Reversible solid-oxide cells (RSOCs) are high-efficiency, fuel-flexible electrochemical devices that interconvert electricity and chemical fuels^1,2. Operating in fuel-cell (FC) mode, RSOCs electrochemically oxidize H₂ or carbon-containing fuels to generate electricity³. In electrolysis-cell (EC) mode, they store surplus electrical energy as H₂ or synthetic fuels⁴. This operational flexibility positions RSOCs as key components for balancing temporal mismatches between renewable energy generation and demand^5,6. In both modes, the oxygen electrode governs a large fraction of the losses and the lifetime⁷. Modern deployment scenarios demand operation at high current density (> 0.5 A·cm⁻²) for tens of thousands of hours, under gas streams that often contain multiple reactive impurities⁸. These demanding conditions expose the oxygen electrode to volatile contaminants, particularly chromium (Cr) species from metallic interconnects, which accelerate structural degradation and interfacial deactivation⁹. Consequently, stability metrics under combined thermal, chemical, and electrochemical stresses remain insufficient for grid-scale RSOC systems.

Cr-induced degradation in perovskite oxygen electrodes proceeds through multiple coupled pathways. In benchmark compositions such as La_0.6Sr_0.4Co_0.2Fe_0.8O₃₋_δ (LSCF), volatile CrO₃ species react with surface-segregated SrO to form insulating SrCrO₄ precipitates¹⁰. This process causes pore blockage, reduces active surface area, and contracts the triple-phase boundary region where gas, electrode, and electrolyte phases meet. The resulting suppression of oxygen reduction and evolution kinetics drives rapid performance decay, with degradation rates often exceeding 1% per 100 hours under practical operating conditions. Current mitigation strategies include protective coatings on metallic interconnects to suppress Cr volatilization^11,12, implementation of Cr getter layers to sequester volatile species¹³, and development of intrinsically Cr-resistant electrode materials³. While each approach offers partial solutions, none adequately addresses the fundamental coupling between surface chemistry, bulk stability, and catalytic activity that governs long-term electrode performance.

Recent advances in oxygen-electrode materials have explored three primary directions. Surface modification strategies employ protective layers such as BaO¹⁴, BaCoO₃¹⁵, or Gd_0.1Pr_0.1Ce_0.8O₂₋_δ¹⁶ to physically or chemically inhibit Cr deposition. These approaches, however, introduce additional processing complexity and create new interfaces susceptible to degradation under thermal cycling. High-entropy oxide compositions, exemplified by ((La_0.25Pr_0.25Nd_0.25Sm_0.25)Ba_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ⁷, and SrCo_0.5Fe_0.2Ti_0.1Ta_0.1Nb_0.1O₃₋_δ¹⁷, enhance configurational entropy to stabilize the lattice structure but often dilute catalytically active sites, compromising intrinsic activity. Exsolution engineering produces reactive nano-domains such as PrNi_0.5Mn_0.5O₃–PrOₓ¹⁸, or Ba₁₋_xCe_0.8Gd_0.2O₃₋_δ–BaCO₃¹⁹ for in situ Cr trapping, yet typically lacks sufficient bulk stabilization and surface chemistry control. These single-mechanism strategies fail to address the interconnected nature of degradation pathways, limiting their effectiveness under RSOC-relevant stresses.

Here, we introduce a triple-barrier design that integrates structural, interfacial, and chemical defenses within a single, scalable perovskite platform, as shown in Fig. 1a. The rationale is to simultaneously address lattice instability, volatile Cr ingress, and acid–alkali imbalance, which are intrinsically coupled in conventional LSCF oxygen electrodes. Starting from the LSCF base composition, partial Sr to Ba substitution produces La_0.6Sr_0.1Ba_0.3Co_0.2Fe_0.8O₃₋_δ (LSB_0.3CF), which suppresses oxide segregation and stabilizes the lattice—establishing the first internal barrier. Further increasing the A-site Ba excess generates La_0.6Sr_0.1Ba_0.35Co_0.2Fe_0.8O₃₋_δ (LSB_0.35CF), facilitating in situ exsolution of BaCoO₃ nano-domains. These exsolved phases preferentially capture Cr species and simultaneously limit contaminant ingress while promoting oxygen-surface kinetics—creating an external barrier. The final optimization involves trace Mo substitution at the B-site, yielding La_0.6Sr_0.1Ba_0.35Co_0.2Fe_0.78Mo_0.02O₃₋_δ (LSB_0.35CFM_0.02). This modification modulates surface basicity to inhibit SrO/BaO segregation and enhance resistance to acidic contaminants such as CrO₃, providing a chemical barrier that complements the structural and exsolved barriers. Together, these three complementary mechanisms form a synergistic triple protection: lattice stabilization resists bulk degradation, exsolved-phase trapping intercepts Cr species at the surface, and acidity regulation mitigates chemical driving forces for poisoning. As demonstrated below, this integrated strategy delivers state-of-the-art oxygen-exchange kinetics while maintaining robust durability under severe Cr exposure, thereby narrowing the activity–durability gap for practical RSOC oxygen electrodes.

To elucidate the oxygen reduction reaction (ORR), the electrochemical impedance spectroscopy (EIS) of LSB_0.35CFM_0.02 (Fig. 3a) at various pO₂ is analyzed using the distribution of relaxation time (DRT). The resulting curves (Fig. S2) reveal three distinct processes: low-frequency (LF), mid-frequency (MF), and high-frequency (HF) contributions. The dependence of polarization resistance (R_p) on pO₂ follows R_p = k(pO₂)ⁿ³⁰. (Fig. 3b) The LF feature is associated with O₂ adsorption (n ≈ 1)^31,32, the MF process arises from the oxygen adsorption/dissociation at the surface (n ≈ 1/2)^33,34, and the HF feature corresponds to O²⁻ charge transfer across the electrode-electrolyte interface (n ≈ 0)^35,36.

Figure S3 presents the EIS of half-cells at 650–800°C after deleting the ohmic resistance (R_o). In this study, the R_p of blank LSCF exceeds previously reported values (0.10–0.35 Ω·cm² at 700°C)^36,37, likely due to screen-printing or sintering procedures that influence electrode porosity. The R_p of LSB_xCFM_y (x = 0, 0.3, 0.35; y = 0, 0.02) decreases with Ba/Mo doping, with LSB_0.35CFM_0.02 manifesting the lowest R_p (0.058 Ω·cm²), 70.4% reduction versus LSCF (0.196 Ω·cm²) at 750°C. Similarly, the activation energy (E_a) declines with increasing Ba/Mo, with LSB_0.35CFM_0.02 showing the lowest E_a (1.567 eV), indicating faster ORR kinetics (Fig. 3c). The improvement is ascribed to the synergistic effect of Ba/Mo co-doping. For LSB_0.3CF, the rational incorporation of Ba improved ORR kinetics by optimizing O₂ diffusion pathways, as evidenced by the significant reduction of LF peak in DRT curves (Fig. S4)³⁸. In LSB_0.35CF, A-site Ba excess induces a secondary BaCoO₃ phase with excellent oxygen adsorption/dissociation ability^15,39, and the resulting particles supply additional active sites⁴⁰, thereby enhancing catalytic activity relative to LSB_0.3CF. Compared with LSB_0.35CF, LSB_0.35CFM_0.02 demonstrates enhanced catalytic activity due to the Mo-for-Fe substitution that modulates the B-site electronic structure, potentially improving electron transport pathways within the Co-O-Fe/Mo framework and accelerating the surface exchange kinetics^41,42.

Figure 3d gives the R_p variation of half-cells with LSB_xCFM_y (x = 0, 0.3, 0.35; y = 0, 0.02) at 750°C, OCV under Cr₂O₃ exposure. LSCF suffers a 91.3% degradation, with R_p escalating from 0.196 to 0.375 Ω·cm² within 120 h (Fig. 3f). The deterioration originates mainly from SrCrO₄ formation through the reaction between CrO₃ and segregated SrO. In contrast, LSB_0.3CF demonstrates a 63.6% degradation, with R_p increasing from 0.088 to 0.144 Ω·cm² (Fig. S7). The improvement arises from the larger Ba²⁺ radius (1.61 Å vs. 1.44 Å for Sr²⁺)^21,22, which increases the A-site ionic radius and elevates the tolerance factor (t, Eq. S1)⁴³. A higher t value drives the perovskite lattice toward an ideal cubic symmetry, thereby alleviating local strain, stabilizing A–O coordination, and suppressing the oxide segregation^38,44. The Ab Initio Molecular Dynamics simulations further corroborate that LSB_0.3CF maintains structural integrity with minimal atomic displacement after 2 ps (Fig. S5), whereas the LSCF model exhibits noticeable Co segregation (Fig. S6). These results establish Ba-induced lattice stabilization as the first structural barrier against Cr-induced degradation. As shown in Fig. 3d, the degradation of LSB_0.35CF and LSB_0.35CFM_0.02 is greatly suppressed, mainly because excessive Ba promotes the exsolution of BaCoO₃, which preferentially captures Cr species and blocks their diffusion within the electrode. As a result, LSB_0.35CF shows a 20.8% rise in R_p from 0.072 to 0.087 Ω·cm² (Fig. S8), whereas LSB_0.35CFM_0.02 displays almost no change (0.058 − 0.060 Ω·cm², Fig. 3g). This impressive stability is further attributed to Mo introduction, which strengthens the lattice framework through strong Mo–O bonding and electronic-structure tuning while simultaneously providing acid–alkali interactions (Fig. 1b)⁴¹. Specifically, MoO₃ (5.2) provides an acidity level that matches well with CrO₃ (6.6) and strongly interacts with alkaline surface species such as SrO (− 9.4) and BaO (− 10.8)²⁰. Such dual affinity suppresses their outward diffusion and establishes an additional chemical barrier against Cr poisoning. The acid–alkali modulation by Mo was verified by CO₂/NH₃-TPD tests. In CO₂-TPD (Fig. 1c), LSB_0.35CFM_0.02 exhibits a pronounced decrease in CO₂ desorption, particularly in the medium-temperature region (300–500°C), indicating a significant reduction of strong alkali sites compared with LSCF⁴⁵. Conversely, NH₃-TPD (Fig. 1d) shows that LSB_0.35CFM_0.02 displays enhanced NH₃ desorption in the low-temperature region (50–200°C)⁴⁶. This trend suggests that Mo incorporation primarily increases weak acid sites while suppressing strong alkali sites, thus shifting the acid–alkali balance toward enhanced acidity and reduced basicity, consistent with the proposed chemical barrier mechanism.

Figure 3e compares the R_p variation of half-cells with LSCF/LSB_0.35CFM_0.02 at 0.5 A·cm⁻² current density under Cr₂O₃ exposure for 1000 h. Under high current load, LSCF degrades rapidly (1.45% h⁻¹), whereas LSB_0.35CFM_0.02 exhibits negligible R_p increase with the rate of 0.25% h⁻¹, confirming superior Cr tolerance under polarization. To clarify the electrochemical-driven degradation mechanism of Cr poisoning, DRT analysis is further performed (Fig. 3h,k). Notably, under OCV conditions, the LF peak increases more significantly than the MF/HF peaks, whereas under 0.5 A·cm⁻² load, the HF peak shows a more pronounced increase. This suggests that, in addition to oxygen adsorption and transport, the interfacial reactions between the electrode and electrolyte become increasingly vulnerable under current load^35,36. The pronounced growth of the HF peak indicates aggravated interfacial charge-transfer resistance. Under polarization, enhanced oxygen vacancy generation and cation migration promote SrO segregation, which provides reactive sites for Cr species⁴⁷. The accumulation and reaction of these species at the interface disrupt electronic/ionic transport pathways, thereby accelerating interfacial deactivation. As a result, polarization not only accelerates surface poisoning but also amplifies Cr-induced interfacial blockage, leading to pronounced deterioration of charge transfer and overall electrode performance under operating conditions⁴⁸.

Both fresh LSCF and LSB_0.35CFM_0.02 electrodes exhibit porous, interconnected skeletons in Fig. 4a,b. After Cr₂O₃ exposure, LSCF develops coarse irregular deposits that severely compromise its structure (Fig. 4c). These deposits are identified as SrCrO₄,²⁵ originating from the reaction of segregated SrO with CrO₃. In contrast, the LSB_0.35CFM_0.02@Cr surface remains comparatively clean and structurally intact (Fig. 4d and S9). The composition of the Cr deposits on LSCF/LSB_0.35CFM_0.02 is further analyzed through Raman and XRD (Fig. S10).

Figure 4e gives the Raman spectra of LSCF/LSB_0.35CFM_0.02 after Cr₂O₃ exposure at 750°C for 120 h. The LSCF@Cr sample exhibits three distinct sets of Raman peaks, with characteristic peaks at 340–430, 600–700, and 850–990 cm⁻¹. The peaks at 340–430 cm⁻¹ are identified as Cr₂O₃, while 850–990 cm⁻¹ peaks correspond to SrCrO₄⁴⁹. The Co₃O₄ peaks observed at 600–700 cm⁻¹ in LSCF@Cr signify Co segregation^50,51. In contrast, no apparent Co₃O₄ segregation is observed in LSB_0.35CFM_0.02. For LSB_0.35CFM_0.02@Cr, the intensity of the SrCrO₄/BaCrO₄ peaks is obviously reduced, and the Cr₂O₃ peaks are also markedly less intense compared to LSCF@Cr. The Raman mappings of LSCF@Cr and LSB_0.35CFM_0.02@Cr in 20×20 µm region are viewed in Fig. 4f,g. These mappings are acquired by the Raman peak of SrCrO₄/BaCrO₄ at 859 cm⁻¹. A substantial quantity of SrCrO₄ is observed on the LSCF@Cr, while only a small amount of Cr contamination is seen on LSB_0.35CFM_0.02@Cr.

X-ray photoelectron spectroscopy (XPS) is performed on LSCF/LSB_0.35CFM_0.02 before and after Cr₂O₃ exposure to detect the surface changes. Figure 4h,i illustrates that the Sr 3d region is fitted into two sets of double peaks, including the surface Sr part (Sr_sur: 134.2 and 135.5 eV) and the lattice Sr part (Sr_lat: 133.4 and 134.8 eV)⁵². After Cr deposition, the intensity of Sr_sur rises due to SrO segregation and SrCrO₄ formation, consistent with the O_sur peak appearing at 529.8 eV⁵³ in Fig. 4j,k. The appearance of SrCrO₄ is further confirmed by the Cr⁶⁺ double peaks at 578.7 and 588.3 eV in Fig. 4l,m, and the double peaks at 576.8 and 586.1 eV belong to Cr₂O₃⁵⁴. After Cr deposition, the percentage of Sr_sur in LSB_0.35CFM_0.02@Cr is smaller than that of Sr_lat, and the percentage of O_sur is reduced compared to LSCF@Cr. The percentage of Cr⁶⁺ in LSB_0.35CFM_0.02@Cr is also smaller than that in LSCF@Cr, confirming a notable reduction in contaminants like SrCrO₄/BaCrO₄ in LSB_0.35CFM_0.02.

Figure 5a depicts the single-cell architecture with an oxygen electrode thickness of ~ 18 µm. As shown in Fig. 5b and S11, the LSB_0.35CFM_0.02 cell delivers superior peak power densities of 0.598–1.352 W·cm⁻² at 700–800°C, considerably higher than those of the LSCF cell (0.248–0.673 W·cm⁻²). These findings confirm that the LSB_0.35CFM_0.02 cell demonstrates superior electrochemical performance, comparable to other reported cells (Fig. S12)^{21,49,55–57}. In EC mode, the single cell with LSB_0.35CFM_0.02 exhibits superior catalytic performance than LSCF cell (Fig. 5c and S13). Under 1.5 V voltage with 40% H₂O–60% H₂ fuel, the LSB_0.35CFM_0.02 cell achieves a current density of 2.08 A·cm⁻², 39.6% higher than the LSCF cell (1.49 A·cm⁻²) at 800°C.

The Cr poisoning tests is first conducted in FC mode at 750°C, 0.5 A·cm⁻² for 500 h (Fig. 5d–g). The R_o of the LSCF cell rises quite at 0.158% h⁻¹, while the LSB_0.35CFM_0.02 cell remains basically unchanged (Fig. 5h). The R_p of the LSCF cell escalates at 0.781% h⁻¹, while the LSB_0.35CFM_0.02 cell shows only a minor increase of 0.029% h⁻¹ (Fig. 5i). In both cases, R_p degradation surpasses R_o, suggesting the importance of an excellent oxygen electrode for SOC. As shown in Fig. 5f,g, the marked growth of the P3 and P4 peaks in DRT curves indicates that Cr poisoning severely deteriorates the surface exchange of O₂ and the bulk transport of O²⁻ in the oxygen electrode⁵⁸. This degradation, driven by the electrochemical potential gradient and oxygen vacancies, originates from SrO segregation that reacts with infiltrating CrO₃ to form insulating SrCrO₄, thereby reducing electrode porosity and hindering the ORR^59,60. Consistently, time-of-flight secondary-ion mass spectrometry (TOF-SIMS) and 3D images in Fig. 6a–c reveal severe and homogeneous CrO₄⁻ deposition in LSCF, whereas only weak signals are observed in LSB_0.35CFM_0.02, confirming that Ba/Mo co-doping effectively suppresses Cr intrusion and accumulation inside the cell. Furthermore, the LSB_0.35CFM_0.02 cell demonstrates superior durability under long-term Cr poisoning (Fig. 5j). Under a 0.5 A·cm⁻² load, both cells initially present a voltage of about 0.95 V. After Cr₂O₃ exposure, the LSCF-cell voltage decreases to 0.53 V, while the LSB_0.35CFM_0.02 cell maintains 0.95 V. Finally, the RSOC with LSB_0.35CFM_0.02 demonstrates super reversibility and stability at ±0.5 A·cm⁻² for 1000 h. As a result, LSB_0.35CFM_0.02 exhibits less Cr-induced degradation and superior output performance during practical RSOC operation (Fig. 6d).

The adsorption process of CrO₃ on the LSCF, LSBCF and LSBCFM surfaces was investigated, with the resulting configurations at various sites in Fig. 7. It can be seen that when CrO₃ is adsorbed on the three surfaces, Cr and O in CrO₃ tend to bond with O and M (La/Sr/Ba) sites on the surface, respectively. The adsorbed Gibbs free energies (G_ads) of CrO₃ at different sites on the three surfaces are summarized in Table 1.

Specifically, on the LSCF and LSBCF surfaces, the adsorption configurations of CrO₃ are optimized at all possible sites, namely (La–O–La, Sr–O–Sr, La–O–Sr, and La–O–Ba, and the corresponding adsorption energies are calculated. Energetically, CrO₃ preferentially adsorbs at the La–O–La and Sr–O–Sr sites of the LSCF and LSBCF surfaces, with the G_ads of − 3.62 and − 3.74 eV, respectively. On the LSBCFM surface, similar to LSBCF, the adsorption of CrO₃ has also been tried at the four sites (La–O–La, Sr–O–Sr, La–O–Sr, La–O–Ba). After structural optimization, the adsorption configurations at La–O–La, Sr–O–Sr, La–O–Sr converged to the same stable structure, in which the three oxygen atoms of CrO₃ bind toward one La and two Sr, with a distance of 2.72 Å and a G_ads of − 2.15 eV. Meanwhile, the G_ads of CrO₃ at the La–O–Ba site on the LSBCFM surface is − 3.58 eV.

Overall, Ba/Mo co-doping significantly weakens the adsorption strength of CrO₃ on the surface, particularly at the surface Sr site (La–O–Sr), elevating Cr tolerance of the electrode. Bader charge is analyzed to illustrate the mechanism of Ba/Mo co-doping in weakening CrO₃ adsorption. In all cases, net charge transfer occurs from the perovskite surface to the CrO₃ species during adsorption. As illustrated in Fig. S14, substantial electron accumulation is observed between Cr in CrO₃ and the surface oxygen sites, as well as between O in CrO₃ and La/Sr sites. This redistribution of electron density corresponds to chemical bonds between CrO₃ and the surface upon adsorption. In summary, Ba/Mo co-doping effectively reduces the adsorption strength of CrO₃ on the LSBCFM surface, particularly at the La–O–Sr sites. This phenomenon mainly originates from the reduced electron density at the oxygen site within the La–O–Sr sites, weakening its bonding interaction with Cr.

In this work, we establish a multiscale design principle for overcoming the fundamental trade-off between catalytic activity and chemical stability in SOC electrodes. Through systematic integration of lattice engineering, interfacial modification, and surface chemistry control, we developed La_0.6Sr_0.1Ba_0.35Co_0.2Fe_0.78Mo_0.02O₃₋_δ as a high-performance oxygen electrode with exceptional chromium tolerance. The material achieves a R_p of 0.058 Ω·cm² at 750°C, 70.4% lower than standard LSCF, while maintaining degradation rates below 0.25% h⁻¹ under 0.5 A·cm⁻² load with Cr₂O₃ exposure for 1000 hours.

Three complementary mechanisms contribute to the enhanced Cr resistance. First, Ba substitution at the A-site increases the tolerance factor toward unity, stabilizing the perovskite lattice and suppressing Sr segregation that serves as the primary precursor for chromate formation. Second, controlled Ba excess triggers in situ exsolution of BaCoO₃ nanodomains that function as sacrificial Cr scavengers while providing additional catalytic sites. Third, trace Mo incorporation modulates surface acid-base properties, creating unfavorable thermodynamics for CrO₃ adsorption. DFT calculations confirm this mechanism, revealing reduced electron density at vulnerable La-O-Sr sites that weakens Cr-O interactions.

Comprehensive characterization through impedance spectroscopy, surface analytical techniques, and computational modeling provides mechanistic insights extending beyond empirical performance metrics. DRT analysis reveals that current-induced degradation preferentially affects electrode/electrolyte interfacial processes, while TOF-SIMS analysis demonstrates minimal Cr penetration in the optimized composition. These findings identify critical degradation pathways and validate the effectiveness of the integrated protection strategy.

This approach transcends conventional single-mechanism mitigation strategies by addressing the interconnected nature of degradation processes in high-temperature electrochemical systems. The design principles demonstrated here, which combine thermodynamic stabilization, kinetic barriers, and chemical incompatibility, provide a framework applicable to diverse electrode materials facing volatile contaminant challenges. Implementation of these materials could substantially extend RSOC operational lifetimes, advancing their deployment for grid-scale energy storage and sustainable fuel production.