non stoichiometry summarised in Table 3.^{152,155–180,186} Several studies have reported a similar occurrence,^155–159 inferring that oxygen content fundamentally controls the structural symmetry of the LnBaCo₂O_5+δ series. Streule et al. suggested that structural transitions observed in double perovskites are related to the order/disorder of the oxygen vacancies.¹⁴⁹ When the ordering of oxygen vacancies occurs, the unit cell along the b-axis doubles, leading to the formation of an orthorhombic a_p × 2a_p × 2a_p (Pmmm) structure. Thus, at certain perovskite oxide composition stoichiometries, vacancy ordering is observed within the sublattice. Since oxygen content varies with temperature, heating double perovskite oxide compositions to certain temperature ranges will drive a disorder in the arrangement of oxygen vacancies.

Table 3 Summary of the crystallographic parameters and room temperature oxygen stoichiometry of selected IT-SOFC double perovskite cathodes^{152,155–180,186}

Composition	Space group	Cell constants				Oxygen contents (5 + δ)
		a (Å)	b (Å)	c (Å)	V (Å^3)
LaBaCo2O5+δ (ref. 152)	Pm3̄m	3.881	—	—	58.456	6.00
LaBaCoCuO5+δ (ref. 160)	Pmmm	3.922	3.9360	11.7073	180.74	—
LaBaCuFeO5+δ (ref. 161)	Immm	5.5586	5.5550	7.8155	241.33	—
LaSrMnCoO5+δ (ref. 162)	Fm3m	7.6891	—	—	454.597
PrBaCo2O5+δ (ref. 163)	P4/mmm	3.909	—	7.638	116.8	5.64
PrBaCoFeO5+δ (ref. 164)	P4/mmm	3.9184	—	7.6568	117.56	5.79
PrBaCo2/3Fe2/3Cu2/3O5+δ (ref. 165)	P4/mmm	3.904	—	7.651	116.63	—
PrBaFe2O5+δ (ref. 166)	Pmmm	3.928	3.934	7.794	120.46	5.884
PrBa0.5Sr0.5Co2O5+δ (ref. 167)	P4/mmm	7.758	—	7.704	463.70	5.498
PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (ref. 168)	P4/mmm	3.871	—	7.757	116.212	6.00
PrBa0.5Sr0.5CoFeO5+δ (ref. 168)	P4/mmm	3.875	—	7.767	116.652	6.00
PrBa0.8Ca0.2Co1.5Fe0.5O5+δ (ref. 169)	P4/mmm	3.871	—	7.703	115.421	5.81
NdBaCo2O5+δ (ref. 170)	P4/mmm	3.896	—	7.619	115.65	5.85
NdBaCoCuO5+δ (ref. 171)	P4/mmm	3.920	—	7.683	118.049	5.782
NdBaCoFeO5+δ (ref. 164)	P4/mmm	3.9090	—	7.6252	116.526	5.67
NdBaCo2/3Fe2/3Cu2/3O5+δ (ref. 172)	P4/mmm	3.923	—	7.696	118.5	5.44
NdBaCo1.6Ni0.4O5+δ (ref. 173)	P4/mmm	3.9022	—	7.6200	116.03	—
NdBaFe2O5+δ (ref. 170)	Pm3̄m	3.930	—	—	60.70	—
NdBa0.5Sr0.5CoCuO5+δ (ref. 171)	P4/mmm	3.871	—	7.664	114.851	5.789
NdBa0.5Sr0.5CoFeO5+δ (ref. 168)	P4/mmm	3.864	—	7.718	115.226	6.00
NdBa0.5Sr0.5Co2O5+δ (ref. 167)	P4/mmm	7.669	—	7.685	452.0	5.235
NdSrCo2O5+δ (ref. 174)	Pbnm	5.3740	5.4201	7.6020	221.443	6.00
SmBaCo2O5+δ (ref. 175)	Pmmm	3.889	7.839	7.563	230.22	5.62
SmBaCo1.6Fe0.4O5+δ (ref. 175)	Pmmm	3.888	7.826	7.599	231.29	—
SmBaCo0.5Mn1.5O5+δ (ref. 176)	Cmmm	7.736	7.799	7.692	464.07	5.98
SmBaCo1.6Ni0.4O5+δ (ref. 173)	Pmmm	0.392	0.389	0.758	0.116	—
SmSrCo2O5+δ (ref. 177)	Pbnm	5.403	5.3830	7.6264	221.788	6.00
YBaCo2O5+δ (ref. 152)	P4/mmm	3.874	—	7.483	112.304	5.41
YBaCo1.4Cu0.6O5+δ (ref. 178)	P4/mmm	11.658	—	7.546	1025.539	—
YBaCo1.8Fe0.2O5+δ (ref. 179)	P4/mmm	3.8807	—	7.519	113.23	—
EuBaCo2O5+δ (ref. 180)	P4/mmm	3.882	—	7.541	229.36	5.40
GdBaCo2O5+δ (ref. 181)	Pmmm	3.876	3.912	7.541	114.367	5.61
GdBaCoCuO5+δ (ref. 171)	P4/mmm	3.894	—	3.510	115.288	5.643
GdBaCoFeO5+δ (ref. 170)	P4/mmm	3.903	—	7.643	116.43	6.00
GdBaFeNiO5+δ (ref. 182)	P4/mmm	3.915	—	7.598	116.5	5.26
GdBa0.5Sr0.5Co2O5+δ (ref. 183)	P4/mmm	3.8624	—	7.5578	112.74	—
GdBa0.5Sr0.5CoCuO5+δ (ref. 171)	P4/mmm	3.866	—	7.576	113.231′	5.66
GdBa0.5Sr0.5CoFeO5+δ (ref. 183)	P4/mmm	3.8710	—	7.6368	114.43	—
GdBa0.5Sr0.5Co1.5Fe0.5O5+δ (ref. 183 and 184)	P4/mmm	3.8596	—	7.5802	112.90	5.75
GdSrCo2O5+δ (ref. 181)	Pnma	5.373	7.572	5.402	219.763	6.00
TbBaCo2O5+δ (ref. 180)	P4/mmm	3.867	—	7.516	112.39	5.40
DyBaCo2O5+δ (ref. 180 and 185)	P4/mmm	3.879	—	7.505	112.95	5.30
HoBaCo2O5+δ (ref. 186)	P4/mmm	3.873	—	7.496	112.44	5.30

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Double perovskite IT-SOFC cathodes: composition–structure–property relationship

Substantial efforts have been dedicated towards tuning the physical properties of double perovskite oxide catalysts through a systematic partial or complete substitution of the A and/or B sites. Whereas the A-site cation substitution primarily controls the electrochemical performance of the perovskite, B-site cation substitutions influence the electronic structure and catalytic activities of the perovskite oxides as the redox reactions occur on the B site. Studies that employ A-site cation substitution typically dope a trivalent lanthanide element with a larger divalent alkaline earth metal.^187–189 The most common A site doping is substituting lanthanide elements with Ba and/or Sr. The A-site layered LnBaCo₂O_5+δ (Ln = Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Y) series has been widely investigated.^{151,152,155,185,190–193} This compositional family shows promising ionic transport properties, surface exchange kinetics and high electrical conductivities. Kim et al. reported a surface exchange rate coefficient, k = 6.5 × 10⁻⁵ cm s⁻¹ and diffusion coefficient of oxygen ions, D = 3.6 × 10⁻⁷ cm² s⁻¹ at 500 °C for PrBaCo₂O_5+δ (PBCO)¹⁹⁴ as compared to the widely known La_0.6Sr_0.4Co_0.2Fe_0.8O_3−δ (LSCF) cathode with k = 3.3 × 10⁻⁹ cm s⁻¹ and D = 1.2 × 10⁻¹⁰ cm² s⁻¹ at 500 °C.¹⁹⁵ They also reported ionic conductivity values of 0.056, 0.04 and 0.11 S cm⁻¹ and total electrical conductivities of 700, 464 and 425 S cm⁻¹ for LnBaCo₂O_5+δ (Ln = La, Nd, Sm) respectively at 900 °C.¹⁵⁵ Their results reveal that the ionic and electronic conductivities of LnBaCo₂O_5+δ decrease with decreasing Ln³⁺ size (Fig. 4a) due to the increase in the concentration of oxygen vacancies (Fig. 4b) and a consequent disturbance of the O²⁻–Co³⁺/Co⁴⁺–O²⁻ electronic band overlap.¹⁵² This trend is observed in other studies which report that while the concentration of oxygen vacancies increases with decreasing radius of the lanthanide ions, the total electrical conductivity and the TEC decrease down the series from Ln = La to Ho, and Y.^151,196 In addition to inducing ordering in the crystallographic arrangement of the A and A′ sites, Ba enhances the ionic conductivity of double perovskite oxides because of its large ionic radii. Several studies have reported enhanced electrical conductivity and electrochemical performance with the substitution of Ba²⁺ for Sr²⁺ in layered double perovskite oxides.^177,197 Sr doping on Ba sites increases the oxygen content and thus an increase in the concentration of electronic holes. Table 4 summarises of the most important properties of selected IT-SOFC double perovskite cathodes.^{152,156–165,167,169–171,175–180,186,197,220,224,229–241}

Fig. 4 (a) Temperature-dependent DC conductivity of LnBaCo₂O_5+δ (Ln = La, Nd, Sm, Gd, Y) samples in air. (b) Variation of oxygen content and cobalt oxidation state with temperature in air: (a) Ln = La, (b) Ln = Nd, (c) Ln = Sm, (d) Ln = Gd, (e) Ln = Y. Reproduced with permission.¹⁵²

Table 4 Summary of the most important properties of selected IT-SOFC double perovskite cathodes^{152,156–165,167,169–171,175–180,186,197,220,224,229–241}

Composition	Conductivity (S cm^−1)		TEC (× 10^−6 K^−1)	Cell performance			Power density (W cm^−2)
	σ e	σ i		Electrolyte	R p (Ω cm^2)	ASR (Ω cm^2)
LaBaCo2O5+δ (ref. 152 and 229)	1443 (600 °C)	—	24.3	SDC	—	0.056 (700 °C)	—
LaBaCuCoO5+δ (ref. 220)	392 (600 °C)	—	—	SDC	—	0.11 (700 °C)	0.60 (800 °C)
LaBaCuFeO5+δ (ref. 220)	123 (600 °C)	—	—	SDC	—	0.21 (700 °C)	0.56 (800 °C)
LaSrMnCoO5+δ (ref. 230)	111 (600 °C)	—	15.8	SDC	0.048 (800 °C)	—	0.565 (800 °C)
PrBaCo2O5+δ (ref. 163 and 224)	844 (600 °C)	—	20.4	LSGM	0.75 (700 °C)	0.061 (700 °C)	0.41 (800 °C)
PrBaCoCuO5+δ (ref. 224)	144 (600 °C)	—	15.2	SDC	—	0.047 (700 °C)	0.79 (700 °C)
PrBaCoFeO5+δ (ref. 164)	203 (600 °C)	—	21.0	LSGM	—	0.098 (750 °C)	0.75 (800 °C)
PrBaFe2O5+δ (ref. 166 and 231)	25.4 (600 °C)	—	17.2	SDC	0.48 (700 °C)	—	0.15 (650 °C)
PrBaCo2/3Fe2/3Cu2/3O5+δ (ref. 165)	144 (600 °C)	—	16.6	GDC	0.038 (800 °C)	—	0.659 (800 °C)
PrSrCo2O5+δ (ref. 197)	2084 (600 °C)	—	—	GDC	—		0.73 (600 °C)
PrBa0.5Sr0.5Co2O5+δ (ref. 167 and 197)	493 (600 °C)	—	—	GDC	—	0.44 (700 °C)	1.08 (600 °C)
PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (ref. 168)	—	—	—	GDC	—	0.056 (600 °C)	2.90 (650 °C)
PrBa0.5Sr0.5CoFeO5+δ (ref. 232)	346 (600 °C)	—	20.9	GDC	0.077 (800 °C)		0.649 (850 °C)
PrBa0.8Ca0.2Co1.5Fe0.5O5+δ (ref. 169)	—	—	20.28	GDC	0.080 (600 °C)	—	1.89 (600 °C)
NdBaCo2O5+δ (ref. 170)	776 (600 °C)	—	21.5	LSGM	—	0.70 (700 °C)	0.55 (800 °C)
NdBaCoCuO5+δ (ref. 171)	99.2 (700 °C)	—	16.9	LSGM	—	0.28 (700 °C)	—
NdBaCoFeO5+δ (ref. 164)	71 (600 °C)	—	19.5	LSGM	—	0.119 (750 °C)	0.67 (800 °C)
NdBaCo2/3Fe2/3Cu2/3O5+δ (ref. 172)	92 (625 °C)	—	15.7	LSGM	—	0.023 (800 °C)	0.719 (800 °C)
NdBaCo1.6Ni0.4O5+δ (ref. 173)	654 (600 °C)	—	19.4	SDC	—	0.077 (700 °C)	0.714 (800 °C)
NdBaFe2O5+δ (ref. 170)	11.7 (600 °C)	—	18.3	LSGM	—	—	—
NdBa0.5Sr0.5Co2O5+δ (ref. 174)	2412 (600 °C)	—	—	GDC	—	0.112 (600 °C)	1.00 (600 °C)
PrBa0.5Sr0.5CoCuO5+δ (ref. 233)	221 (850 °C)	—	17.58	GDC	0.06 (800 °C)	—	0.521 (800 °C)
NdSrCo2O5+δ (ref. 174)	2420 (600 °C)	—	—	GDC	—	0.05 (700 °C)	—
SmBaCo2O5+δ (ref. 175 and 234)	434 (800 °C)	—	21.1	LSGM	—	0.054 (750 °C)	0.777 (800 °C)
SmBaCo1.6Fe0.4O5+δ (ref. 175)	342 (600 °C)	—	20.8	—	—	—	—
SmBaCo0.5Mn1.5O5+δ (ref. 176)	28 (600 °C)	—	18.7	LSGM	0.081 (900 °C)	—	0.377 (850 °C)
SmBaCo1.6Ni0.4O5+δ (ref. 173)	412 (600 °C)	—	16.6	SDC	—	0.11 (700 °C)	0.572 (800 °C)
SmBa0.5Sr0.5Co2O5+δ (ref. 235 and 236)	280 (900 °C)	—	21.9	GDC	0.092 (700 °C)	—	1.31 (800 °C)
SmSrCo2O5+δ (ref. 177 and 237)	2137 (600 °C)	—	22.7	GDC	—	0.28 (600 °C)	0.713 (600 °C)
SmSrCoMnO5+δ (ref. 237)	45.9 (700 °C)	—	13.7	—	—	—	—
YBaCo2O5+δ (ref. 179)	189 (600 °C)	—	16.3	LSGMC	—	0.11 (700 °C)	0.873 (800 °C)
YBaCo1.4Cu0.6O5+δ (ref. 178)	47.4 (700 °C)	—	14.7	LSGM	—	0.12 (700 °C)	0.642 (800 °C)
YBaCo1.8Fe0.2O5+δ (ref. 179)	128 (600 °C)	—	17.3	LSGMC	—	0.13	0.768 (800 °C)
YBa0.5Sr0.5Co2O5+δ (ref. 238)	371 (600 °C)	—	—	GDC	—	0.36 (700 °C)	—
GdBaCo2O5+δ (ref. 239)	375 (600 °C)	—	18.1	LSGM	0.34 (700 °C)	—	0.117 (750 °C)
GdBaCoCuO5+δ (ref. 171)	90.239 (800 °C)	—	16.3	LSGM	—	—	0.468 (800 °C)
GdBaCoFeO5+δ (ref. 170)	73.6 (600 °C)	—	18.8	LSGM	—	1.09 (700 °C)	0.450 (800 °C)
GdBaFeNiO5+δ (ref. 182)	—	—	14.7	SDC	—	0.922 (700 °C)	0.287 (800 °C)
GdBa0.5Sr0.5Co2O5+δ (ref. 183)	640 (600 °C)	—	—	GDC	—	0.13 (600 °C)	—
GdBa0.25Sr0.75CoCuO5+δ (ref. 171)	46.7 (800 °C)	—	16	LSGM	—	0.80 (700 °C)	0.53 (800 °C)
GdBa0.5Sr0.5CoFeO5+δ (ref. 183 and 184)	—	0.01 (600 °C)	—	GDC	—	0.067 (600 °C)	1.31 (600 °C)
GdSrCo2O5+δ (ref. 181)	1155 (600 °C)	—	18.8	LSGM	—	—	0.350 (800 °C)
Sr2FeTiO6−δ (ref. 240)	2.83 (600 °C)	—	16.8	SDC	—	0.204 (700 °C)	0.335 (800 °C)
Ba2CoMo0.5Nb0.5O5+δ (ref. 241)	1.2 (800 °C)	—	—	SDC	—	0.09 (750 °C)	—

---

A concern with A′-site alkaline earth cations, particularly Sr and Ba, is their preferential segregation towards the cathode surface^198–202 and their chemical instability in CO₂ as they have a high tendency to form carbonates on reaction with atmospheric CO₂.^41–43 The challenge this phenomenon presents is that surface enriched strontium and barium oxides are susceptible to reaction with other gaseous contaminants such as Cr and CO₂ to form secondary phases on the electrode surface which are electronically insulating and impede surface oxygen exchange, consequently hampering the conductivity and electrochemical performance of the cathodes. Segregation originates from the minimization of elastic energy or lattice strain caused by the size mismatch between the dopant and the host cations, driving the smaller or larger sized dopant to free surfaces or interfaces.²⁰³ Composition, temperature, electrochemical polarisation, and oxygen partial pressure have been observed to influence the degree of segregation of Ba and Sr to the surface of the cathode.^{201,204–206} To suppress surface segregation, Kwon et al.²⁰⁷ suggest selecting an A-site cation dopant such as Ca which exhibits a small size difference with the host cation. Additionally, they proposed a B-site cation with a small ionic radius such as Co or Ni as they discovered that the B-site ionic size is a major contributor to the segregation energetics. Xia et al.²⁰⁸ reported that co-doping Ca on the Ba and Pr sites in Pr_0.9Ca_0.1Ba_0.8Ca_0.2Co₂O_5+δ not only reduced the segregation of Ba to the electrode surface but reduced the material's TEC (18.5 × 10⁻⁶ K⁻¹) and improved its electrochemical performance (0.069 Ω cm²) and maximum power density (712 mW cm⁻² at 800 °C) as compared with PrBaCo₂O_5+δ with TEC, ASR and maximum power density values of 22.4 × 10⁻⁶ K⁻¹, 0.094 Ω cm², and 629 mW cm⁻², respectively. Chen et al.²⁰⁹ investigated the long-term stability of PrBa_0.8Ca_0.2Co₂O_5+δ (PBCC) in a Ni–BaZr_0.1Ce_0.7Y_0.1Yb_0.1O₃ anode supported cell with SDC electrolyte. The cell was run at a cell voltage of 0.7 V at 700 °C with humidified H₂ (∼3% H₂O) as fuel and air with ∼1% CO₂ as the oxidant. After ∼50 h of operation, the PBCC cathode showed stable power output with a degradation rate ∼1/24 of that of state-of-the-art LSCF cathode under the same operating conditions, indicating good tolerance to CO₂. Anjum et al.¹⁹⁹ observed a substantial reduction in Ba surface segregation in nanostructured GdBaCo₂O_5+δ (GBCO) (∼10 nm radius) as well as reduced impedance in comparison with the chemically synthesised bulk sized GBCO electrode. Thus, they proposed applying nano-structuring strategies to control surface cation segregation.

The most studied transition metals for B site substitutions include Co, Fe, Mn, Ni and Cu. Co exhibits a characteristic high catalytic activity making it a B site element choice for many compositions.¹⁶⁵ Furthermore, Co containing samples have demonstrated higher electrical conductivity values than other transition metal compositions. Zhang et al. describes the electronic conductivity of cobalt based double perovskite oxides to occur via the hopping of electrons along the Co⁴⁺–O²–Co³⁺ bonds within the perovskite structure.¹⁹⁰ The electronic conductivity of perovskite oxides with Co as the only B site ion generally depicts metallic behaviour within the temperature range of 300–900 °C. This phenomenon is explained by the decrease in the concentration of charge carrier accompanied by the loss of lattice oxygen, and the low to high spin transition of Co³⁺ ions with increasing temperature. Fe doping decreases the TEC, improves oxygen diffusivity and catalytic activity, and thermal stability.^169,183,210 However, electrical conductivity decreases with increasing Fe content. For such samples, several factors contribute to the observed reduction in electrical conductivity. Firstly, substituting the slightly larger Fe³⁺ for Co³⁺ alters the orbital configuration of the valence electrons in which the overlap between (Co,Fe)^3+/4+: 3d and O²⁻: 2p orbitals decreases resulting in lower electron delocalization and a consequent impediment to electron hopping. Moreover, because a charge compensation of Fe³⁺–Fe⁴⁺ preferentially occurs over Co³⁺–Co⁴⁺ and a low mobility of Fe ions, the electrical conductivity substantially decreases when Fe doping exceeds its percolation limit. Fe ions have such a lower mobility than Co ions that the reported mobility of electrical holes of LaFeO₃ is lower than that of LaCoO₃ by approximately three orders of magnitude.²¹¹ Choi et al.¹⁶⁸ investigated the synergetic effect of co-doping both A- and B-sites on the electrochemical properties of PrBaCo₂O_5+δ. Partially substituting Ba with Sr and Co with Fe in PrBaCo₂O_5+δ created oxygen vacancy (pore) channels in the [PrO] and [CoO] planes that significantly enhanced the oxygen ion diffusion and surface oxygen exchange resulting in a peak power density of ∼2.9 W cm⁻² at 650 °C. Similarly, doping Fe in PrBa_0.8Ca_0.2Co₂O_5+δ improved the catalytic activity of the cathode with a maximum power density of 1.89 W cm⁻² at 600 °C.¹⁶⁹ In both compositions, the stronger Fe–O bonding strength increased the mobile oxygen species in the Ln–O layer and thus, improved their catalytic activities. In terms of electrochemical performance, PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ and PrBa_0.5Ca_0.5Co_1.5Fe_0.5O_5+δ are very promising materials for IT-SOFC cathodes as they exhibit the highest maximum power density values.

Co containing perovskite oxides undergo chemical expansion with temperature increase, thus their TECs do not match the current IT-SOFC electrolytes. LnBaCo₂O_5+δ perovskites have been reported to possess TEC values in the range of 15–29 × 10⁻⁶ K⁻¹,¹⁵² while La_0.8Sr_0.2Ga_0.83Mg_0.17O_3−δ (LSGM) and Sm_0.2Ce_0.8O_1.95 (SDC) have TEC values of 11.4 × 10⁻⁶ K⁻¹ and 12.4 × 10⁻⁶ K⁻¹ respectively.^212,213 This large thermal expansion stems from the reduction of Co⁴⁺ ions to the larger Co³⁺ ions through the loss of lattice oxygen. Additionally, the spin state transition of Co³⁺ ions from low-spin Co³⁺(t⁶_2ge⁰_g) (0.54 Å) to the intermediate spin Co³⁺(t⁵_2ge¹_g) or high spin Co³⁺(t⁴_2ge²_g) (0.61 Å) state (Fig. 5) with increasing temperature strongly contributes to the TEC mismatch of Co containing cathodes. While several studies have attempted to develop compositions with compatible TECs, it has been observed that doping Sr on the Ba sites increases the TEC of the samples in the high temperature region.^181,214,215 Sr increases the concentration of Co³⁺ which undergoes spin-state transition within the 300–900 °C temperature range, hence an expansion of the lattice. Conversely, doping the A-site with Ca has been reported to improve the TEC of the composition.^208,216,217 Unlike Sr, doping with Ca facilitates the formation of Co⁴⁺ species which are known to exist in the low-spin state without transitioning to a higher spin state.²⁰⁸ Also, introducing a small amount of Ln- and/or Ba-deficiency has also been reported to slightly reduce the TEC.^163,218,219 Substituting with other transition metals on the B-site is another strategy to reduce the TECs of cobalt based perovskites. The effect of doping Mn, Fe, Ni, and Cu on the TEC of double perovskites has been investigated in several studies.^{170,172,173,176,183,220–222} Substitution of Fe for Co in LnBaCo_2−xFe_xO_5+δ (Ln = Nd and Gd) decreased their TECs from 21.5 × 10⁻⁶ and 19.9 × 10⁻⁶ K⁻¹ to 20.0 × 10⁻⁶ and 18.8 × 10⁻⁶ K⁻¹, respectively for x = 1.0 within the temperature range of 80–900 °C.¹⁷⁰ Thermal expansion in Fe doped perovskites is mitigated by the decrease in the concentration of oxygen vacancies as a result of a stronger Fe–O bond, and the formation of Fe³⁺ in place of the high spin Co³⁺ ions.¹⁸³ Cu doped GdBaCo₂O_5+δ and PrBaCo₂O_5+δ decreased from 18.3 × 10⁻⁶ and 24.1 × 10⁻⁶ K⁻¹ to 15.1 × 10⁻⁶ and 15.2 × 10⁻⁶ K⁻¹, respectively.^223,224 Ni substitution for Co in LnBaCo_1.6Ni_0.4O_5+δ (Ln = Pr, Nd, and Sm) exhibited a lower TEC than undoped LnBaCo₂O_5+δ (Ln = Pr, Nd, and Sm) by ∼4%, 8%, and 13%, respectively within 30–1000 °C.¹⁷³ TEC data reported in several studies (Table 4) show that doping Fe on Co sites reduces the TEC by a marginal quantity while Cu or Ni introduction considerably reduces the TEC. Co-doping Fe and Cu in GdBaCo_2/3Fe_2/3Ni_2/3O_5+δ yielded a TEC value of 14.6 × 10⁻⁶ K⁻¹ in comparison with GdBaCo₂O_5+δ (19.9 × 10⁻⁶ K⁻¹).²²²

Fig. 5 Electron energy levels of the Low Spin (LS), Intermediate Spin (IS) and High Spin (HS) states of Co³⁺ with 3d⁶ electron configuration.²²⁵

Another approach to minimize thermal expansion mismatch with electrolytes is incorporating electrolyte powders into the cathode compositions to form composite electrodes. Not only does this method decrease thermal expansion, it increases the triple phase boundary and consequently, active sites for the ORR. For example, SDC, which was introduced in LnBaCo₂O_5+δ in a 25–75 wt% composite, respectively reduced the TEC of the cathode from 21.5 × 10⁻⁶, 21.0 × 10⁻⁶, 19.1 × 10⁻⁶, and 17.6 × 10⁻⁶ K⁻¹ to 20.1 × 10⁻⁶, 18.5 × 10⁻⁶, 17.2 × 10⁻⁶, and 16.7 × 10⁻⁶ K⁻¹ for Ln = Pr, Nd, Sm, and Gd, between 30 and 1000 °C, respectively.²²⁶ Some studies have proposed the complete substitution of Co as a strategy to decrease the TEC of cobalt based cathodes. Thus, cobalt-free double perovskite oxides such as LaBaCuFeO_5+δ, SmBaCu₂O_5+δ, NdBaFe_2−xMn_xO_5+δ and GdBaFeNiO_5+δ have been investigated as potential IT-SOFC cathodes.^{182,220,227,228} The TEC of GdBaFeNiO_5+δ was reduced to 14.7 × 10⁻⁶ K⁻¹ as compared to GdBaCo₂O_5+δ (17.6 × 10⁻⁶ K⁻¹).¹⁸²Table 4 shows summary of the most important properties of selected IT-SOFC double perovskite cathodes.^{152,156–165,167,169–171,175–180,186,197,220,224,229–241}

Ruddlesden–Popper type oxides as IT-SOFC cathodes

Ruddlesden–Popper (RP) oxides are a family of perovskite related structures with a general formula of A_n+1B_nO_3n+1, alternatively written as AO(ABO₃)_n. Often described as a two-dimensional (2D) variation of three-dimensional (3D) perovskite oxides, their crystal structure consists of n perovskite ABO₃ layers stacked between AO rock-salt layers along the c-axis in an alternate arrangement (Fig. 6). The first member of the RP series A₂BO₄ (n = 1) exhibits a K₂NiF₄-type tetragonal structure, however, its ideal I4/mmm space group may transform due to a distortion of the BO₆ octahedra around the c-axis.²⁴² 2D perovskite layers are formed through the corners shared by the BO₆ octahedra, while the AO layers are located between perovskite layers along the c-axis.

Fig. 6 Schematic crystal structures of n = 1, 2 and 3 members of layered structure Ruddlesden–Popper type A_n+1B_nO_3n+1 perovskite oxides. The denotation of n represents the number of stacked perovskite ABO₃ layers separated by a rock salt AO layer.²⁵⁰

Layered K₂NiF₄-type Ln₂NiO_4+δ (Ln = La, Pr, Nd) nickelates and cuprates have garnered interest for exhibiting high oxygen ion diffusivity,^195,243 surface exchange kinetics, sufficient electrical conductivity²⁴⁴ and moderate thermal expansion coefficients,^245–247 earning them a space amongst promising alternative IT-SOFC cathode materials. For example, Boehm et al. demonstrated higher oxygen bulk diffusion (D* ∼ 4.15 × 10⁻⁷ cm² S⁻¹) and surface exchange coefficient (k* ∼ 7.57 × 10⁻⁶ cm S⁻¹) in RP Ln₂NiO_4+δ nickelates than conventional La_0.6Sr_0.4Fe_0.8Co_0.2O_3−δ (D* ∼ 5.40 × 10⁻⁹ cm S⁻¹, k* ∼ 9.26 × 10⁻⁸ cm S⁻¹) at 700 °C.²⁴⁸ While ABO₃ oxides are generally oxygen-deficient perovskites, RP oxides can accommodate oxygen interstitial defects in the AO layers,²⁴⁹ thus, their oxygen content can be hyper-stoichiometric as well as hypo-stoichiometric. These phenomena strongly influence the oxygen transport properties as oxygen ion migration can occur via mechanisms related to oxygen vacancies or interstitials.

Like single and double perovskites, cation substitutions influence the ionic and electronic conductivities, surface oxygen catalytic activity, and thermal expansion coefficients amongst other physical properties of RP phases. In addition to doping, the number of perovskite layers within the rock salt layers (AO) (Fig. 6)²⁵⁰ regulates their physical properties.²⁵¹ Generally, the electrical conductivity of RP perovskite oxides (n = 1) in air ranges up to a few hundred S cm⁻¹, depending on the temperature.^252,253 Despite the poor electrical conductivity (<100 S cm⁻¹) reported in some (n = 1) RP perovskite compositions, good catalytic activities have been observed. La_1.5Pr_0.5Ni_0.95−xCu_xAl_0.05O_4+δ (x = 0.1) exhibited a conductivity value of ∼30 S cm⁻¹ in air due to the reduction of charge carrier concentration, however its ASR and maximum power density were 0.04 Ω cm² and 530 mW cm⁻² at 800 °C, respectively.²⁵³ Its high catalytic activity was ascribed to Cu and Al doping which increased the oxygen vacancies, favouring the adsorption and transport of oxygen ions. Higher order RP (n = 2 and 3) MIEC oxides, however, are more electrically conducting and exhibit better electrochemical activities than the (n = 1) RP series due to the higher oxygen migration barrier of RP n = 1 than n = 2 and 3.^251,254,255 An extensive review on RP perovskite cathodes for SOFCs has been published by Ding et al.²⁵⁶

Other crystal structures for IT-SOFC cathodes

As stated in previous sections, perovskite or perovskite-related oxides containing Co have been widely investigated as candidate cathode materials owing to their high mixed ion–electronic conducting properties. However, their practical application has been hindered due to high TEC values when compared to conventional oxide-ion electrolytes. The resultant thermal expansion mismatch between the cathode and electrolyte becomes a major degradation issue during the thermal cycling of SOFCs as it can lead to delamination and cracks in the cathode layer. The increase in ionic radius due to the low-spin to high-spin transition of Co³⁺ ions (in the octahedral-site) is the cause of abnormally high TECs. Hence, Mn, Fe, Ni, and Cu have been partially substituted at the Co site in order to reduce the TEC values.^{4,60,257–260} These substitutions result in compositions that exhibit lower TEC values but the electrochemical performances also decrease when compared to parent Co-based compositions. Hence, researchers have investigated other crystal systems possessing matching TEC values with other cell components.

Swedenborgite-type RBaCo_4−xM_xO₇ (R = Y, Ca, In, Lu, Yb, Tm, Er, Ho, Dy; M = Co, Zn, Fe, Al, Ga) has shown potential as an oxygen storage material at low temperatures (200–400 °C). However, phase decomposition at elevated temperatures of 700–800 °C has prevented their application as SOFC cathodes.^261–267 Manthiram's group was the first research group that systematically studied the effect of various dopants on the phase stability and electrochemical performance of swedenborgite-type oxides as SOFC cathodes. In the RBa(Co,M)₄O₇ structure, Ba²⁺ and R ions adopt 12- and 6-fold oxygen coordination, respectively and the structure consists of corner-shared (Co, M)O₄ tetrahedra (Fig. 7a). Low TEC has been attributed to the presence of tetrahedral-site Co^2+/3+ ions which do not experience spin-state transitions at elevated temperatures as they are already in the high spin state.²⁶⁴ The low anisotropic TEC along the a-axis is the main contributor to the low bulk TECs as revealed by neutron diffraction studies in YBaCo₃ZnO_7+δ, Y_0.9In_0.1BaCo₃ZnO_7+δ, and Y_0.9In_0.1BaCo₃Zn_0.6Fe_0.4O_7+δ.²⁶⁸ The change in Co–O bond length in CoO₄ polyhedra was suppressed by doping with In, Zn, and Fe which resulted in a reduction in the anisotropic and bulk TECs.²⁶⁸

Fig. 7 Schematic illustration of the crystal structure of (a) YBaCo₄O₇. Reprinted with permission from ref. 275. Copyright 2006 American Chemical Society. (b) Y_3−xCa_xFe₅O_12−δ (Ia3̄d, cubic). Reprinted with permission from ref. 276. Copyright 2020 Elsevier.

The phase stabilities of RBaCo_4−xM_xO₇ series were assessed by long-term phase stability measurement by heating the samples at 600, 700, 800, and 900 °C for 50–120 h, and high-temperature X-ray diffraction (XRD) measurements. From Table 5, it can be seen that Zn substitution increased the phase stability at high temperatures for RBa(Co,M)₄O₇ (R = Y, Ca, In; M = Zn, Fe, Al).²⁶⁹ By looking into the decomposition products (BaCoO_3−δ and Co₃O₄) of the YBaCo₄O₇ sample, it was suggested that at elevated temperatures cobalt prefers to adopt octahedral coordination instead of tetrahedral coordination. In both BaCoO_3−δ and Co₃O₄ decomposition products, Co is in the octahedral coordination. As Zn²⁺ prefers the tetrahedral-site, the partial substitution of Zn²⁺ for Co^2+/3+ stabilised the YBaCo_4−xZn_xO₇ (x ≥ 1) phase with corner-shared CoO₄ tetrahedra.²⁶⁹ Similar to RBaCo_4−xM_xO₇, the high temperature phase stability of the Y_1−xCa_xBaCo_4−yZn_yO₇ system improved with increasing Zn content, while Ca contents ≥0.5 deteriorated the phase stability.²⁷⁰ Similar tests were also performed on the Y_0.5In_0.5BaCo_4−xZn_xO₇ (x = 1, 1.5, and 2) series.²⁷¹ Here, it was seen that employing a mixture of Y and In (50% each) promotes phase stability and overcomes the phase-decomposition problems due to the increased oxygen content and decreased lattice size.²⁷¹

Table 5 High temperature phase stability of selected metal oxides.^269–274a

Composition	Long-term stability test (120 h)
	800 °C	700 °C	600 °C
a ✗ = not stable, ✓ = stable.
YBaCo4O7 (ref. 269)	✗ (50 h)	✗ (50 h)	—
YBaCo4−xZnxO7; x = 0, 0.5, 1.0 ≤ x ≤ 2.0 (ref. 269)	—	—	—
	✓	✗
	✓ (50 h)	✓ (50 h)
YBaCo3ZnO7 (ref. 269)	✓ (50 h)	✓ (50 h)
YBaCo3ZnO7 (ref. 270)	✓	✗	✗
Y0.75Ca0.25BaCo2.5Zn1.5O7+δ (ref. 270)	✓	✓	✓
Y0.5Ca0.5BaCo2.25Zn1.75O7+δ (ref. 270)	✓	✓	✓
Y0.25Ca0.75BaCo2.5Zn1.5O7+δ (ref. 270)	✓	✓	✗
CaBaCo3ZnO7+δ (ref. 270)	✓	✗	✗
Y0.5In0.5BaCo3ZnO7±δ (ref. 271)	✓	✓	✓
Y0.5In0.5BaCo2.5ZnO7±δ (ref. 271)	✓	✓	✓
Y0.5In0.5BaCo2Zn2O7±δ (ref. 271)	✓	✓	✓
InBaCo3ZnO7±δ (ref. 271)	✗	✗	✓
YBaCo3.2Ga0.8O7+δ (ref. 272)	✓	✓	✓
YBaCo3.3Ga0.7O7+δ (ref. 272)	✓	✗	✗
InBaCo3.3Ga0.7O7+δ (ref. 272)	✗	✗	✗
CaBaCo3.3Ga0.7O7+δ (ref. 272)	✗	✗	✗
Y0.9In0.1BaCo3.3Ga0.7O7+δ (ref. 272)	✓	✓	✓
Y0.5In0.5BaCo3.5Ga0.5O7+δ (ref. 272)	✓	✓	✓
Y0.7In0.3BaCo3.3Ga0.7O7+δ (ref. 272)	✗	✗	✗
In0.7Ca0.3BaCo3.3Ga0.7O7+δ (ref. 272)	✗	✗	✗
YBaCo4−xAlxO7+δ (ref. 273)	✗	✗	✗
Y0.75Tb0.25BaCo3.2Ga0.8O7+δ (ref. 274)	✓	✓	✓

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In the Ga doped-YBaCo_4−yGa_yO_7+δ (y = 0.6–0.8) series, YBaCo_3.2Ga_0.8O_7+δ exhibited good stability in long-term studies suggesting the positive effect of Ga doping to reduce the temperature range of decomposition and improving the phase stability at 800 °C.²⁷² On the other hand, serious decomposition of InBaCo_3.3Ga_0.7O_7+δ into Co₃O₄, In₂O₃ and CaBa–Co_3.3Ga_0.7O_7+δ indicated again that the instability of Co³⁺ in the tetrahedral sites and its preference for octahedral coordination is the cause of phase instability.²⁷² The Y-doped Y_1−xIn_xBaCo_3.3Ga_0.7O_7+δ (x = 0.1–0.9) series also remains stable at high temperatures indicating that the synergistic effect of In and Y could also maximize the stability at a certain Ga content.²⁷² However, Y_1−xCa_xBaCo_3.3Ga_0.7O_7+δ and In_1−xCa_xBaCo_3.3Ga_0.7O_7+δ samples were not stable long-term, suggesting that there is no synergistic effect of In and Ca codopants.²⁷² In recent studies, doping and co-doping effects of trivalent cations (Al³⁺, Ga³⁺, and Fe³⁺) on the phase stability and electrochemical performance for the ORR have been reported.²⁷³ It was seen that Al based compositions, YBaCo_4−xAl_xO_7+δ showed severe decomposition above 700 °C.²⁷³ Among the trivalent dopants in the YBa(Co, Ga, Al, and Fe)₄O_7+δ series, the order of dopants towards the phase stabilization capability can be summarized as Ga³⁺ > Al³⁺ > Fe³⁺.²⁷³ Additional studies with the Tb doping showed that Tb has a relatively weaker stabilization capability compared to Y.²⁷⁴

However, owing to low oxygen permeation flux majority of the studies for the ORR were performed with GDC composites. For example, the non-composite YBaCo₃ZnO₇ cathode showed an ASR of 0.15 Ω cm² at 700 °C, whereas the YBaCo₃ZnO₇ + GDC composite cathode showed a lower ASR of 0.06 Ω cm² at 700 °C.²⁶⁹ Studies on various YBaCo₃ZnO₇ + GDC composite cathodes with various GDC contents showed that 50 : 50 wt% showed the lowest ASR values indicating that incorporation of GDC offers an extended TPB and oxide-ion bulk diffusion and thereby enhances the catalytic activity for the ORR.²⁷⁷ All the studies showed almost similar ASR values at 700 °C for composite cathodes (0.06–0.08 Ω cm²).^269,274 Stability against CO₂ was also investigated for (Y,Tb)Ba(Co,Ga)₄O_7+δ swedenborgite oxides,²⁷⁷ where it was seen that the ASRs of (Y,Tb)Ba(Co,Ga)₄O_7+δ–Gd-doped CeO₂ (GDC) composite cathodes only increased by ∼120% at 600 °C when exposed to 5% CO₂ in air,²⁷⁷ whereas literature studies have shown that the ASR of Co-containing perovskite oxides increases >500% when exposed to 5% CO₂ in air.²⁷⁷ The better CO₂ tolerance was attributed to the presence of low number of oxygen vacancies in (Y,Tb)Ba(Co,Ga)₄O_7+δ.²⁷⁷

Yttrium iron garnet, Y₃Fe₅O₁₂ (YIG) finds applications in ferrimagnetic oxide, microwave and magneto-optic devices.^278–280 Doping Y with Ca²⁺ increases specific oxygen permeability (10⁻¹¹ mol s⁻¹ cm⁻²).²⁸¹ The other advantage associated with YIG is low TEC values (10.6 × 10⁻⁶ K⁻¹).²⁸² Given these advantages few studies have employed doped garnets as SOFC cathodes. Zhong et al. showed that the Y_2.5Ca_0.5Fe₅O_12−δ (YCFO)–Ce_0.8Sm_0.2O_1.9 (SDC, 40 wt%) composite electrode cathode showed an ASR of 0.55 Ω cm² at 650 °C,²⁸² where oxygen ion diffusion, oxygen dissociative adsorption, and gas-phase diffusion were assigned as rate-limiting steps based on equivalent circuit modeling.²⁸² The maximum power density (MPD) of 438 mW cm⁻² with SDC electrolyte (40 μm) was seen at 650 °C.²⁸²

A recent study by Zhang et al. reported the systematic effect of Ca-doping on the electrical and electrochemical properties of Y_3−xCa_xFe₅O_12−δ (x = 0, 0.05, 0.1, 0.3, 0.5 and 0.7), where the x = 0.3 member exhibited the highest oxygen non-stoichiometry (δ = 0.19) and X-ray absorption spectroscopy (XAS) studies confirmed the formation of hole carriers () as a result of Ca doping. With an increase in Ca amount until x = 0.1, the electrical conductivity (1.58 S cm⁻¹ at 750 °C) increased and then decreased due to a decrease in the concentration of the charge carriers. The lowest ASR of 1 Ω cm² was seen for x = 0.3 garnet–LSGM composite electrode at 750 °C in air. pO₂ dependent ASR and impedance spectroscopy genetic programming (ISGP) analysis showed that oxygen dissociation and partial reduction of adsorbed oxygen molecule are the rate limiting steps for the ORR.²⁷⁶

Composite cathodes

Juhl et al. were one of the first researchers who correlated the performance and structure of composite SOFC cathodes.²⁸³ At an overvoltage of −50 mV an ASR of 0.5 Ω cm² at 850 °C was obtained for the LSM–YSZ composite cathode.²⁸³ With increase in the thickness of the composite cathode layer, the polarisation resistance decreased more prominently at 700 °C, indicating that the bulk of the LSM–YSZ (60 : 40) composite layer is active for the ORR.²⁸³ A uniform mixture of the LSM, YSZ, and pores was recognized as the key for considerable improvement of the composite layer which can lead to easy percolation of electrons, oxide ions and gas through the layer and hence extending the active TPB area.²⁸³

Leng et al. employed the LSCF–GDC (50–50 wt%) composite cathode with a power density of 625 mW cm⁻² at 600 °C.²⁸⁴ The effect of sintering temperature showed that the best performance at 600 °C was achieved for a sintering temperature of 975 °C.²⁸⁵ A high sintering temperature of 1100 °C resulted in a large area of dense regions with less micropores which significantly reduced the reaction area and increased the resistance of oxygen species diffusion along the surface of grains.²⁸⁵ On the other hand, the lower temperature sintered cathode was more porous, with a lot of macro- and micro-pores, which led to an increase in reaction area. However, reducing the temperature to 850 °C weakened the connection between agglomerated particles. This increased the resistance of bulk/surface diffusion of oxygen species including oxygen ion as well as electron transfer through the porous cathode.²⁸⁵

Chen et al. prepared a novel Pr₂NiO₄ (PNO)–Pr_0.2Ce_0.8O_1.9 (PCO) composite cathode through solid-state mixing and a modified sol–gel method,²⁸⁶ where the PNO–PCO composite cathode obtained by the sol–gel method exhibited better electrochemical performance due to uniform particle size distribution and porosity with GDC electrolyte (ASR = 0.09 Ω cm² at 800 °C), and the NiO–GDC/GDC/PNO–PCO single cell yielded an MPD of 0.57 W cm⁻² at 800 °C.²⁸⁶ Co-doped double perovskite-type cobaltite Pr_0.9Y_0.1BaCo_1.8Ni_0.2O_6−δ (PYBCN)–SDC (PYBCN–SDC) composite exhibited an ASR value of 0.045 Ω cm² at 800 °C where the lower ASR was attributed predominantly to a large concentration of O²⁻ vacancies in the cobaltite component of the composite.²⁸⁷ Jafari et al. showed an ASR at 0.008 Ω cm² (750 °C) for the La_0.6Ca_0.4Fe_0.8Ni_0.2O_3−δ–YSZ (LCFN–YSZ) composite.²⁸⁸

Zhao et al. prepared composite-cathode LSCF–GDC using the nanoparticle of GDC impregnated to the LSCF and obtained the lowest resistance of 0.07 Ω cm² at 600 °C.²⁸⁹ Liu et al. performed long term studies on the LSCF–GDC composite cathode and showed that ASR increased from 0.38 Ω cm² to 0.83 Ω cm² after testing at 750 °C for 500 h. It was also shown that the degradation rate of the LSCF conventional cathode was higher when compared with composite-cathode LSCF–SDC.²⁹⁰ Xi et al. impregnated Sm_0.5Sr_0.5CoO_3−δ (SSC) into PrBaCo₂O_5+δ (PBC) and obtained an ASR of 0.16 Ω cm² and power density of 385 mW cm⁻² at 700 °C.²⁹¹ The infiltrated LCFN–SDC (70 : 30) composite cathode showed an ASR of 0.15 Ω cm² at 800 °C where the improvement was attributed to enhanced activity for surface oxygen dissociation and diffusion processes achieved due to the specific electrode architecture by the nano SDC decorated on the LCFN backbone.²⁹² A single perovskite oxide Sm_0.5Sr_0.5CoO_3−δ (SSC) with high ORR activity was combined with MIEC SmBaCo₂O_5+δ (SBC) to exhibit an ASR of 0.021 Ω cm² at 750 °C.²⁹³

In the field of proton conducting SOFC cathodes, composite cathode materials are mainly divided into proton-blocking composite cathodes (PBCCs) and proton-conducting composite cathodes (PCCCs).²⁹⁴ In PBCC, electrochemical reactions are mainly restricted at the cathode–electrolyte interface, as dissociated oxygen ions are transferred along the surface of the cathode or through the bulk of the cathode to TPBs. Whereas in PCCCs, transport of all three charge carriers (oxygen vacancies, electronic defects and protons) occurs simultaneously resulting in enhanced active area for electrochemical reactions.²⁹⁴ Simply, it is expected that PCCCs should exhibit better electrochemical performance than the PBCCs owing to the presence of more TPBs. However, some experiments show the opposite trend. As water is generated at the cathode side, the PCCC will adsorb more water and hence there will be a reduction in active TPBs for the ORR; on the other hand, PBCC will show better performance even though it has less TPBs.

When comparing two types of PBCC with BaZr_0.1Ce_0.7Y_0.2O_3−δ (BZCY) electrolyte, La₂NiO_4+δ–LaNi_0.6Fe_0.4O_3−δ (LNO–LNF) showed a lower ASR of 0.103 Ω cm² with an MPD of 490 mW cm⁻² at 700 °C than Sm_0.2Ce_0.8O_2−δ–LaNi_0.6Fe_0.4O_3−δ (SDC–LNF).²⁹⁴ Additionally, introducing an anode functional layer (AFL) between the anode and electrolyte increased the power outputs to 708 mW cm⁻², at 700 °C for the LNO–LNF cathode. The LNO–LNF cathode exhibited better performance than SDC–LNF due to its oxygen transport mechanism which occurs through interstitial oxygen defects (Fig. 8).²⁹⁴ Chen et al. generalised the percolation theory for typical H⁺–SOFC composite cathodes with e⁻, O²⁻ and H⁺ mixed conducting properties based on the (LSCF–SDC–BZCY) composite cathode.²⁹⁵ Duan et al. employed a proton-, oxygen-ion-, and electron–hole-conducting BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ as a composite cathode and obtained a high power density of 455 W cm⁻² for 40 wt% BCZYYb + 60 wt% NiO|BCZYYb +1.0 wt% NiO|BCZY63 + BCFZY0.1, and 405 W cm⁻² for 40 wt% BCZYYb + 60 wt% NiO|BCZYYb +1.0 wt% NiO|BCZY63 + BCFZY0.1 cell at 500 °C.²⁹⁶ Dai et al. employed a one step co-firing process to prepare SSC–BZY and SSC–BaCe_0.7Zr_0.1Y_0.2O_3−δ (BCZY) composite-cathode,²⁹⁷ where SSC–BZY showed an ASR of 0.3 Ω cm² and SSC–BZCY showed an ASR of 0.58 Ω cm² at 650 °C.²⁹⁷ The lower ASR of SSC–BZY was attributed to the highly porous microstructure, which increased the rate of gas diffusion at the cathode.

Fig. 8 Schematic diagrams of the overall ORR at (a) the LNO–LNF cathode (b) SDC–LNF cathode with BZCY electrolyte, and an ORR model at the cathode (c) LNO–LNF/electrolyte (BZCY) interface and (d) SDC–LNF/electrolyte (BZCY) interface. Reproduced from ref. 294 with permission from the Royal Society of Chemistry.

Mechanism of the oxygen reduction reaction (ORR)

The overall ORR can be generally expressed by an overall equation as shown in Table 6 ^{283,285,298–301} and Fig. 9. The ORR consists of multielementary reactions such as adsorption of oxygen molecule onto the surface, dissociation of the oxygen molecule and diffusion of the adsorbed species, subsequent reduction of adsorbed oxygen species and incorporation into the cathode/electrolyte lattice.²⁹⁸ The rates of these reactions determine the cathode's ultimate electrochemical performance. To understand the reason behind the different electrochemical behaviours of various compositions, understanding the rate determining step (RDS) of the ORR is very important.

Table 6 pO₂ dependency of ASR revealing the process associated with different reactions involved in the ORR^{283,285,298–301}

Process	n	Equation
Overall	ASR = ASRo(pO2)^−n
Adsorption of oxygen molecules^298,299	1	O2(g) ⇆ O2,ads
Transfer of electrons^299	0.39
Ionization of atomic oxygen and CT reaction at TPB^300	0.28–0.36
Charge-transfer at TBP^301	0.24 to 0.32
Oxygen surface diffusion of dissociative adsorbed oxygen at the TPB (La0.75Sr0.25)0.9MnO3–8YSZ (50 : 50)^283	0.5	O2,ads(g) ⇆ 2Oads
Oxygen surface diffusion of dissociative adsorbed oxygen at the TPB (La0.75Sr0.25)0.9MnO3–8YSZ (50 : 50)^283
Charge transfer reaction^285	0.25
O^2− transfer from the TPB to the electrolyte^285	0

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Fig. 9 Schematic representation of the variation of ASRs as a function of pO₂.

Based on the conducting mechanism and the pathways for the ORR, cathode materials can be categorized into two groups: (i) pure electronic conducting materials, and (ii) mixed ionic-electronic conductors (MIECs).^{6,8,119,302,303} In the first group, after the oxygen molecule has been adsorbed and dissociates on the perovskite surface, it migrates to the TPBs through surface diffusion where oxide ions form and incorporate into the electrolyte by electron transfer. Fig. 10a schematically shows the oxygen pathway in a pure electronic conductive catalyst. The length of TPBs plays an important role in controlling the catalytic activity of the electrode. In this case, a porous electrode is required to provide more TPB sites for the ORR.^{6,8,302–304} The second group consists of materials showing both electronic and ionic conductivity towards the ORR. As a result, the adsorbed oxygen can be transferred to the electrolyte via both surface and bulk diffusion not limited to the TPBs as illustrated in Fig. 10b. MIECs are particularly attractive for application in IT-SOFCs where catalytic activity is required in the lower temperature range.³⁰⁵

Fig. 10 Schematic of the possible pathways for the oxygen reduction reaction (ORR) in (a) pure electronic conductor and (b) mixed-ionic electronic conductor (MIEC).

Each elementary step of the overall ORR occurring at the MIEC cathodes has a specific relationship with the pO₂ as shown in Table 6.^{283,285,298–301} For example, the double perovskite type Y_1−xCa_xBaCo₂O_5+δ (YCBC)/LSGM/YCBC symmetrical cell showed a dependence value of 0.5 at temperatures 700 °C to 800 °C, indicating that dissociation of molecular oxygen into atomic oxygen is the RDS at these temperatures.²¹⁶ The pO₂ dependence of ASR for LSM-i-ESB (LSM-infiltrated ESB (Bi_0.8Er_0.2)₂O₃) at 650 °C shows a dependence value of 0.1 for ASR_HF, and dependence of 0.7 for ASR_LF.³⁰⁶ This indicates that ASR_LF is related to surface chemical reactions, whereas ASR_HF is related to oxygen ion transport between solid phases. In the case of Nd_0.75Sr_0.25Co_0.8Fe_0.2O_3−δ (NSCF) + LSGM symmetrical cells at low pO₂ range (<0.1 atm), ASR_HF showed dependence of 1, and ASR_LF showed dependence of 0.24 at 700 °C. At high pO₂ range (>0.1 atm), ASR_HF showed dependence of 0.58; ASR_LF showed dependence of 0 at 700 °C,³⁰⁷ indicating that in the low pO₂ range, the charge-transfer reaction dominates the ORR, whereas surface diffusion (of dissociative adsorbed oxygen) dominates the ORR in the higher pO₂ region.

Distribution of relaxation times

From a mechanistic point of view, the ORR involves various subprocess where ASR values are affected by both fundamental material properties (ionic and electronic conductivity, oxygen surface exchange and bulk diffusion rates) and synthesis parameters (particle size, surface morphology of the grains, porosity, tortuosity, and the interface between the electrode and electrolyte). Conventionally, the impedance data are analysed through equivalent circuit modelling (ECM), where a viable equivalent circuit (EC) is fitted to the measurement dataset.^308–310 In ECM, known electrical analogues, such as resistance, capacitance and inductance are used to build the EC, along with more complex functions like semi-infinite and bounded diffusion (Warburg type) elements, and Gerischer element (GE).^308,311 However, deconvolution of AC impedance data becomes difficult when processes with similar relaxation times are present, especially in the case of the ORR which involves several elementary reactions (Fig. 11).

Fig. 11 Nyquist plot showing the semicircle deconvolution of simulated data with different relaxation times (RC).

Alternatively, the impedance data can be transformed into a distribution function of relaxation times (DFRT).^312–323 DFRT shows data as peaks on a log(τ) axis and each peak corresponds to specific electrochemical process. DFRT does not involve preconceived notions and is hence model free. By looking into the trends of peak position and height as a function of temperature, partial pressure and/or polarization, information on the electrochemical processes can be easily deducted and visualised. The DFRT, G(τ), can be obtained by solving the following expression:

where Z(ω) is the dataset, R_∞ is the high frequency cut-off resistance, and R_pol is the polarization resistance or overall resistance of the dispersion. The time constant, τ, is the inverse of the frequency: τ = (2πf)⁻¹ = ω⁻¹.

Marshenya et al. studied the impedance data under OCV conditions for Pr_0.9Y_0.1BaCo_1.8Ni_0.2O_6−δ–Ce_0.8Sm_0.2O_1.9 (PYBCN–SDC) composite cathode through DRT (Fig. 12).²⁸⁷ The DRT plot clearly showed that at lower temperature the ORR is dominated by a single electrochemical process and with increase in temperature, four different additional processes appear (Fig. 12e). The additional peaks at high temperatures were attributed to the presence of impurity phases seen in high temperature mixtures of PYBCN–SDC powders. Although electrochemical processes responsible for each peak were not explained, it was suggested that the impurity peaks might have affected the electrochemical parameters of oxygen exchange between the cathode and ambient atmosphere.

Fig. 12 (a–d) Electrochemical impedance spectra, (e) DRT plots at different temperatures for the symmetrical cell with PYBCN–SDC (70–30 wt%) composite electrodes at different temperatures in the air. Reprinted from ref. 287. Copyright 2019, with permission from Elsevier.

Fig. 13a and b show the impedance plots under OCV conditions and Fig. 13c and d show the DRT plots for symmetrical cells of LaBa_0.5Sr_0.5Co_1.5Fe_0.5O_5−δ (LBSCF)–GDC and NdBa_0.5Sr_0.5Co_1.5Fe_0.5O_5−δ (NBSCF)–GDC composite cathodes at different temperatures in air.³²⁴ The authors first fitted the impedance plot with ECM and then employed DRT to deconvolute data to further understand individual electrochemical processes. The authors assigned oxygen ion charge transfer from the electrolyte to the cathode at TPB to high frequency (>10³ Hz) peaks, surface exchange or ion transfer at the cathode to IF (1–10³ Hz) peaks, and the gas diffusion process was attributed to LF (10²–1 Hz) peaks. By analysing the DRT plots, the authors argued that since the integral areas for HF and IF peaks of NBSCF were smaller than those of LBSCF (Fig. 13c and d), NBSCF possessed higher oxygen surface exchange and diffusion ability. Although not clearly seen, the authors mentioned that since LBSCF DRT peaks showed slightly larger temperature dependence, LF peaks attributed to the oxygen diffusion process were termed RDS. It is important to mention that individual peaks were assigned to different electrochemical processes by referencing literature studies.

Fig. 13 Electrochemical impedance spectra of (a) LBSCF and (b) NBSCF symmetrical cells under OCV conditions at different temperatures in air. DRT analysis of ASRs for LBSCF and NBSCF cathodes at (c) 700 °C and (d) 650 °C. Reprinted from ref. 324. Copyright 2021, with permission from Elsevier.

Wei et al.³²⁵ employed DRT to distinguish the contributions of different polarization processes of anode-supported button cells with Ba_0.9Co_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (B9CFZY) and B9CFZY–BaZr_0.1Ce_0.7Y_0.2O_3−δ (BZCY) cathodes. Fig. 14 shows the fitted impedance data of anode supported cells under OCV conditions with B9CFZY and B9CFZY–BZCY cathodes, although ECM used for fitting the experimental data was not specified. Fig. 15 shows the DRT plot at 700–550 °C, where at lower temperature four peaks were seen, and at 700 °C three peaks were seen. The peaks were labelled as P1, P1_add, P2, and P3. Comparing the DRT plots in Fig. 14a and b, it can be seen that P1 peaks were similar in both cases and were assigned to hydrogen charge transfer in the anode. P2 and P3 were assigned to oxide ion diffusion to TPBs or active sites in the cathode, and oxygen gas adsorption/dissociation, while referencing literature studies. In the B9CFZY–BZCY cell, P2 and P3 were smaller than the corresponding peaks for the cell with the B9CFZY cathode, indicating the positive effect of adding BZCY to the B9CFZY cathode which boosts oxygen gas adsorption, dissociation and transfer. As P3 was the major contributor to polarization resistance, the RDS was assigned to oxygen species involved in the reaction. P1 seen at lower temperatures was assigned to the incorporation and transfer of O²⁻ in the lattice. The assignment of individual electrochemical processes to peaks in DRT plots was based on other studies, where impedance spectra of the cell were further characterized and analysed as a function of anodic and cathodic gas composition.³²⁶

Fig. 14 The fitted impedance spectra of the cells under OCV conditions with (a) B9CFZY and (b) B9CFZY–BZCY cathodes. Reprinted with permission from ref. 325. Copyright 2019, with permission from Elsevier.

Fig. 15 DRT analysis of the impedance spectrum data under OCV conditions for the anode supported cells with (a) B9CFZY and (b) B9CFZY–BZCY cathodes, and the values of polarization resistance corresponding to the different peaks for the anode supported cells with (c) B9CFZY and (d) B9CFZY–BZCY cathodes. Reprinted with permission from ref. 325. Copyright 2019, with permission from Elsevier.

Almar et al. investigated the ORR of the BSCF/GDC symmetrical cell cathode under OCV conditions, which exhibits fast oxygen-exchange kinetics leading to the impedance spectra showing a semicircle for the Gerischer process instead of the typical tear drop shape^327,328 by DRT (Fig. 16 and 17). P1 showed low thermal deactivation and pO₂ dependence of 0.98, and hence was associated with molecular oxygen diffusion within the cathode setup, the contacting gold meshes and the porous cathode. P2 showed thermal activation and pO₂ dependence of 0.66 and hence was associated with the surface-exchange reaction. P3 also showed thermal activation with pO₂ dependency of 0.09, with a capacitance of 0.05 to 0.08 F cm⁻² from 600 to 900 °C associated with interfacial capacitances. Hence P3 was attributed to oxide transfer losses across the cathode/electrolyte interface. P4 also showed thermal activation but was pO₂ independent and hence was attributed to electronic current losses between the electrode and the current collector (gold mesh).

Fig. 16 Electrochemical impedance plots for the BSCF/GDC/BSCF symmetrical cell under OCV conditions with in situ sintered electrodes from 900 to 600 °C at a constant pO₂ of 0.21 atm: (a) impedance spectra (ohmic losses were subtracted for clarity reasons) and (b) corresponding DRTs. Reprinted with permission from ref. 328. Copyright 2017, with permission from The Electrochemical Society.

Fig. 17 Electrochemical impedance plots for the BSCF/GDC/BSCF symmetrical cell under OCV conditions with in situ sintered electrodes at 700 °C in the pO₂ range from 0.02 to 1 atm, (a) impedance spectra (ohmic losses were subtracted for clarity reasons) and (b) corresponding DRTs. Reprinted with permission from ref. 328. Copyright 2017, with permission from The Electrochemical Society.

Mroziński et al. employed the DRT method to validate the ECM fit for Sr_0.86Ti_0.65Fe_0.35O₃ (STF35)/GDC symmetrical cells under OCV conditions.³²⁹Fig. 18a and b show DRT plots at different pO₂ (10%, 1%, and 0.1%), where three peaks were seen at HF, MF, and LFs depending on the pO₂. At low pO₂ (0.1%) additional contribution at LFs was seen, and the HF peak was ascribed to the Gerischer process.³²⁹ The temperature dependent DRT plot at 0.1% oxygen content in Fig. 18b shows that LF contribution is present at all temperatures. From these observations, ECM with the Gerischer element was proposed and the chi-squared parameter was mostly <10⁻⁵ (Fig. 18c and d). Fitting with different ECMs gave bad fittings along with higher chi-squared values, where adsorption of oxygen species was determined as RDS after analysing the dependence of each peak on pO₂ and temperature, and calculating the activation energy and capacitance values.³²⁹

Fig. 18 DRT plots of a symmetrical STF35 electrode under OCV conditions as a function of (a) pO₂ at 800 °C, (b) temperature at 0.1% O₂. Impedance spectra fitted with an ECM at (c) 800 °C and (d) 700 °C in 0.1% O₂. Reprinted with permission from ref. 329. Copyright 2019, with permission from Elsevier.

DRT analysis of the pO₂ dependence study for La_0.85Sr_0.15MnO_3±δ (LSM) infiltrated (Bi_0.8Er_0.2)₂O₃ (ESB)/GDC symmetrical cells under OCV conditions is shown in Fig. 19, where R1 (red) indicates the process of ion transport, R2 (blue) indicates surface chemical reactions, and R3 (dark yellow) indicates gas diffusion.³⁰⁶ The intermediate peak R2 associated with surface chemical reactions shows strong dependency on pO₂ and its intensity in DRT plots also shows strong correlation with pO₂ and is considered the rate-limiting step. The above-mentioned examples show that DRT analysis has been successfully employed to deconvolute impedance plots for cases with similar relaxation times, resulting in better understanding of individual electrochemical processes occurring in SOFC cathodes. It was also seen in various studies that DRT plots served as a complementary tool for ECM fitting of the impedance data.

Fig. 19 Electrochemical performance of LSM-i-ESB/GDC symmetrical cells under OCV conditions. (a) pO₂ dependence of impedance spectra at 650 °C, (b) DRT analysis of impedance spectra under different pO₂ at 650 °C, (c) variation of ASR with pO₂ at 650 °C, (d) temperature dependence of impedance spectra in synthetic air, (e) DRT analysis of impedance spectra for symmetrical cells calcined at 650 °C (closed symbols) and 800 °C (open symbols), and (f) Arrhenius plot of cathode ASR. Reprinted with permission from ref. 306. Copyright 2018, with permission from American Chemical Society.

Summary

The evolution of IT-SOFC cathodes has undergone tremendous progress in recent years, particularly towards designing alternative cathode materials that are excellent electrocatalysts, chemically and thermally stable, compatible with IT-SOFC electrolytes, and demonstrate improved electrochemical properties. We have discussed recent advances in the development of IT-SOFC cathodes with respect to material developments and interfacial engineering. LSCF, one of the most investigated MIECs, has demonstrated good ionic and electronic conductivities suitable for IT-SOFC cathodes. However, under the operating conditions, its TEC is incompatible with IT-SOFC electrolytes, and it suffers from surface segregation of Sr and reactivity with other contaminants in the cell. To prevent contaminant-poisoning of LSCF, surface functionalization routes such as infiltration have been proposed. Additionally, double perovskite oxides, especially Co-containing perovskites, have been investigated as cathodes for IT-SOFCs and have demonstrated higher surface exchange kinetics and diffusion coefficient of oxygen ions and higher electronic conductivities in comparison with ABO₃ disordered perovskite oxides. Additionally, mixing high performing cathode materials with electrolytes in a composite ratio further improves the electrochemical performance and long-term stability of the cathode and increases the power density of the cell. The reduction of oxygen occurs through several intermediate reactions whose kinetics are strongly influenced by the properties of the material, synthesis parameters and operating conditions. By distinguishing the electrode processes and the corresponding relaxation time, DFRT/DRT can be employed to comprehend the intricate ORR process in IT-SOFCs. For example, the DRT analysis on the dependence of LSM-i-ESB on pO₂ suggests that pO₂ is the rate-limiting step in this process.

Author contributions

V. T. planned the review paper. S. A. wrote the section on single perovskite oxides and ORR mechanisms, A. N. wrote the sections on double perovskites and layered Ruddlesden–Popper perovskites and K. S. contributed to the sections on other crystal structures, composite cathodes, and distribution of relaxation times. All authors commented on the paper.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors would like to thank the Alberta Innovates for supporting this work. One of us (V. T.) would like to dedicate this paper to Professor John Kilner.

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