Review on Ruddlesden–Popper perovskites as cathode for solid oxide fuel cells

There are currently three recognized oxygen ion diffusion mechanisms: the direct interstitial mechanism, interstitialcy mechanism and the vacancy mechanism of diffusion [60]. The direct interstitial mechanism refers to the migration of interstitial ions to the adjacent interstitial site without causing any permanent displacement of the other ions [60]. Frayret et al [52] studied the direct diffusion mechanism of La₂NiO₄ by studying the electronic structure using density functional theory (DFT), he only considered the interstitial oxygen between adjacent holes in the La₂O_2.25 layer and concluded that direct diffusion mechanism seems to be determined by the lattice relaxation and oxygen polarizability. Different from the direct interstitial mechanism, interstitialcy mechanism might include the interstitial ions further transfer away from the equilibrium sites [57]. For example, Chroneos et al [61] used molecular dynamic simulation to predict the oxygen diffusion process of La₂NiO₄, as shown in the figure 3, for predicting interstitialcy mechanism, the oxygen interstitial replaces a top oxygen ion from the NiO₆ octahedron, and the oxygen ion transfers to the adjacent oxygen interstitial position. Minervini et al [62] proposed a 'push–pull' mechanism to describe oxygen diffusion through atomic simulation calculations and suggested the generation of Frenkel anion defects (oxygen vacancies and interstitial oxygen) leads to an increase in the extra oxygen content. The oxygen diffusion in perovskites can also be illustrated by the vacancy mechanism of diffusion, oxygen ion diffuses by jumping to a neighboring vacancy [63]. However, vacancy diffusion of the RP-structure is different from perovskite oxide. For example, Parfitt et al [64] studied the oxygen transport in Pr₂NiO_4+δ through molecular dynamics calculations and found that the oxygen diffusion in Pr₂NiO_4+δ was highly anisotropic, occurred almost exclusively through the gap mechanism of the a–b plane, while the vacancy transport mechanism along the c axis was limited.

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To further understanding the mechanism of the oxygen diffusion, the oxygen migration species and migration paths needs to be considered [65]. Oxygen migration jumps by the interstitialcy (push–pull) mechanism, where interstitial oxygen can effectively pass through the peroxide interstitial state [66]. Both neutron scattering experiment [67] and molecular dynamics calculation [61] proved that oxygen ions in Ln₂NiO₄ (Ln = La, Pr) perovskite are mainly transmitting along the a–b plane through the interstitial process. The oxygen transport of RP materials is mainly accomplished through interstitial and interstitial mediated diffusion mechanisms, while vacancies play a smaller role in diffusion, even for materials with lower stoichiometry [59, 68, 69]. For example, Huan et al [70] further studied the ORR activity of RP materials, and found that it is the concentration of interstitial oxygen ions, the mobility of interstitial oxygen ions and lattice oxygen that governs high ORR activity rather than the oxygen vacancies. Minervini et al [62] had predicted in La₂NiO₄ that both monovalent and divalent oxygen ions are involved in interstitial migration by atomic calculations. Buttery et al [71] confirmed interstitial oxygen migration in Ln₂NiO₄ including ${\text{O}}_2^{2 - }$species migration. Xu et al [72] found the two oxygen interstitial diffusion mechanisms in La_2−xSr_xNiO_4+δ materials by DFT. In addition to the proposed interstitial-mediated mechanism, another type of oxide and peroxidation was also discovered. Figure 4 shows the oxide–oxide, oxide–peroxide diffusion mechanism. In the transition state of the oxide–oxide mechanism, the oxide interstitial (Oint²⁻)atom 'kicks out' the oxygen at the top of the bond with the Ni atom, resulting in the temporary formation of a vacancy. However, the peroxide interstitial (Oint¹⁻) bonds with a top oxygen atom in the rock salt layer to form an ${\text{O}}_2^{2 - }$ dumbbell, without generating oxygen vacancies.

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Moreover, the transport of oxygen has a high degree of anisotropy, where the diffusion of oxygen gaps in the a–b plane is generally higher than that in the c-direction by at least an order of magnitude [74–76]. For example, Xia [73] used DFT method to explore two typical migration paths of oxygen ions in S₃Fe₂O_7−δ and calculated the energy barriers for different oxygen ion migration directions, found that the energy barrier for oxygen ion migration in the [010] direction is lower than that in the [001] direction, confirming that oxygen ion migration occurs mainly in the a–b plane. As demonstrated in figure 5, the maximum migration energy barrier for the oxygen migration along path I is 1.40 eV, which is along the [001] direction: T1→2, T2→3, T3→4 and T4→3, while lower energy barriers of 0.90 eV for oxygen diffusion with path II (T1→2 and T2→5) and path III (T6→7 and T7→8) in [010] direction. Zhang et al [53] comparatively studied the interstitial oxygen transport of RP materials with n = 1 and n = 2, and calculated the relative energy of diffusion on the a–b plane using the CI-NEB (see figure 6), the oxygen migration barrier in n = 2 is about 1.46 eV, which is significantly lower than 1.74 eV in n = 1. For LNO materials, the oxygen ion mobility increases with the number of layers, which is consistent with previous report [77, 78].

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The migration of protons in perovskite or perovskite-like materials follows the two-step Grotthuss mechanism which is accepted by most people [66, 79–82]. The main characteristics of proton defect transmission includes rotational diffusion and octahedral jumping (protons jump from one oxygen ion to another adjacent oxygen ion in the same octahedron), that is, protons only show long-distance diffusion, while oxygen remains in its crystal position [79]. For example, Hermet et al [80] revealed the proton transport of BaCeO₃ perovskite that there are two processes of proton transport: reorientation and transfer (or hopping), protons may jump directly from one octahedron to another (jumping between the octahedrons rather than the intraoctahedral hoppings). By using theoretical (first-principle) and experimental methods, Wang et al [81] investigated two possible proton migration processes, and found that highest proton migration barrier was about 0.63 eV, which is in agreement with the experimental findings. It should be pointed out that the exact mechanism of proton transport in RP-type perovskite has not been determined by research activities [82]. Researchers generally use the proton transport mechanism in perovskite materials to explain the proton transport in RP-type materials. Proton diffuse to RP-structure Sr₃Fe₂O_7−δ (SFO) surface via the combination of two typical paths: the proton intra octahedral hopping and the proton reorientation around O ions, as illustrated in figures 7(a) and (b), respectively [65]. The transfer of protons across the rock-salt layer (O6 to O5) is the most difficult process with the highest energy barrier of 0.67 eV.

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The possible location of proton defect formation in RP perovskite oxide was demonstrated by DFT [83]. As shown in figure 8, Sr₃Fe₂O_7−δ proton defects may form on three different oxygen atoms, d1, d2 and d3, respectively, and the lowest proton formation energy at d3 indicates that proton defects are more likely to be formed at d3 [83]. Proton transfer is achieved by the two-step Grotthuss diffusion mechanism [84, 85], which involves the rotation of the proton around an oxygen site (R) and proton transfer between adjacent oxygen ions (T) [86]. For example, Wang et al [83] studied SFO cathode and found that the proton migrates along the direction of [001] and [010] by means of rotation (R) and transfer (T) as shown in figures 9(a) and (b). It reveals that proton diffusion in perovskites along the direction [010] has a greater advantage than through the SrO layer [001]. An overall proton migration pathway along the [001] and [010] direction can be summarized as the following process: T12/R23/R34/T45/T56 and T12/T23/R34/R45. The [001] direction offer the highest migration barrier of 0.62 eV higher than the migration of [010] direction (0.28 eV). It is generally accepted that a smaller energy barrier along the path favors proton diffusion. Chen et al [87] performed by comprehensive calculation that the proton migration in Sr₃Fe₂O_7−δ is more likely to occur in the [010] layer. In addition, Zhang et al [53] disclosed the proton transfer of RP series LNO (n = 1 and 2) materials, and it turned out that the highest migration barrier (1.66 eV) in n = 1 was greater than that in n = 2 (1.29 eV), which led to the faster proton diffusivity of LNO materials with n = 2, which is caused by the difference of the effective charge of proton, in which the average value of proton n = 2 (0.522e) is higher than that in n = 1 (0.459e).

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From the experimental perspective of exploring the diffusion mechanism, electrical conductivity relaxation (ECR) methods are often used to determine chemical oxygen bulk diffusion coefficient (D_Chem) and chemical surface exchange coefficient (k_Chem). For example, Hu et al [44] studied the oxygen transport performance of La_1.9−xBi_xSr_0.1CuO_4−δ using ECR technology, it was shown that LBSC2 (x = 0.2) has much higher k_Chem and D_Chem compared with LBSC1 (x = 0.1). At 800 °C, k_Chem for LBSC2 is 7.7 × 10⁻⁴ cm s⁻¹, about 15% higher than that of LBSC1, whereas D_Chem is 1.5 × 10⁻⁵ cm² s⁻¹, twice as that for LBSC1. He ascribed the enhanced oxygen transport performance to the extra oxygen vacancies formation and higher active oxygen concentrations by Bi doping. Furthermore, the faster apparent values of D_Chem and k_Chem achieved with increased time of annealing was studied by Saher et al [88], it was found that the two parameters increased with the increase of the annealing time and tended to be stable after 80 h. The apparent values of D_Chem and k_Chem were about one order of magnitude higher than the initial value. ¹⁸O/¹⁶O isotope exchange using secondary ion mass spectrometry and temperature programmed isotope exchange (TPIE) are the latest methods for measuring D_Chem and k_Chem [89]. Gu et al [90] adopted the isotope oxygen exchange model to explore the effectiveness of the oxygen exchange process of La₂NiO₄, they confirmed O₂ decomposition as rate-determining step. Sadykov et al confirmed that oxygen diffusion is the rate-determining step the in Pr_2−xCa_xNiO_4+δ material by the TPIE data model [41]. Chen et al [76] pointed out that the oxygen diffusion coefficient of (La_0.5Sr_0.5)₂CoO₄ with different crystal orientations: LSC50 (100) and LSC50 (001) were 1.20 × 10⁻¹³ cm² s⁻¹ and 2.97 × 10⁻¹⁴ cm² s⁻¹ respectively, and the oxygen diffusion coefficient along the a–b plane in LSC50 was at least 4 times faster than that along the C axis.

The ionic/electron conduction in RP perovskite materials occurs due to the redox reaction or the change in the valence state of the transition metal at the B site caused by the difference in the valence state of the host and dopant [91]. Figure 10 summarizes the electrical conductivity of some RP structure materials when n = 1 and n>1. On the whole, the conductivity for RP materials when n = 1 varies widely, ranging from a few to hundreds of S cm⁻¹. Comparing the electrical conductivity of LNO series, the electrical conductivity increases as the number of n increases (La₄Ni₃O_10−δ > La₃Ni₂O_7−δ > La₂NiO_4+δ ) [78]. This is related to the higher oxygen migration barrier of the RP with n = 1 than that of the n = 2 material, as reported by Zhang et al [53]. When measuring conductivity, the conductivity increases with increasing temperature at low temperatures, but the conductivity decreases at higher temperatures [25, 73, 92]. The former is due to thermally activated semiconductor properties, while the latter is related to oxygen release from oxides [93]. The loss of lattice oxygen at high temperatures leads to decrease in charge carrier concentration or the annihilation of electron holes, as shown in equation (3) [94]:

$$\begin{equation}{\text{O}}_{\text{O}}^{\text{x}}{{ + 2}}{{\text{h}}^.}{\text{ = V}}_{\text{O}}^{..}{{ + }}\frac{{\text{1}}}{{\text{2}}}{{\text{O}}_{\text{2}}}{\text{(g)}}{{.}}\end{equation} \tag{ (3) }$$

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Manthiram found similar trends in temperature dependence of conductivity in Sr_2.7−xCa_xLn_0.3Fe_2−yCo_yO_7−δ and Sr_3.2−xCa_xLn_0.8Fe_1.5Co_1.5O_10−δ [92, 100]. Shown in figure 11 is the relationship between the conductivity and the temperature [73]. The decrease in conductivity at high temperature is due to the formation of oxide ion vacancies, causing the reduction of Fe⁴⁺ to Fe³⁺ or Co⁴⁺ to Co³⁺, resulting in a diminution of the carrier concentration [101]. However, the electronic conductivity of undoped RP perovskite oxides at intermediate temperature is not sufficient for its application as cathodes for SOFC. Therefore, people are using doping strategy to increase its performance. A-site doping is generally used to replace the rare earth elements with alkaline earth metals (Sr, Ba, Ca), aiming at improving the electronic conductivity of the material. The results showed that there is an optimal ratio of A-site doping that can enhance the electrical conductivity. For example, Pikalova et al [36] found that the conductivity of Ca doped Nd₂NiO_4+δ materials increased firstly and then decreased with the Ca²⁺ doping amount.

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Replacing small amount of Nd³⁺ with Ca²⁺ leads to the increase of electron hole concentration and the p-type electron conductivity due to charge compensation (see equation (4)) [38, 39]. However, the excessive Ca²⁺doping would lead to decreasing the conductivity. The conductivity deterioration can be associated with the possible formation of electrically neutral defect associates [102, 103] or the formation of NiO impurity phase. The same phenomenon is also observed as to Pr_2−xCa_xNiO_4+δ [39, 41].

$$\begin{equation}2[{\text{O}}_{\text{i}}^{{\prime} ^{\prime} }\left] {\, + \,} \right[{\text{Ca}}_{{\text{Nd}}}^{\prime} \left] {\, \leftrightarrow \,} \right[{\text{Ni}}_{{\text{Ni}}}^{\prime} \left] {\, + \,2} \right[{\text{V}}_{\text{O}}^{..}].\end{equation} \tag{ (4) }$$

Yang et al [104] found that the electrical conductivity of Pr_2−xSr_xNiO₄ at 800 °C increased from 70.64 to 121.51 S cm⁻¹ with increasing Sr doping. Manthiram et al [100] studied the effect of the A-site ion radius of Sr_3.2−xCa_xLn_0.8Fe_1.5Co_1.5O_10−δ on the conductivity and found the electrical conductivity decreased with decreasing A-site doping radius mainly because of the decrease of the overlapping of O_2p orbital and the Co/Fe_3d orbital [105]. Hu et al [44] reported that the conductivity of La_1.9−xBi_xSr_0.1CuO_4−δ decreased with the bismuth doping as charge carriers would be consumed to form extra additional oxygen vacancies in equation (5), as evidenced by the increased non-stoichiometric ratio of oxygen δ via iodometric titration and thermalgravimetric analysis.

$$\begin{equation}{\text{V}}_{\text{O}}^{..} + \frac{1}{2}{O_2}(g) = {\text{O}}_{\text{O}}^{\text{x}} + 2{h^.}.\end{equation} \tag{ (5) }$$

In addition to the alkaline earth metal doping, the B-site transition element, such as Cu and Co doping also has a certain effect on the electrical conductivity. Zhang et al confirmed Cu-doped Nd_1.5Pr_0.5NiO_4+δ material with significantly improved conductivity of 140 S cm⁻¹ [95] and it was explained that the replacement of the Ni²⁺ (0.69 Å) with larger Cu²⁺ (0.73 Å) facilitated the hole jumping to the next adjacent B cation because of the decreased distance of the adjacent BO₆ units [106]. Zhao et al [107] also found electrical conductivity improvement of Cu doping in La_1.5Pr_0.5NiO₄. However, the conductivity of Cu-doped La₂NiO_4+δ decreased, which was considered by Tarutin et al to be ascribed to the decreased charge carrier concentration in Cu-doped compounds [35]. The influence of doping on the conductivity is complex and diverse, it is determined not only by the doping ions, amount of doping and the doping location, it is also affected by the preparation condition and the temperature dependence.

The thermal expansion coefficient (TEC) becomes an indispensable parameter when considering a promising cathode for SOFC as the lifetime of cell will be seriously attenuated if a mismatch of the thermal expansion behavior within the cell occurs [108]. The coefficient of thermal expansion refers to the change of the thermal behavior of the material, in other words, different physical behaviors on temperature [91]. The thermal expansion of perovskite and perovskite-related materials is attributed to physical expansion and chemical expansion [89]. The physical expansion is the crystal expansion caused by the harmonic atomic vibration that comes from the electrostatic attraction in the lattice [109]. The chemical expansion is caused by the partial reduction of B-site cations or the spin transition of B-site cobalt ions [100]. Figure 12 displays the TEC of some layered perovskite materials, mainly RP materials. Broadly speaking, the RP with multi AO rock layers exhibits larger TEC compared to that with one AO rock layer, and the TEC values of La_1.2Sr_0.8NiO₄ and Pr_1.2Sr_0.8NiO₄ are the lowest, which are 8.7 × 10⁻⁶ and 8.0 × 10⁻⁶K⁻¹, respectively [96]. Sr_3.2La_0.8Fe_1.5Co_1.5O_10−δ exhibited different TEC of 9.24 × 10⁻⁶ K⁻¹ and 17.11 × 10⁻⁶ K⁻¹ in 30 °C–300 °C and 300 °C–900 °C, respectively, because of the formation of oxygen vacancies due to the reduction of cobalt and iron to larger ions [100]. This is also consistent with the findings at Sr_2.7−xCa_xLn_0.3Fe_2−yCo_yO_7−δ materials reported by Manthiram et al [92]. The Co-containing RP perovskites usually have a higher TEC, mostly because of the transition of Co³⁺ cation from low spin (LS, 0.545 Å, t⁶ _2ge⁰ _g) to high spin (HS, 0.61 Å, t⁴ _2ge² _g) [110]. The larger high spin Co³⁺ brings about an increase in TEC corresponding to lattice expansion [111]. The TEC value of the cathode must be between 10 ∼ 13 × 10⁻⁶K⁻¹ in order to be compatible with other SOFC components [112]. Researchers generally use dopants to decrease the TEC of cathode materials. For example, Manthiram et al confirmed that Ca substituted LaSr₃Fe₃O_10−δ and Sr_3.2Ln_0.8Fe_1.5Co_1.5O_10−δ can shorten the bond length between cations and oxygen anions, resulting in cell volume shrinkage and decrease in TEC [92, 100]. It was found by Hu et al that Bi doped LaSr_0.1CuO_4−δ (LSC) has reduced the TEC from 13.3 × 10⁻⁶ K⁻¹ to 9.2 × 10⁻⁶ K⁻¹ [44]. It is significantly lower than the TEC of typical electrolytes, such as 12.2 × 10⁻⁶ K⁻¹ for Ce_0.8Sm_0.2O_2-δ (SDC) and 10.5 × 10⁻⁶ K⁻¹ for yttria stabilized zirconia (YSZ) [112, 113].

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Besides A-site doping, doping on B-site can also influence the TEC of the material. Zhang et al [95] found that TEC for Cu-doped Nd_1.5Pr_0.5NiO_4+δ has reduced to 13.6 × 10⁻⁶ K⁻¹. Tarutin et al [35] also concluded the decrease in TEC gradually with increasing Cu doping in La₂NiO₄. Zhao et al [107] confirmed that introducing a small amount of Cu at the Ni site can effectively reduce the value of TEC. The doping of B-site transition metals in RP perovskites does not always bring a positive effect. Replacing Co with Mn and Fe in RP perovskites could lower the TEC value, in the meantime, it will also cause the decreased electronic and ionic conductivity. Therefore, in the development of SOFC cathode materials, it is necessary to find a suitable compromise between the high conductivity and the matched thermal performance with other cell components [108].

The operating temperature for oxygen-conducting type IT-SOFC is in the range of 650~800 °C. Composition optimization of the single phase and composite cathodes fabrication are the most common schemes to applying RP perovskites as cathode for SOFC in different scenarios. A-site and B-site doping in perovskites has considered to be effective method to modify the ORR catalytic activity [119]. Recently, most scholars have adjusted the relationship between structure and performance by changing the composition of RP oxides to improve cathode performance. Generally, substitution at the A position in RP perovskites refers to partial replacement of rare earth elements with alkaline earth metals (Sr, Ba, Ca) [39, 120–122]. For example, Yang et al [104] prepared Pr_2−xSr_xNiO₄ by a glycine-nitrate process, it was found that the area specific resistance (ASR) of the Pr_2−xSr_xNiO₄ is the minimum when x = 0.4. In addition, Sharma et al [123] found that the R_p of Pr-doped La_2−xPr_xNiO_4+δ cathode became smaller with the increase of x at 600 °C, while the thermal stability decreased. This requires compromise between stability and electrochemical performance. When x = 1, LaPrNiO_4+δ phase had excellent long-term stability in a symmetrical state (operating at 700 °C for 500 h). In addition, a number of studies have shown that reasonable substitution of transition elements such as Co, Fe and Cu can effectively improve the performance of the cathode [83, 107, 124–126]. The ORR catalytic activity of cobalt-based RP is the highest, but there are some problems such as high coefficient of thermal expansion and poor stability. Therefore, Fe was used instead of cobalt to obtain suitable TEC and stability. Boulahya et al [25] reduced the TEC of the Eu₂SrCo_1.5Fe_0.5O₇ cathode by partially replacing Co with Fe. The area specific resistance at 700 °C in air was 0.26 Ω cm². The composites of Eu₂SrCo_1.5Fe_0.5O₇ and Gd_0.1Ce_0.9O_2-δ (CGO) were heated at 900 °C (70:30 wt%) for 12 h without any reaction or decomposition, indicating that these cathode compounds had good long-term stability. The La₂Ni_1−xCu_xO_4+δ material showed highest conductivity at Cu substitution amount of 0.4, and the lowest ASR value could be obtained by using La_0.9Sr_0.1Ga_0.8Mg_0.2O_3-δ (LSGM) electrolyte for its matched TEC with LSGM [120, 127]. Some researchers also improved the electrocatalytic performance and stability of fuel cells by co-doping [40, 98, 128, 129]. Navarrete et al [129] studied co-doping of LNO with Pr and Co and found that La_1.5Pr_0.5Ni_0.8Co_0.2O_4+δ cathode exhibited the lowest electrode polarization resistance and activation energy values in the temperature range of 450~900 °C and a maximum power density of ∼400 mW cm⁻² at 700 °C using pure hydrogen and air.

Although doping has been conducted to optimizing the ORR catalytic activity, however, it is still hard to compromise with sufficient electrical conductivity and suitable TEC. By designing composite cathode via mechanical mixing and infiltration provides novel insight in further improving the electrochemical performances and stability. RP perovskites have good oxygen surface exchange property but low electrical conductivity and large TEC, by mixing with the electrolyte, the TEC and the oxygen ionic conductivity will be modified, thus effectively improving the thermal compatibility between cathode and electrolyte and enhancing the stability and durability of cell performance. The La₂NiO₄–SDC composite electrode was successfully fabricated by an infiltration method that was studied to be a stable cathode as no obvious degradation was observed after running at 650 °C for 210 h [130], the maximum power density at 800 °C with hydrogen as fuel gas was 550 mW cm⁻², which was higher than similar cell by using Pt-YSZ cathode [131]. Sharma et al [132] detected a great decrease of R_p by adding 50 wt% CGO electrolyte into La₃Ni₂O_7−δ (L3N2) from 1.00 Ω cm² to 0.03 Ω cm² at 700 °C, and no chemical reactivity between L3N2 and CGO was observed when being heated in air at 900 °C for 2 weeks. Shirani et al [133] systematically studied the performance of (Nd_0.9La_0.1)_1.6Sr_0.4Ni_0.75Cu_0.25O_3−δ (NLSNC4) and Ce_0.8Sm_0.2O_2-δ(SDC) composite cathode materials and found that the R_p of the composite cathode decreased firstly and then increased with the SDC content at 800 °C, and NLSNC4 with 50 wt% SDC showed the minimum ASR of 0.12 Ω cm², however, its conductivity decreased with the increase of SDC content. Chiu et al [115] studied the effect of a series of weight percentages of SDC on the electrochemical performance of Pr₂CuO₄ (PCO)-SDC composite cathode, Pr₂CuO₄ (PCO)-10 wt% SDC composite cathode showed the lowest area specific resistance of 0.35 Ω cm² and the maximum power density of 347.50 mW cm⁻². Chen [134] reported a new type of Nd₂NiO_4−δ (NNO) composite cathode impregnated with scandia stabilized zirconia (SSZ) showed excellent power density of 1.26 W cm⁻² and the ASR of the cathode was 0.04 Ω cm² at 800 °C. RP perovskites cathode has high oxygen surface exchange coefficient and stability, so it is one of the choices as protective layer for SOFC cathode. Some people have studied the composite of RP perovskites and perovskites by solution infiltration method. Li et al [135] observed 40 % polarization resistance drop of infiltrated La₂NiO₄@La_0.6Sr_0.4Co_0.2Fe_0.8O₃ (LSCF) to 0.31 Ω cm² from 0.54 Ω cm² at 750 °C. The La₂NiO₄@PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ (LNO@PBSCF) cathode with LNO protective layer also exhibited good tolerance to Cr poisoning, the LNO coating effectively prevented the reaction between the volatile Cr species and the PBSCF cathode when treated at 750 °C for 1200 min [136].

RP perovskites exhibit encouraging performance as cathode in O-SOFC. Recently, researchers are also working on the application of RP perovskites as cathode in proton-conducting SOFC, which is a potential alternative to traditional oxide ion-conducting SOFC [89]. Compared with traditional SOFC, H-SOFC has lower working temperature (450~700 °C) and usually composed of a proton conductive electrolyte, through which the protons transport from the anode to the cathode and produce water. Therefore, the cathode for H-SOFC is even more demanding. Proton migration is accomplished through Grotthuss mechanism in SOFC [66], which mainly occurs in the a–b plane. The electrolyte for H-SOFC is different from O-SOFC. Iwahara [137, 138] firstly reported promising proton conductors such as SrCeO₃ and BaCeO₃. Currently, the commonly used proton conductors are barium ceratebased materials, such as BaZr_1−xY_xO_3−δ (BZY) and BaCe_1−xZr_xY_yO_3−δ (BCZY) [89]. In H-SOFC, it will be beneficial if the cathode has both proton and oxide ion conductivity, because it can expand the reaction area in the cathode. In order to realize this idea, recent research work has focused on applying traditional O-SOFC cathodes into H-SOFC to test its electrochemical performance and stability under humidified atmosphere. For example, Yang et al [96] reported that Ln_1.2Sr_0.8NiO₄ (BZCY) (Ln = La, Pr) can maintain a stable structure within 100 h in humid air (20% vol H₂O) at 800 °C. Peng et al [125] studied the R_p of Co-doped La₃Ni_1.9Co_0.1O₇ cathode changed from 1.12 Ω cm² to 0.35 Ω cm² at 650 °C, and its maximum power density at 700 °C was 0.398 Wcm⁻². Zhang [53] reported that Co or Cu doped La₂NiO₄/La₃Ni₂O₇ accelerated the proton diffusion of the material due to its low migration energy along the path. Wang et al [116] proved that Mn-doped Pr₂BaNi₂O_7−δ at the B site has enhanced oxygen diffusion and proton diffusion property. The single cell composed of BZCYYb-based electrolyte had a peak power density of 1070 mW cm⁻² at 700 °C and had no degradation at 0.7 V at 600 °C for 100 h. Miao et al [97] successfully synthesized Sr and Fe co-doped La₂NiO₄ cathode, the cell based on La_1.2Sr_0.8Ni_0.6Fe_0.4O_4+δ cathode exhibited high maximum power density (MPD) of 781 mW cm⁻² and low polarization resistance 0.078 Ω cm² at 700 °C, which also showed excellent stability after operating at 600 °C for 100 h. Therefore, co-doping Sr and Fe could be considered as promising strategy in improve electrochemical performance without increasing TEC.

Recently, the composite cathode composed of RP perovskites with proton-conducting electrolyte has shown encouraging performance in H-SOFC. For example, Wang et al [83] found that the MPD of Sr₃Fe₂O_7−δ -5 wt% BaZr_0.3Ce_0.5Y_0.2O_3−δ composite cathode at 600 °C is 530 mW cm⁻², which is better than PrBaCo₂O_5+δ -BZCYZ (266 mW cm⁻²) [139] and Sm_0.5Sr_0.5CoO_3-δ-BaZr_0.3Ce_0.5Y_0.16Zn_0.04O_3−δ (BZCYZ) (364 mW cm⁻²) [140]. More importantly, no discharge decay was observed during the 100 h test. Li et al [141] confirmed that 43.6 wt.% (Pr_0.9La_0.1)₂(Ni_0.74Cu_0.21Nb_0.05)O_4−δ (PLNCN) impregnated BaZr_0.1Ce_0.7Y_0.2O_3−δ (BZCY) electrode achieved the maximum power density of 0.77 W cm⁻² at 700 °C, which is significantly higher than other cathodes such as LSCF-BZCY [142], and it did not show any degradation at 600 °C for 200 h. Fu et al [143] added BZCY to the (LaSr)_0.9Fe_0.9Cu_0.1O₄ (LSFCu) cathode to increase the ion conductivity and minimize the thermal mismatch between the electrode and the electrolyte. The area specific resistance at 800 °C is 0.19 Ω cm², and the peak power density is 573 mW cm⁻² at 800 °C in humidified H₂.