Strategic atomic trapping at heterointerfaces for protonic ceramic cells

Domain: physics
Published: 2025-11-24
Editorial outcome: minor_revision

Abstract
Engineering dynamic heterointerfaces with atomic precision is critical for unlocking the full potential of reversible protonic ceramic electrochemical cells in sustainable energy conversion, while minimizing precious metal use in composite electrodes. Here, we introduce an atomic trapping strategy that restructures the interfacial chemistry of Ba0.5Sr0.5Co0.8Fe0.2O3-δ perovskite and Ru@CeO2-δ fluorite heteroelectrodes, achieving catalytic synergy through manipulating Ru coordination environment. A scalable co-sintering protocol induces thermodynamically driven Ru migration from the CeO2 lattice into the perovskite matrix, creating coupled interfaces. This optimizes interfacial electron redistribution, generates interfacial oxygen vacancies, and improves triple conductivity and hydration kinetics. The electrode with low Ru loading exhibits bifunctionality, delivering a peak power density of 1.51 W cm−2 and an electrolysis current density of −2.21 A cm−2 at 650 °C. It demonstrates notable durability with minimal degradation (0.09 mV h−1) over 400 h at 600 °C, providing a universal strategy for next-generation solid-state energy devices.

Introduction
The pursuit of robust electrochemical energy infrastructures has propelled reversible protonic ceramic electrochemical cells (R-PCECs) into the spotlight, due to their specific ability to switch reversibly between fuel cell and electrolysis modes with high energy efficiency at intermediate temperatures (400–600 °C)1–4. These devices hold great promise for bridging renewable electricity with storable chemical fuels—producing green hydrogen from intermittent solar and wind energy, while efficiently reconverting it to electrical power during peak grid demand5,6. The wide-spread applications of this technology require the development of efficient electrode with high activity, durability and cost effectiveness. Noble metal oxides like RuO2 are highly active for catalyzing oxygen reduction reaction (ORR) and oxygen evolution reaction (OER), while their high price inhibits the use in large amount7,8. In addition, a fundamental contradiction remains that inhibits the development of robust R-PCECs, i.e., the very interfacial architectures that enable multifunctionality also inherently drive operational degradation3,9. State-of-the-art air electrodes, responsible for facilitating both ORR and OER during bidirectional operation, suffers significant performance losses due to interfacial detachment—a mechanically driven failure mode that accelerates catalytic deactivation and disrupts ionic conduction networks under prolonged operation10,11. Therefore, development of efficient electrode materials with minimized use of precious metal elements and re-enforced interface is crucial in the development of R-PCEC.

Precious metal-free perovskite oxides, such as Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF), PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (PBSCF), BaCo0.4Fe0.4Zr0.1Y0.1O3-δ (BCFZY), and Ba0.875Fe0.875Zr0.125O3-δ (BFZ), have dominated R-PCEC electrode designs, leveraging their high ionic conductivity and flexible oxygen non-stoichiometry12–14. Despite considerable progress, the performance and stability of R-PCECs remain insufficient for practical applications. This limitation primarily stems from weakly bonded heterogeneous structures at the interfaces between cell components, which hinder ionic and electronic conduction, leading to increased interfacial resistance at intermediate to low temperatures9,15. Nanocomposites have emerged as a useful air electrode for R-PCEC with improved proton conductivity and reduced thermal expansion, while they introduce internal interface issues inside the electrodes16. Interface engineering is then proposed, which presents a promising approach to enhance cell performance in terms of both activity and durability by optimizing the contact area, coupling interactions, charge distribution, and lattice structure at these critical interfaces16,17.

Current interface engineering, ranging from chemical etching to nanocomposite interlayers, however, merely treats symptoms (e.g., contact resistance) rather than curing the root disease3,9. For example, roughening electrolyte surfaces through acid treatment or incorporating phase-separated composite interlayers has been shown to improve electrode/electrolyte interactions, reducing interfacial resistance and improving electrochemical performance3. While recent advances in interface tailoring offer glimpses of hope, but demand energy-intensive processes incompatible with industrial scale-up18–20. More critically, the far greater design complexity arises from the internal interfaces within composite electrode matrices, where the synergistic interactions between the multiple phases at the interface dictate overall electrochemical activity and stability21. This emphasizes the deterministic role of internal heterointerfaces. Recent advances have primarily addressed simpler electrode/electrolyte interfaces, leaving the vastly more complex internal electrode architecture largely unoptimized, which is a critical gap that fundamentally limits R-PCEC development.

Here, we report an active atom-scale interface engineering strategy through manipulating the Ru migration that greatly increases the activity and durability for both ORR and OER in R-PCECs while realizing minimal use of the precious metal element. We first doped a minor amount of Ru into the lattice of CeO2 (RC), and then composited it with BSCF perovskite to form a composite electrode. A subsequent thermal processing then deliberately triggered controlled Ru migration, and selective capture of migrating Ru atoms from the RC phase by BSCF was happened (Fig. 1a, b), which facilitated the formation of covalent metal-oxygen bridges at the rich interfaces between the two phases, thereby achieving interfacial charge redistribution, generating rich interfacial oxygen vacancies, and imparting favorable triple conductivity and hydration elasticity (Fig. 1c, d). Leveraging these synergistic effects, our R-PCECs demonstrated absolute advantages in bifunctionality, achieving an area-specific resistance of 0.08 Ω cm2 at 600 °C, with negligible decay in reversible variable current density cycling (96 cycles) and thermal cycling tests. More importantly, we interpreted these macroscopic triumphs through experiments and theoretical analysis, establishing a universal interface regulation rule, providing a scalable and viable solution for advancing R-PCEC technology.

Results
Evaluation of heterogeneous structures
Composite electrodes often show competitive activity for ORR in protonic ceramic cells due to the potential of synergy between the different phases. Sm0.2Ce0.8O2-δ, due to its high oxygen ion conductivity, is widely used in composite electrodes. Here, a Ru-doped CeO2 catalyst (Ru@CeO2-δ, denoted as RC) as a second phase for composite electrode is synthesized through a co-precipitation method followed by calcination at 800 °C for 2 h. Rietveld refinement of X-ray diffraction (XRD) data confirms the pure-phase fluorite cubic structure, without any detectable RuO2 peaks (Fig. 2a and Supplementary Fig. 1), indicating the successful incorporation of Ru into the fluorite phase22. The refined XRD data reveal that the RC catalyst has an Fm-3m space group, with lattice contraction (a=5.4085 Å vs. 5.4110 Å for pristine CeO2), indicating that Ru atoms are successfully introduced into the bulk phase (Δr=0.08 Å) (Supplementary Fig. 2 and Supplementary Table 1)23. This lattice strain facilitates oxygen vacancy formation, as evidenced by Raman spectroscopy (Supplementary Fig. 3)24. Extended X-ray absorption fine structure (EXAFS) analysis (Fig. 2b) and atomic-resolution High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) (Fig. 2c) reveal the coordination environment of Ru. The Fourier-transformed (FT) EXAFS spectra exhibit a dominant scattering peak at 1.5 Å, attributed to Ru-O coordination with an average coordination number of 3.9 (Supplementary Table 2), while no discernible Ru-Ru scattering contributions observed at 2.4 Å (Supplementary Fig. 4)25. These structural features, corroborated by HAADF-STEM and energy-dispersive X-ray spectroscopy (EDS) elemental mapping results (Fig. 2d), conclusively confirm the atomic dispersion of Ru species within the CeO2 matrix.

We then combine RC with BSCF to form a composite electrode. The optimized BSCF: RC (70: 30 wt%; denoted as BSCF-RC30) composite preserves dual-phase integrity, as verified by Rietveld quantification (69.68 wt% perovskite, 30.32 wt% fluorite) (Fig. 2e and Supplementary Fig. 5)26. This suggests that BSCF-RC30 both maintains the integrity of the individual components and potentially benefits from the synergistic effect between the two phases. Structural analysis further reveals minimal variation in perovskite lattice parameters across BSCF, BSCF-C30 (Ru-free BSCF and CeO2 composite at a 70:30 mass ratio), and BSCF-RC30 (Supplementary Fig. 6), attributable to the negligible ionic radius mismatch between Ru4+ and host Co/Fe cations (Δr<0.03 Å). For the fluorite phase, BSCF-RC30 shows a pronounced lattice contraction (a=5.4059 Å) relative to both BSCF-C30 (a=5.4102 Å) and the RC catalyst (a=5.4085 Å) (Supplementary Table 1). This additional shrinkage is attributed to charge-compensating oxygen vacancies generated during the trapping of Ru, which shortens the average Ce-O bond length, facilitating the lowering of the proton hopping barrier and enhancing bulk proton conduction27. In addition, high-temperature XRD (HT-XRD) analysis is conducted to evaluate the thermal stability of the RC and BSCF-RC30 catalysts. As shown in Supplementary Fig. 7, both catalysts exhibit phase stability up to 800 °C, while the typical shift of the diffraction peaks is related to the materials’ thermal reduction. The lattice fringe spacings of 0.28 and 0.23 nm, corresponding to the (200) planes of the fluorite phase and the (111) planes of the cubic perovskite structure, respectively (Fig. 2f), are observed in STEM image. The strong contact and positive interactions between the two phases, as seen in STEM images, are expected to enhance electrode stability and improve ionic diffusion at the heterogeneous interface (Supplementary Fig. 8)28.

Crucially, the spatial redistribution of Ru species during interface formation is quantitatively mapped through EDS elemental analysis (Fig. 2g). Remarkably, Ru atoms migrate from the fluorite phase into adjacent perovskite domains during the co-sintering29. This thermodynamically driven spontaneous migration trend persists at multiple independent interfaces (Supplementary Fig. 9), demonstrating the thermodynamic preference for Ru doping into the perovskite lattice. Complementary spectroscopic investigations corroborate this spontaneous atomic trapping phenomenon and reveal the evolution of the coordination environment. Raman analysis identifies three critical signatures (Supplementary Fig. 10): (1) an intrinsic F mode at 465 cm−1, (2) emergent peaks at 700/970 cm−1 (strain-induced Ru-O-Ce distortion in RC), and (3) complete disappearance of fluorite-associated signals in BSCF-RC30, which is crystallographic evidence of Ru depletion from CeO22,30. The FT EXAFS spectra of BSCF-RC30 display a predominant Ru-O peak (Fig. 2h), whereas the trapped Ru atoms transition from pseudo-tetrahedral geometry in CeO2 (CNRu-O=3.9) (Fig. 2b) to near-ideal octahedral coordination in BSCF (CNRu-O=6.3) (Fig. 2i and Supplementary Table 2)31. The 61% increase in coordination number confirms the Ru integration into perovskite BO6 frameworks, constituting a structural reorganization with significant kinetic implications.

Building on the aforementioned experimental characterizations, we further elucidate the underlying Ru migration mechanism at the interface through the climbing image nudged elastic band (CI-NEB) method26,32. The atomic coordinates for the optimized computational model were provided in Supplementary Data 1. For better comprehension, the interface before and after BSCF capture of Ru from RC are labelled as BSCF-RC and BSCFR-C (BSCFR: BSCF matrix after capture of Ru, C: pure CeO2 phase), respectively. As illustrated in Fig. 2j, the Ru atom has a relatively low energy barrier of 1.14 eV to migrate from RC side to BSCF side, confirming the kinetic feasibility of the proposed Ru capture mechanism. Furthermore, Density functional theory (DFT) calculation also demonstrates that under the appearance of oxygen vacancy, the configuration with Ru atom on the surface layer (L1) is thermodynamically more stable (ΔE=−1.38 eV) than that with Ru atom on the subsurface layer (L2), providing driving force for the continuous Ru migration to the BSCF phase (Fig. 2k and Supplementary Fig. 11). To validate the critical role of oxygen vacancy concentration in atomic trapping dynamics, we systematically investigate a counterexample of La2NiO4 (LN)-RC interface33,34 with almost no oxygen vacancy, and reveals drastically different migration kinetics: Ru atom face a prohibitive energy barrier of 2.00 eV to migrate from RC side to LN side (Supplementary Fig. 12), higher than that in oxygen-deficient BSCF. This is consistent with STEM energy-dispersive X-ray spectroscopy (STEM-EDS) mapping observation of Ru retention in the CeO2 matrix after the same heat treatment (Supplementary Fig. 13). Parallel DFT/CI-NEB simulations and STEM-EDS mapping of oxygen-vacancy-rich BSCF-RC interface versus the vacancy-poor LN-RC analogue demonstrate that lattice vacancies provide both a low-barrier migration channel and the deep B-site trap needed for irreversible Ru capture30. Crucially, this thermodynamically driven atomic trapping strategy through oxygen-vacancy pathways, thereby enhancing the triple conductivity and interfacial activity of the composite catalysts. This mechanism is fundamentally different from conventional reverse-trapping techniques, such as the acid-base extraction of alkaline-earth ions (e.g., Sr2+ removal by MoO3), which rely primarily on surface effects29.

Heterogeneous interfacial effect
To comprehensively resolve the electronic and structural interplay arising from thermodynamically driven Ru migration, we employed a synergistic multiscale characterization framework integrating X-ray absorption spectroscopy (XAS), Aberration-corrected TEM, Electron energy loss spectroscopy (EELS), Time-of-Flight Secondary Ion Mass spectroscopy (ToF-SIMS), and DFT calculations. This collective evidences delineate an activation mechanism at the BSCF-RC30 heterointerface, mediated by localized electron redistribution, vacancy-coupled oxygen dynamics, and proton transfer modulation.

X-ray absorption near-edge structure (XANES) analysis provides critical insights into the valence state evolution of heteroatoms during interfacial Ru migration. Notably, although Co and Fe K-edge spectra remain indistinguishable between pristine BSCF and post-treated BSCF-RC30 (Fig. 3a, b)35,36, the Ru K-edge spectra reveal a striking valence shift from +3.34 in isolated RC nanoparticles (precursor phase) to +4.75 in the heterostructured system (Supplementary Fig. 14), which quantitatively demonstrates interfacial charge transfer during structural evolution (Fig. 3c)7,23. Also, this is confirmed by X-ray photoelectron spectroscopy (XPS) of Ru 3p orbitals for both RC and BSCF-RC30 (Supplementary Fig. 15). Notably, the retention of the near +4 state of the Ce ion (Supplementary Fig. 16) and unaltered Co/Fe coordination environments excludes bulk-phase compositional changes, pinpointing the interfacial region as the exclusive locus of electronic reconfiguration37,38.

Aberration-corrected STEM and EELS decrypt the spatial charge compensation dynamics underpinning. Atomic-resolution elemental analysis reveals strong chemical bonding between BSCF and RC phases, with Ru species preferentially stabilized at perovskite lattice sites (Supplementary Fig. 17)7. Line-scan profiles across the BSCF-RC interface identify two interrelated phenomena (Fig. 3d). First, the penetration of Ru atomic signal into BSCF domains at 2–3 nm (point A: Ru M-edge), accompanied by localized oxygen depletion (O K-edge) at the interface (point B), suggesting that oxygen vacancies are generated at the interface as a result of Ru migration9,27. A change in oxygen content is typically associated with variations in the oxidation states of metal ions due to charge transfer. As shown in Fig. 3e, f, EELS spectra of the Co L-edge and Ce M-edge further demonstrate charge transfer at the heterogeneous interface. The Co L-edge shifts to higher energies, indicating an increase in the oxidation state of Co, while the Ce M-edge shifts to lower energies, signifying a reduction in the oxidation state of Ce. The Ce 4+ characteristic peaks (883 eV) detected at positions E and F adjacent to the bulk RC phase exhibited enhanced intensity compared to those at the interfacial position D, thereby providing additional evidence for the valence state reduction of Ce at the interface39. These antiphase electron transfers establish interfacial dipolar fields that stabilize trapped Ru species while lowering the barrier for oxygen exchange, which is key to unlocking bifunctional ORR/OER activity.

Furthermore, TEM-EDS mapping of BSCF-RC30 collected after calcination at 600 °C for 200 h still detected Ru only in the BSCF domains (Supplementary Fig. 18). These results verify that atomic trapping arises from the combined thermodynamic driving force for Ru migration and the presence of robust anchoring sites (oxygen vacancies and strong Ru-O-Co/Fe covalent bonds) which lock Ru irreversibly after migration, rather than permitting the random, reversible diffusion expected from a purely chemical-potential-driven process40. The oxygen vacancy reservoir directly governs the catalyst’s oxygen handling capacity, so we quantified oxygen vacancy differences using thermogravimetric analysis (TGA) and iodine titration (Supplementary Fig. 19). TGA curves acquired in air (50–1000 °C) show that BSCF-RC30 (Ru-trapped composite) exhibits a higher total weight loss (~1.65%) compared to Ru-free BSCF-C30 (~1.40%). Since the RC phase itself is thermally stable (negligible mass loss), this additional weight loss in BSCF-RC30 directly reflects enhanced oxygen vacancy formation induced by Ru migration. In addition, O2 temperature-programmed desorption (O2-TPD) also verifies that BSCF-RC30 hybrid has a good oxygen dynamic exchange capacity than the BSCF-C30 composite (Supplementary Fig. 20). Iodometric titration show that the average oxygen non-stoichiometry (δ) of BSCF-RC30 increases by ~28% before and after thermal treatment (Supplementary Fig. 21). This increase is further confirmed by deconvoluted XPS O 1s spectra (Supplementary Fig. 22)41. The results show that the treated BSCF-RC30 exhibits an increase in the ratio of active oxygen species (O2−/O−) from 32% to 41% compared with the untreated BSCF-RC30, which is consistent with more oxygen vacancies at the heterointerface.

Hydration behavior of the interface, another critical performance determinant in reversible proton-conducting cells. ToF-SIMS analysis (Fig. 3g) reveals the distribution of CeO−, CoO−, and OH− ions in BSCF-RC30 after treatment at 500 °C for 5 h in humidified air (10% H2O-air)42. Notably, the RC phase exhibits a stronger hydration capacity than the BSCF phase, with the interface showing a particularly high concentration of OH− ions. The enhanced hydration performance at the interface is attributed to the increased oxygen vacancy content, which accelerates the introduction of protons. This selective hydration stems from synergistic contributions: the high surface area of nanoporous CeO2 domains (63.47 m2 g−1) provides abundant active sites for the hydration reaction, while the interfacial oxygen vacancies provide the dissociation sites (Supplementary Fig. 23). H2O-TPD analysis (Supplementary Fig. 24) shows that the RC catalyst displays the strongest H2O desorption peak compared to both BSCF-RC30 and BSCF oxides, further indicating effective hydration properties. The infrared (IR) spectroscopy (Supplementary Fig. 25) also confirms the presence of prominent hydroxyl signals at 3400 cm−1 in the RC and BSCF-RC30 samples, while bulk BSCF regions retain only weak adsorbed water signatures, providing additional evidence of the improved hydration capability of the BSCF-RC3043.

To further elucidate the activation mechanism and interface effect of the heterogeneous catalyst, we perform DFT calculations based on the experimental data and establish the BSCFR-C heterogeneous interfacial structure. The calculations confirm the strong interfacial coupling formed via charge transfer between Ce atoms in CeO2 phase and O atoms in BSCFR phase (Fig. 3h). Compared to the valence state in the bulk phases, this interfacial charge redistribution manifests as a coordinated increase in Co oxidation state with concurrent Ce valence drop at the phase boundary (Fig. 3i), establishing an electronic bridge that enhances catalytic functionality. Furthermore, compared to the hydration energy (Ehdr) of −1.00 eV and −0.07 eV in CeO2 and BSCFR bulk structure, BSCFR-C interface represents the minimum Ehdr of −1.13 eV (Fig. 3j), which aligns well with the EELS and ToF-SIMS findings (Fig. 3d–g)19,26,44. Proton conduction at heterogeneous interface is also explored (Supplementary Fig. 26). After Ru atom trapping, the proton migration energy of BSCFR (Hopping: 0.16 eV; Rotation: 0.18 eV) is notably lower than the pristine BSCF (Hopping: 0.52 eV; Rotation: 0.29 eV), realizing a low proton migration energy similar to that of CeO2 (Hopping: 0.22 eV; Rotation: 0.13 eV)26. Although the average proton migration barrier at the interface is slightly higher than that of the overall BSCFR and CeO2, the localized oxygen vacancy channel allows efficient Grothaus-type proton rotation with a migration barrier as low as 0.21 eV, which is close to that of the bulk phase (Fig. 3k and Supplementary Table 3)45,46.

For electrical conductivity, the BSCF-RC30 heterogeneous catalyst achieves an electrical conductivity of 28.5–38.2 S cm−1 at 300–800 °C (Supplementary Fig. 27), which retains the major metal conduction contribution and competitive with reported triple-conducting air electrodes, such as BCFZY (0.4–2.1 S cm−1 at 400–800 °C) and BaCo0.7(Ce0.8Y0.2)0.3O3-δ (BCCY) (1.3–4.3 S cm−1 at 400–800 °C). This multimodal conduction synergy transforms the engineered interface into a three-phase reaction hotspot, where the ORR and OER processes couple efficiently with proton diffusion, an active property not found in any of the individual components.

Electrochemical activity analysis
To evaluate the electrocatalytic performance of the derived air electrodes in R-PCECs, symmetric cells employing BZCYYb electrolyte were fabricated47. Although the single fluorite-structured RC contains atomically dispersed Ru active sites, its insufficient oxygen ion/electronic conductivity hinders charge transport, leading to a very large polarization resistance (Rp) (~2500 Ω cm2 at 550 °C) when used as a standalone air electrode, which is not sufficient to meet the performance requirements of R-PCECs (Supplementary Fig. 28). BSCF perovskite has high oxygen ion/electron conductivity, abundant oxygen vacancies, and its lattice can stabilize the migration of Ru atoms by forming Ru-O-B (B=Co/Fe) bonds. Upon compositing, BSCF’s conductivity compensates for RC’s charge transport limitations, while the interfacial active sites (generated by Ru trapping) synergize with BSCF’s oxygen vacancies to enhance oxygen reduction/evolution kinetics. EIS measurements of bare BSCF and BSCF-RCx (x=20, 30, 40 wt%) heterogeneous electrodes were conducted between 500 and 700 °C under dry air conditions. At 600 °C, the ASR of the bare BSCF electrode measures 0.52 Ω cm2, while BSCF-RC20, BSCF-RC30, and BSCF-RC40 electrodes exhibit reduced ASR values of 0.35, 0.28, and 0.36 Ω cm2, respectively (Supplementary Fig. 29). This systematic improvement in oxygen reduction reaction activity with RC catalyst incorporation highlights enhanced heterogeneous interfacial kinetics. Among these, the BSCF-RC30 electrode demonstrates optimal performance due to balanced RC content (30 wt%) and interfacial dispersion, maximizing synergistic effects. Under practical operating conditions (5% H2O-air at 600 °C), the BSCF-RC30 electrode achieved a remarkably low ASR of 0.08 Ω cm2, outperforming bare BSCF (0.28 Ω cm2) and other composites (BSCF-RC20: 0.18 Ω cm2; BSCF-RC40: 0.13 Ω cm2) (Supplementary Fig. 30). These results confirm that Ru atom trapping simultaneously enhances ORR activity and oxygen activation capability in humid environments. For completeness, we repeated the symmetric-cell test with an independently fabricated specimen (Cell #2). As shown in Supplementary Fig. 31, the resulting ASR values deviate by no more than ±6% from the previous ones over 500-700 °C, confirming that the low-resistance trend is highly reproducible and intrinsic to the atomic-trapping interface. Activation energy analysis derived from ASR-temperature dependence (500-700 °C) reveals values of 1.04, 1.02, 0.94, and 0.96 eV for bare BSCF, BSCF-RC20, BSCF-RC30, and BSCF-RC40 under 5% H2O-air conditions, respectively (Fig. 4a). The minimized activation energy (0.94 eV) of BSCF-RC30 underscores its enhanced catalytic efficiency and potential for low-temperature applications. The distribution of relaxation time (DRT) as an effective means to analyze the EIS was employed to further explore the electrochemical processes at the air electrode5,48. DRT deconvolution identifies three polarization contributions: high-frequency (>103 Hz, charge transfer), mid-frequency (10–103 Hz, surface ion exchange and bulk ion diffusion), and low-frequency (<10 Hz, surface mass transport) (Fig. 4b). Under 5% H2O-air at 600 °C, the MF resistance of the bare BSCF electrode is significantly higher than that of the heterogeneous electrodes with Ru atom trapping, indicating that the interfacial effects and synergistic interactions in BSCF-RC30 facilitate improved surface ion exchange and bulk phase diffusion. Furthermore, the BSCF-RC30 electrode exhibits effective surface mass transfer properties compared to other electrodes, further emphasizing the enhanced electrocatalytic performance due to atom trapping.

To confirm the promotion of catalytic activity through Ru atom trapping, additional electrodes were prepared for comparison. These included CeO2-modified BSCF (BSCF:C=70:30 wt%, denoted as BSCF-C30), Ru-doped Ba0.5Sr0.5(Co0.8Fe0.2)0.98Ru0.02O3-δ (BSCFR0.02), and BSCFR0.02-C30 (BSCFR0.02:C=70:30 wt%) air electrodes, and their electrochemical performances were tested using symmetric cells at 500 °C with 5% H2O-air (Supplementary Fig. 32). The corresponding ASRs are 1.22, 1.31, and 1.19 Ω cm2 for BSCF-C30, BSCFR0.02, and BSCFR0.02-C30, respectively, while the ASRs of BSCF and BSCF-RC30 electrodes are 1.50 and 0.57 Ω cm2. The BSCF-RC30 electrode exhibits an ASR of only 0.57 Ω cm2, markedly lower than BSCF-C30 (1.22 Ω cm2), BSCFR0.02 (1.31 Ω cm2), and the combined BSCFR0.02-C30 (1.19 Ω cm2). This pronounced advantage arises from Ru trapping, which establishes a highly active CeO2/BSCF heterointerface, while the efficient charge transfer and abundant vacancies at the interface not only provide active sites for hydration and oxygen activation, but also promote ion exchange between the two phases, which is an interfacial synergistic effect that neither simple CeO2 modification nor Ru pre-doping can provide. These results demonstrate that, while Ru doping or CeO2 composites can enhance BSCF’s electrochemical activity, the catalytic performance of simple CeO2 modification is far less significant than the RC catalyst. The improved performance of BSCF-RC30 is attributed to the trapping of Ru atoms, which not only enhances oxygen activation but also improves hydration and interaction kinetics between phases.

To assess the universality of the improved interfacial effect and catalytic activity, we extended the strategy to other Co-based air electrodes rich in oxygen vacancies, including PrBaCoO5+δ (PBC), BCFZY, SrCo0.8Fe0.1Nb0.1O3-δ (SCFN), and La0.6Sr0.4Co0.2Fe0.8O3-δ (LSCF), known for their marked oxygen activation ability. Indeed, elemental mapping of the as-prepared BCFZY-RC30 composite (Supplementary Fig. 33) likewise shows Ru exclusively anchored on the BCFZY side, confirming that oxygen-vacancy-assisted Ru trapping is both universal and robust across diverse perovskite hosts. EIS tests were performed on symmetric cells for hybrid electrodes compared to pristine counterparts (Supplementary Fig. 34). The ASRs of PBC-RC30, BCFZY-RC30, SCFN-RC30, and LSCF-RC30 electrodes are measured as 2.16, 2.00, 2.45, and 2.3 Ω cm2, respectively, under dry air conditions at 500 °C. These values show a significant decrease compared to their pristine counterparts, PBC (4.16 Ω cm2), BCFZY (4.38 Ω cm2), SCFN (4.90 Ω cm2), and LSCF (5.16 Ω cm2), with reductions of 48.1%, 54.3%, 50.0%, and 55.4%, respectively (Fig. 4c), confirming the efficacy of the atom trapping strategy in enhancing ORR activity across various perovskite-based materials wealthy in oxygen vacancies. Further EIS tests under a 5% H2O-air atmosphere (Supplementary Fig. 35) demonstrate that the catalytic activity of the air electrodes is significantly enhanced, particularly in humid air. For instance, the ASRs of the SCFN electrode are 0.10, 0.23, 0.49, 1.37, and 4.14 Ω cm2 at 700, 650, 600, 550, and 500 °C, respectively. In contrast, the ASRs of the SCFN-RC30 electrode at the same temperatures are found to be 0.04, 0.09, 0.16, 0.32, and 0.76 Ω cm2, respectively. This reduction in ASR by 60% at 700 °C and 81.6% at 500 °C further emphasizes the benefits of incorporating the RC catalyst for enhancing the electrochemical activity of the air electrode in humid conditions. Notably, LN-RC30 electrode shows no performance enhancement due to absent Ru atom trapping, emphasizing the critical role of this mechanism. The ASR of the LN electrode in humid air ranged from 0.18 to 8.10 Ω cm2 at 700–500 °C (Supplementary Fig. 36). However, lack of interfacial activity and poor ion transport are the main reasons for the LN-RC30 electrode to exhibit low electrochemical activity under dry and humid air conditions, with ASRs of 0.52 and 0.65 Ω cm2 at 600 °C (Supplementary Fig. 37).

As shown in Supplementary Fig. 38, the activation energies of PBC, BCFZY, SCFN, and LSCF air electrodes increase upon mixing with the RC catalyst compared to pristine counterparts under dry air conditions. Interestingly, under 5% H2O-air conditions, the activation energies of PBC-RC30, BCFZY-RC30, SCFN-RC30, and LSCF-RC30 air electrodes are lower than those of the pristine counterparts (Fig. 4d–g). The observed differences in activation energy patterns under dry and humid air conditions can be attributed to the varying contributions of the perovskite phase and the RC catalyst to the electrochemical reactions on the surface of R-PCECs air electrodes. In dry air, hybrid electrodes exhibit higher activation energies due to limited oxygen vacancy mobility in the RC phase. Conversely, humid conditions enable proton migration through CeO2-rich interfaces, reducing temperature dependence.

The stability of the electrodes is critical for ensuring long-term performance in R-PCECs. Stability tests on symmetric cells with BSCF and BSCF-RC30 electrodes were conducted in a 5% H2O-air atmosphere at 600 °C for up to 210 h. As displayed in Fig. 4h, the ASR of the BSCF electrode increases by 34.9% over the test period, with a growth rate of 4.81×10−4 Ω cm2 h−1. In contrast, the ASR of the BSCF-RC30 electrode shows no significant increase, demonstrating its evident catalytic stability. The stable performance of the BSCF-RC30 electrode is attributed to the formation of a protective layer of CeO2, which mitigates Ba and Sr segregation upon contact with water vapor (Supplementary Fig. 39). Additionally, CeO2 exhibits both effective water adsorption and hydration capabilities and effectively mitigates the Ba and Sr segregation resulting from the contact of the BSCF phase with water vapor. Not only that, but the BSCF-RC30 electrode behaves similarly stable under dry air (Supplementary Fig. 40). Further stability tests were conducted under severe thermal cycling conditions. As depicted in Fig. 4i and Supplementary Fig. 41, the ASR values of the BSCF-RC30 electrode increase by 0.06 and 0.25 Ω cm2 at 600 and 500 °C, respectively, over 48 cycles of 180 h. In contrast, the ASR values of the BSCF electrode under the same conditions increase by 0.12 and 0.73 Ω cm2 at 600 and 500 °C, respectively. The small increase in ASR of the BSCF-RC30 electrode during thermal cycling demonstrated the tolerance of the catalyst. This improved performance is primarily attributed to the BSCF-RC30 electrode’s stable phase structure, lower thermal expansion coefficient (TEC), and sustained catalytic activity (Supplementary Fig. 42). XRD patterns show that the BSCF-RC30 hybrid remains pure cubic and fluorite after sintering at 600 °C for 200 h, without the formation of RuO2 or alkaline earth metal second phases (Supplementary Fig. 43). Complementarily, XPS shows no shift in the Ru 3p binding energy, indicating an unchanged oxidation state, while TEM-EDS mapping reveals Ru confined to the BSCF lattice/heterointerface with negligible signal in the CeO2 domain (Supplementary Fig. 44, 18). The unchanged crystal structure, electronic state, and spatial distribution collectively demonstrate that Ru remains securely anchored, with no detectable segregation or reverse migration during extended operation, underscoring the long-term robustness of the atomic-trapping strategy.

The performance of R-PCECs in PCFC mode
The enhanced heterogeneous interface activity resulting from Ru atom trapping allows the BSCF-RCx air electrode to demonstrate competitive performance in R-PCECs. To evaluate the impact of different composite ratios on the performance of the air electrodes, button cells with the geometry Ni-BZCYYb|BZCYYb|BSCF-RCx (x=0, 20, 30, and 40 wt%) were fabricated. Dry hydrogen and humid air were used as the proton and oxygen sources for the fuel electrode and air electrode, respectively (Supplementary Fig. 45). The current-voltage polarization and power density curves for the R-PCECs with BSCF-RCx (x=0, 20, 30, and 40 wt%) electrodes in PCFC mode are presented in Fig. 5a and Supplementary Fig. 46. Among them, the button cell with BSCF-RC30 electrode obtain peak power densities (PPD) of 1.51, 1.10, 0.82, and 0.54 W cm−2 at 650, 600, 550, and 500 °C, respectively (Fig. 5a). The single-cell performance was likewise reproduced on a second device prepared under identical conditions. Peak power densities differ by only 2.2% at 650 °C and <3% at 550 °C (Supplementary Fig. 47), underscoring the robustness and reliability of the composite-electrode design. Illustrated in Fig. 5b, the EIS of this cell, measured at open circuit voltage (OCV), shows remarkably low Rp values of 0.09, 0.13, 0.18, and 0.35 Ω cm2 at 650, 600, 550, and 500 °C, respectively. These findings strongly suggest that the high ORR activity of the air electrode contributes significantly to the performance of this cell. In fuel cell mode on R-PCECs, the performance of the BSCF-RC30 electrode rivals that of recently reported top-tier air electrodes such as PrNi0.5Co0.5O3-δ (PNC), BCCY, and PBSCF (Supplementary Table 4). Notably, BSCF, BSCF-RC20, BSCF-RC30, and BSCF-RC40 air electrodes exhibit PPDs of 0.75, 0.92, 1.51, and 1.17 W cm−2 at 650 °C in PCFC mode, respectively, demonstrating that incorporating RC significantly enhances the performance of BSCF. The addition of RC catalyst, with complex impacts on conductivity, surface mass transfer, ion conduction, and hydration rate of the air electrode, leads to non-linear performance increases as the mass ratio of RC in the electrode rises. At 650 °C, the BSCF-RC30 cell exhibits a performance improvement of over 100% compared to employing a BSCF electrode (Fig. 5c). As the operating temperature decreases, the performance advantage of the BSCF-RC30 cell becomes more apparent, particularly at 500 °C, where it outperforms cells utilizing BSCF, BSCF-RC20, and BSCF-RC40 air electrodes by approximately 121%, 86%, and 41%, respectively.

The EIS analysis of the cells with various air electrodes provides insights into the differences in ORR activity of the BSCF-RCx cells at test temperatures (Supplementary Fig. 46). The corresponding Rp to temperature dependence is shown in Fig. 5d. As anticipated, the air electrode BSCF-RC30 demonstrates the lowest activation energy at 0.55 eV, while the activation energies of BSCF, BSCF-RC20, and BSCF-RC40 are 1.01, 0.61, and 0.58 eV, respectively. These results emphasize that heterogeneous interfaces with abundant active sites and favorable ionic conduction rates can significantly mitigate the temperature dependence of air electrodes, enhancing the potential for R-PCECs at lower operating temperatures. Concerning the electrode stability in fuel cell mode for R-PCECs, Fig. 5e portrays the endurance of the button cell featuring BSCF-RC30 and the bare BSCF electrodes operating at a constant current density of 400 mA cm−2 at 600 °C. Over the 350-hour testing period, the voltage of the BSCF cell declines notably from 0.82 to 0.53 V. This decline is primarily attributed to the significant degradation of BSCF electrode performance, induced by extensive Ba and Sr segregation during the prolonged testing period and inadequate thermal compatibility with the BZCYYb electrolyte (Supplementary Fig. 48). In contrast, the cell equipped with the BSCF-RC30 electrode displays no apparent deterioration throughout the 400-hour stability test, affirming the competitive thermochemical robustness of the heterogeneous electrode created.

Electrolysis performance and tolerance
R-PCECs featuring varying composite ratios of air electrodes were subjected to testing in PCEC mode, with the air electrode and fuel electrode exposed to 10% H2O-air and dry H2, respectively. Figure 6a displays the current-voltage polarization curves for the cell employing BSCF-RC30 as the air electrode. At a voltage of 1.3 V, the cell’s current densities reach −2.21, −1.49, −1.07, and −0.58 A cm−2 at temperatures of 650, 600, 550, and 500 °C, respectively. Consistent with the electrochemical activity results obtained from symmetric cells testing each electrode independently, the BSCF-RC30 electrode outperforms others in the PCEC mode within the specified temperature range. For instance, at 650 °C, the performances of the BSCF cell, BSCF-RC20 cell, and BSCF-RC40 cell are −0.98, −1.13 and −1.73 A cm−2, respectively. and the BSCF-RC30 cell outperforms them by 125%, 96% and 28%, respectively (Fig. 6b and Supplementary Fig. 49). In addition, the EIS of the cell with the BSCF-RC30 electrode in the PCEC mode is measured at 1.3 V. The resulting Rp values are 0.08, 0.14, 0.24, and 0.43 Ω cm2 at 650, 600, 550, and 500 °C, respectively, underscoring the high OER activity of the BSCF-RC30 electrode (Fig. 6c). According to our knowledge, the BSCF-RC30 electrode exhibits competitive performance in PCEC mode compared with other reported advanced air electrodes (Supplementary Table 4)46,49–57.

For R-PCECs, assessing the air electrode’s electrocatalytic performance involves conducting stable cyclic operations between PCFC and PCEC modes at various voltages. In this study, we subjected R-PCECs with a BSCF-RC30 electrode to cycling tests at 650 °C, employing voltages in the sequence of 0.8, 1.1, 0.9, 1.2, 1.0, and 1.3 V in both fuel cell and electrolysis cell modes, each for a duration of 10 min. As depicted in Fig. 6d, the cell exhibits consistent and stable performance over 96 cycles, with no noticeable degradation observed in either the fuel cell or electrolysis cell mode. This suggests that the BSCF-RC30 air electrode maintains stable electrocatalytic activity and suitability for application. Additionally, Supplementary Fig. 50 illustrates our examination of the step-like short-term stability of the BSCF-RC30 cell in PCEC mode at 600 °C, revealing competitive stability at voltages of 1.1, 1.2, and 1.3 V.

Faraday efficiency (FE) plays a pivotal role in determining the energy conversion efficiency and economic viability of R-PCECs. However, practical electrolysis processes involve various loss mechanisms, such as electrode reactions, electrolyte ion transport, and impurity reactions, resulting in actual chemical reactant conversion rates falling short of theoretical maxima. The hydrogen production rate and FE of R-PCEC with BSCF-RC30 air electrode were assessed and calculated under the conditions of 600 °C and 10% H2O-air (Fig. 6e). With an increase in current density from −0.8 to −1.4 A cm−2, the hydrogen production rate rises from 4.12 to 8.06 mL cm−2 min−1. Notably, the FE peaks at 88.28% at a current density of −1.2 A cm−2, maintaining around 80% even with further increases in current density. This underscore that a higher FE indicates the potential of R-PCECs for efficient energy conversion, reduced energy consumption, and the promotion of sustainable energy development.

Typically, R-PCECs function within the medium temperature range of 400–650 °C. However, the cyclic variation between high and low temperatures during operation induces continual shifts in internal thermal stress and thermal expansion within the cells. This can lead to issues such as rupture, delamination, or damage, ultimately resulting in degradation or complete failure of the cells. Hence, ensuring robust thermal stability is imperative for extending the lifespan of R-PCECs. Here, we assess thermal stability by monitoring voltage changes in R-PCECs under PCFC (0.4 A cm−2) and PCEC (−0.8 A cm−2) modes during temperature cycling. As demonstrated in Fig. 6f, the cell maintains stable voltage levels through 10 cycles at 500, 550, and 600 °C in both PCFC and PCEC modes. This suggests that the air electrode not only exhibits a stable crystal structure and catalytic activity but also demonstrates competitive compatibility with the electrolyte. Furthermore, the tolerance and long-term stability of R-PCECs under high water pressure were investigated. The cell demonstrates stable operation for 500 h at a current density of −1.4 A cm−2 and 600 °C (Fig. 6g). Notably, during the 200-hour test under 50% high water pressure, the cell exhibits stable operation without significant performance degradation. This indicates that the BSCF-RC30 composite air electrode boasts electrochemical stability and tolerance. To elucidate the microscopic origin of this steam tolerance, we next examined whether deleterious surface species develop on the air electrode under humid air conditions. To clarify why the BSCF-RC30 electrode withstands high steam pressures so well, we examined its surface after 24 h in 10% H2O-air at 600 °C. BSCF-RC30 retains a clean morphology, showing no detectable carbonate or hydroxide deposits, which is attributed to the strong interaction of the heterogeneous interface that inhibits the outward migration of alkaline earth metals from the BSCF bulk phase and the high-surface-area CeO2 wrapped on the surface to avoid contact with H2O/CO2 (Supplementary Fig. 51). In contrast, pristine BSCF forms flaky carbonates after air calcination and densely distributed hydroxide nanoparticles after prolonged wet treatment, which can shorten the triple-phase boundary and slow down proton transfer58.

Methods
Materials
Ba(NO3)2, Sr(NO3)2, Co(NO3)2⋅6H2O, Fe(NO3)3⋅9H2O, Zr(NO3)4⋅5H2O, Ce(NO3)3⋅6H2O, Gd(NO3)3·6H2O, Nb(NO3)5, La(NO3)3·6H2O, Ni(NO3)2·6H2O, ethylenediaminetetraacetic acid (EDTA), citric acid (CA) monohydrate, isopropyl alcohol, ethylene glycol, glycerol, and ammonia were all analytically pure and purchased from Sinopharm Chemical Reagent Co., Ltd. Pr(NO3)3·6H2O, Y(NO3)3⋅6H2O, and Yb(NO3)3⋅5H2O with a purity of 99.9% were purchased from Shanghai Aladdin Biochemical Co., Ltd. The Ru-based purity of the ruthenium trichloride aqueous solution was 35.0–42.0%, also purchased from Shanghai Aladdin Biochemical Co., Ltd.

Synthesis of catalysts
The powders of BSCF, PBC, LSCF, SCFN, BCFZY, LN, and BZCYYb were all prepared via a citrate-EDTA sol-gel route. All powders/pellets were stored at room temperature under dehumidified conditions (desiccator) until use. In a typical run (illustrated here for BSCF), stoichiometric amounts of Ba(NO3)2, Sr(NO3)2, Co(NO3)2·6H2O, and Fe(NO3)3·9H2O were first dissolved in deionized water. Separately, citric acid and ethylenediaminetetraacetic acid were co-dissolved in ammonia; this chelating solution was then added to the metal nitrate mixture, and the pH was adjusted to ~7. The overall molar ratio of CA:EDTA:total metal ions was maintained at 2:1:1. Under continuous stirring and gentle heating, the suspension gradually transformed into a clear gel. This gel was dried in an oven at 180 °C to remove residual water, yielding a fluffy precursor powder, which was finally calcined in air at 1000 °C for 5 h to afford the crystalline perovskite oxide. The Ru@CeO2-δ catalyst was obtained by a co-precipitation procedure. Ce(NO3)3·6H2O was dissolved in water at 80 °C, and concentrated ammonia was added dropwise until the solution reached pH ~9, inducing cerium hydroxide precipitation. Then, the corresponding RuCl3 solution was added at a metal atomic ratio of 5% and stirred at 80 °C for 3 h. After cooling to room temperature, the slurry was aged for 12 h, then subjected to three cycles of centrifugation and ultrasonic washing to remove residual chloride. The washed precipitate was dried and finally calcined at 600 °C for 3 h to yield the RC catalyst powder. All composite electrodes were prepared by ball-milling at 400 rpm for 30 min. BSCF and RC were mixed at 80:20, 70:30, and 60:40 (wt%) to obtain BSCF-RC20, BSCF-RC30, and BSCF-RC40, respectively. Likewise, PBC, LSCF, SCFN, BCFZY, and LN were each combined with RC at 70:30 (wt%) to yield M-RC30 composites (M denotes the corresponding oxide). In addition, BSCF and BSCFR0.02 were each combined with CeO2 at 70:30 (wt%) to produce BSCF-C30 and BSCFR0.02-C30 electrodes.

Cells fabrication
The hydrogen-electrode support was first produced by die-casting a blend of BZCYYb, NiO, and starch (mass ratio 6.5: 3.5: 1). A 0.35 g portion of this powder was then compacted into a 15 mm-diameter disk under 100 MPa via uniaxial pressing. These green pellets were sintered in air at 1000 °C for 2 h to yield the pre-treated support layer. To deposit the BZCYYb electrolyte, 1 g of its powder was dispersed in a mixture of 10 mL ethanol, 2 mL ethylene glycol, and 0.1 g ethyl cellulose, and milled at 400 rpm for 60 min. The resulting slurry was drop-cast onto the support by spin-coating at 8000 rpm for 30 s, repeating this cycle three times. The coated half-cells were then fired at 1450 °C for 5 h to form the dense electrolyte. For full-cell assembly, an air-electrode ink was prepared by ball milling the electrode powder with isopropanol, ethylene glycol, and glycerol in a ratio of 1 g: 10 mL: 2 mL: 1 mL. This slurry was sprayed onto the electrolyte side of each half-cell at 150 °C, then calcined at 800 °C for 2 h, giving a functional cell with an air-electrode area of 0.28 cm2. Symmetric cells were made by die-casting 0.4 g of electrolyte powder containing 1 wt% NiO sintering aid into disks and sintering at 1450 °C for 10 h to obtain dense pellets. The dense pellets were polished, then sprayed with electrode paste and sintered (800 °C, 2 h) to complete a symmetrical configuration with a single-sided air-electrode area of 1.06 cm2. The air-electrode loading in both single and symmetric cells after calcination was approximately 20 mg cm−2. Finally, silver wire and silver glue (DAD-87, Shanghai, China) were used as the current collectors and sealant. Data for replicated work have come from parallel experiments using the same electrode preparation method.

Characterizations
X-ray diffraction patterns were collected on a Rigaku Smartlab diffractometer using Cu-Kα radiation (λ=1.5418 Å), and the diffraction data were refined in TOPAS-4.2 to extract precise lattice constants. F20 STTEM/STEM was employed to observe the microstructure and crystal structure of the samples, and HR-TEM, Mapping and STEM at 200 kV were used to capture images. Structural features of the cells were examined using a Hitachi S-4800 SEM. The size change of the air electrode materials in air was detected by Netzsch 402 C/3/G and TEC were calculated. EELS data were acquired using a Hitachi HT7700 operated at 120 kV. BET and XPS tests were carried out using a Quantachrome Autosorb-iQ3 and a PHI5000 VersaProbe spectrometer, respectively. TGA (Model STA449F3, NETZSCH) and O2-TPD (MS, Hiden, HPR20) were used to measure the mass loss and oxygen volatilization of the dried powder samples. ToF-SIMS (ToF-SIMS 5 iontof), H2O-TPD (MS, Hiden, HPR20), and IR (IS50 FTIR) were used to analyze the hydration capacity of the samples. Samples were pre-treated in a 10% H2O-air atmosphere at 500 °C for 5 h, followed by natural cooling to room temperature. Spectra were acquired at room temperature in static ambient air using a Fourier-transform IR spectrometer. For the H2O-TPD test, samples were heated from room temperature to 1000 °C at a rate of 3 °C min−1 under flowing nitrogen (N2, 50 mL min−1). Desorbed water species (H2O, m/z=18) were monitored in real time using a mass spectrometer. Finally, XAS were recorded at a soft-X-ray beamline of the Taiwan Synchrotron to investigate local electronic and structural environments.

Electrochemical measurements
EIS was carried out on symmetric cells at open circuit voltage using a voltage amplitude of 10 mV over the frequency range from 106 to 0.01 Hz. All measurements used a 100 mL min−1 gas flow, and humidity was regulated via a high-pressure constant-current pump (LC100 HPLC). Data were recorded on a Solartron 1287 potentiostat coupled with a 1260A frequency response analyzer. Single-cell evaluation was performed by sealing the cell in a quartz tube with silver adhesive: the fuel electrode faced the tube’s interior (dry H2 at 80 mL min−1) while the air electrode was exposed to humidified air (10 vol% H2O, 100 mL min−1). I-V-P and I-V characteristics in both fuel-cell and electrolysis modes were measured using a Keithley 2420 source meter. For EIS in fuel-cell mode, spectra were acquired at OCV between 106 and 0.1 Hz with a voltage amplitude of 20 mV; in electrolysis mode, impedance was recorded at 1.3 V from 106 down to 0.1 Hz with a voltage amplitude of 100 mV. All impedance data collected and processed using Zview software. The polarization curves and stability tests in R-PCECs were obtained without iR-correction. For the Faraday efficiency test, a single cell employing the same current collection and sealing method was used. The fuel electrode was exposed to hydrogen during temperature rise and converted to argon when the OCV was stabilized, while the air electrode was constantly exposed to a 10% H2O-air environment. The cell was tested for FE at a constant current density (between −0.8 and −1.4 A cm−2 in gradual increments of −0.1 A cm−2). The gas on the effluent side was analyzed in real time by a GC-9860-5CNJ gas chromatograph equipped with a thermal conductivity detector after 30 min of stabilization after the current was applied. To determine electrical conductivity, 0.8 g of air-electrode powder was uniaxially pressed at 100 MPa into a 15 mm × 5 mm rectangular bar, then sintered in air at 1130 °C for 10 h. Conductivity was measured by the DC four-probe method under flowing air (100 mL min−1) over the temperature range 300–800 °C. Except for the BSCF-RC30 electrode, which was tested in both parallel symmetrical cells and single cells, all other electrochemical tests and comparisons were performed only once.

DFT calculation details
All DFT calculations were performed in Vienna Ab initio Simulation Package (VASP)59. Plane-wave pseudopotentials were constructed with the full potential projector augmented wave (PAW) method60, and the exchange correlation was treated with the generalized gradient approximation (GGA) Perdew-Burke-Ernzerhof (PBE) functional61, with spin polarization and aspherical contributions both considered. A simplified rotationally invariant approach to DFT+U was also applied62, with Hubbard U-value of 6.3, 5.3, 3.3, and 5 for Ni, Fe, Co, and Ce, respectively. In total, 4 bulk phases are included in DFT calculation, namely, BSCF, BSCFR, C, and LN. A cutoff energy of 500 eV was used during the geometry optimization. The energy convergence criterion of 10−5 eV and the force convergence criterion of 0.02 eV/Å were used for electronic and ionic steps, respectively. The K-mesh was 4×4×2 for LN and 4×4×4 for other bulk phases (Supplementary Table 5). Hydration energy (Ehdr) of BSCF and BSCFR are based on the unit cell, while that of C and LN are based on the 2×2×2 and 2×2×1 supercells, respectively. The energies of these simulation cells (i.e., unit cells or supercells) were labeled as Ecell. For a typical hydration reaction shown in following Eq. (1)63, Ehdr can be obtained with Ecell-H, Ecell, EH2O denoting the energy of simulation cell with an extra proton, isolated water, and oxygen gas molecule using the following Eq. (2):

H2O + Oox + V(cid:1)o(cid:1) = 2OH(cid:1)o (1)

Ehdr = 2Ecell-H − Ecell − Edef − EH2O (2)

Due to the very limited amount of proton, the charge on all simulation cells were kept neutral. The energy barriers of proton migration (i.e., hopping or rotation) were determined by CI-NEB method. All surface slabs and interface structures constructed along the z-direction, with a vacuum layer over 15 Å. Before the construction of interface structure, 4-layer surface slabs of C(110), BSCF(110), and LN(103) were first cleaved. After that, BSCF-C interface structure was built based on BSCF(110) surface slab and 2×2 supercell of C(110) surface slab, while LN-C interface structure was based on 2×1 supercell of LN(103) and 2×5 supercell of C(110) surface slab. The K-meshes of 2×1×1 and 1×1×1 were used for BSCF-C and LN-C interfaces, respectively, and the corresponding lattice parameters and mismatches were provided (Supplementary Table 6). During the geometry optimization and the following calculation of Ehdr and proton transport energy barriers in these interfaces, only atoms in the two interface layers were allowed to move, while atoms in other layers were fixed. Also, loose convergence criteria of 10−5 eV and 0.05 eV/Å were used due to the high computational cost of interface structures. Regarding Ru migration across the interface, Ru atom was placed as a 4-coordinated interstitial atom in the side of C and as a substitution atom of Co in the BSCF side, respectively. The corresponding two configurations were used as the starting and ending points in the CI-NEB calculation of Ru migration energy barriers.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript reports a thermodynamically driven “atomic-trapping” strategy that reorganizes the heterointerface between Ru-doped CeO2 (RC) and Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF), yielding an ultralow-Ru composite air electrode for reversible protonic ceramic electrochemical cells. By exploiting heat-induced Ru migration, the authors create a strongly coupled interface and enrich the interface with oxygen vacancies, accelerating proton transport in the bulk phase. The resulting electrode delivers 1.51 W cm-2 in fuel-cell mode and -2.21 A cm-2 at 1.3 V in electrolysis mode at 650 °C, with excellent stability and tolerance under complex conditions that surpass state-of-the-art triple-conducting electrodes. In my opinion, the work meets Nature Communications’ criteria for conceptual novelty, mechanistic insight, and practical relevance. Therefore, I recommend accepting after some revisions once the points below are addressed.\n1. The atomic-trapping mechanism is fascinating. Would you consider adding a concise comparison with the previously reported reverse-trapping phenomena? A brief explanation of the driving forces could further underline the originality of your approach.\n2. The manuscript states RC as a single-atom catalyst. Could you explain the key evidence supporting atomic dispersion? If this hypothesis is valid, using the Ru@CeO2-δ labeling method may convey this concept more intuitively.\n3. To gain a more comprehensive understanding of the influence of atomic-trapping on crystal structure, it is useful to provide Rietveld-refined XRD patterns and lattice parameters for BSCF and BSCF-C30.\n4. Given that many single-atom catalysts excel in low-temperature O2 activation, it is recommended to consider testing symmetric-cell or half-cell data for a pure RC electrode to justify why the composite route is superior. At the same time, a succinct explanation of the benefits of choosing BSCF as the perovskite matrix will help highlight the synergistic effect of the RC and BSCF pairing.\n5. For clarity and reproducibility, it may be helpful to (a) define abbreviations such as BSCF-RC, C, BSCFR-C, and BSCFR at first mention, and (b) provide key experimental details (e.g., gas composition, heating rate, treatment temperature, etc.) for the IR and H2O-TPD measurements in the Methods section.", "reviewer_2": "The study presents a detailed investigation into using \"atomic-scale interface engineering\" through Ru migration to enhance the electrochemical performance of protonic ceramic electrochemical cells (R-PCECs). The approach is well-motivated, and the experimental and computational analyses are comprehensive. Also, the manuscript is generally well-written and supported by extensive characterization and theoretical modeling. While the manuscript is suitable for publication, a few points would benefit from clarification and minor revision.\n1. The role of oxygen vacancies is central to the proposed mechanism. If available, quantifying oxygen non-stoichiometry before and after Ru trapping (e.g., through TGA or iodometric titration) would strengthen the interpretation.\n2. A brief discussion on the long-term stability of the trapped Ru species under operational conditions would be helpful. In particular, clarifying whether any redistribution, segregation, or structural changes of Ru were observed after extended electrochemical testing would be helpful.\n3. The hydration performance is well-discussed. However, it would add value to mention whether any surface species (e.g., hydroxyls, carbonates) form under operating conditions and how they might influence electrode behavior, potentially supported by relevant literature.\n4. Including standard deviations or a brief discussion on the reproducibility of key electrochemical measurements (e.g., ASR, PPD) across multiple samples would help evaluate data reliability.\n5. In Figure 3h, including an iso-value label and plane orientation in the differential charge density plot would make it clearer.\n6. For CI-NEB calculations, such as Figures 2j, 3k, and Supplementary Figure S11b, the x-axis label “transport index” appears unclear. It would be more appropriate to use “reaction coordinate” or “image number.” Furthermore, for clarity, please explicitly indicate the reactant, transition state, and product on each plot and ensure a consistent number of intermediate images across cases. The curve smoothing in Figure S11b may introduce artifacts.", "reviewer_3": "This manuscript from Liu et al presents a compelling strategy for enhancing the performance of air electrodes for reversible protonic ceramic electrochemical cells (R-PCECs). The authors describe a phenomenon termed “atomic trapping,” where Ru species, initially doped into a CeO2 support, are shown to migrate into the lattice of a BSCF perovskite during thermal processing. This process is reported to create a highly active and stable heterointerface, leading to remarkable bifunctional performance in both fuel cell and electrolysis modes. The work combines detailed electrochemical characterization with advanced microscopy, spectroscopy, and DFT calculations to elucidate the underlying mechanisms, which are attributed to optimized charge redistribution, interfacial oxygen vacancy formation, and enhanced triple-phase boundary kinetics.\nThe work is of high quality and the results are impressive however, there are several key points that should be addressed to strengthen the manuscripts conclusions before it can be considered for the publication.\n1. The sentence starting from line 34 in the abstract section is lengthy and missing key vital points. For instance, the authors did not indicate the temperature at which the peak power density in the FC mode and current density in the EC mode were achieved. It is recommended that the sentence be simplified and clarified to avoid misleading readers into assuming it was obtained at 600 oC.\n2. The central claim revolve around “atomic trapping”. While the evidence for Ru migration is strong, the term “trapping” implies a specific dynamic mechanism that may not be fully supported by the data, which could also be interpreted as standard thermodynamically driven diffusion. The authors must provide a more robust justification for using this specific term over a simpler description or provide more direct evidences of a “trapping” phenomenon.\n3. The link between Ru migration and presence of oxygen vacancies in the perovskite is critical point. The LN-RC counterexample is illustrative, but the argument would be more convincing if the role of oxygen vacancies as the primary driving force was more definitively established, beyond just correlation.\n4. “Theses” in line 189 is a typo.\n5. The authors claim this strategy is a “universal paradigm” for Co-based perovskites rich in oxygen vacancies. To support this, they demonstrate enhanced performance by adding RC to PBC, BCFZY, SCFN, and LSCF electrodes (Figure 4c). However, the manuscript lacks direct evidence that the same Ru migration mechanism occurs in these other systems. Without supporting STEM-EDS or spectroscopic data for these other composites, the assertion that performance enhancement stems from the same “atom trapping” mecanism remains assumption. To validate such a strong claim of universality, direct evidence of Ru migration should be provide for at least one other perovskite system.\n6. Is the total Ru loading in the BSCF-RC30 electrode comparable to the Ru content in the BSCFR0.02 electrode?\n7. Why the in-situ migration of Ru from CeO2 creates a more active interface than simply pre-doping Ru into BSCF lattice? Is it the formation of a specific interfacial structure that cannot be achieved by other means? Elaborating on this point would significantly improve impact of this great work.\n8. There appear to be a contradiction in the Ea trends reported. For the BSCF-RCx series, the BSCF-RC30 electrode shows the lowest Ea (0.94 eV) in dry air (lines 305-308). However, for the other perovskite systems (PBC, BCFZY, etc.), the text state that the hybrid electrodes exhibit higher activation energies in dry air compared to their pristine counterparts (lines 352-354 and Supp. Fig. 32). The explanation provided, that limited oxygen vacancy mobility in the RC phase increases Ea, does not account for why the BSCF-RC30 system behaves differently. This inconsistency should be clarify. Is there some unique interaction at the BSCF-RC interface that is not present in the others?\n9. For Fig 5., please fix the error with the captioning of the a, b, and c parts."}, "author_response": "Point-to-Point Responses to Reviewers’ Comments and Suggestions\nFirstly, we would like to express our gratitude for the time and effort the reviewers dedicated to reviewing our manuscript (NCOMMS-25-40595-T). We also highly appreciate the thoughtful comments and suggestions provided by the reviewers, which have substantially enhanced the quality and clarity of our manuscript. We have carefully revised the manuscript to thoroughly address all the comments and concerns raised. The modifications have been marked in yellow in the revised manuscript. Our point-to-point responses are presented below.\nReviewer #1 (Remarks to the Author):\nThis manuscript reports a thermodynamically driven “atomic-trapping” strategy that reorganizes the heterointerface between Ru-doped CeO (RC) and Ba Sr Co Fe O (BSCF), yielding an ultralow-Ru composite air electrode for reversible protonic ceramic electrochemical cells (R-PCECs). By exploiting heat-induced Ru migration, the authors create a strongly coupled interface and enrich the interface with oxygen vacancies, accelerating proton transport in the bulk phase. The resulting electrode delivers 1.51 W cm-2 in fuel-cell mode and -2.21 A cm-2 at 1.3 V in electrolysis mode at 650 °C, with excellent stability and tolerance under complex conditions, that clearly surpass state-of-the-art triple-conducting electrodes. In my opinion, the work meets Nature Communications’ criteria for conceptual novelty, mechanistic insight, and practical relevance. Therefore, I recommend accept after minor revision once the points below are addressed.\nResponse:\nWe express our sincere gratitude to the reviewer for providing us with detailed and insightful comments. These comments have greatly contributed to enhancing the quality of this work. We have made every effort to address these comments in the revised version of our manuscript. Please find our responses to the comments point by point as follows.\nComment 1:\nThe atomic-trapping mechanism is fascinating. Would you consider adding a concise comparison with the previously reported reverse-trapping phenomena? A brief explanation of the driving forces could further underline the originality of your approach.\nResponse to C1:\nWe gratitude to the reviewer for the valuable comment. We are pleased that a comparison with previously reported reverse-trapping phenomena would strengthen the mechanistic understanding and emphasize the uniqueness of our strategy. In prior work, Zhuang et al. [Nat. Catal. 2022, 5, 300-310] proposed a “reversely trapping atoms” mechanism, in which Sr atoms were selectively extracted from the surface of (La Sr ) Co Fe O (LSCF) via a solid-state acid-base reaction with acidic MoO under high-temperature conditions. The formation of SrMoO and subsequent Sr-deficient LSCF surface resulted in altered covalent bonding and increased oxygen vacancy concentration, thus enhancing oxygen redox activity. This mechanism is essentially driven by a strong acid-base interaction, where the acidic nature of MoO drives the selective deintercalation of alkaline-earth elements (e.g., Sr) from the perovskite A-site [Zhao et al., Adv. Mater. 2024, 36, 2405052]. In contrast, the atomic trapping mechanism we propose is mainly controlled by thermodynamic stabilization and oxygen vacancy-mediated migration. Rather than being driven by surface reactivity or acid-base reactivity, the Ru atoms originally residing in the fluorite phase (Ru@CeO ) are thermally activated and migrate into the perovskite lattice Ba Sr Co Fe O (BSCF), seeking energetically favorable coordination environments. This migration is facilitated by the presence of abundant oxygen vacancies in BSCF, which serve both as migration channels and low-energy trapping sites. As confirmed by our DFT calculations, the migration barrier for Ru from RC to BSCF is only 1.14 eV, and the presence of surface/subsurface oxygen vacancies further stabilizes Ru occupation at B-sites in the perovskite lattice (see Fig. 2j, k). Unlike the surface-limited effect in reverse-trapping, our mechanism leads to a bulk-interface level reconstruction, forming covalent metal-oxygen bridges and enabling sustained enhancement in both electronic structure and ion conduction.\nRelated discussions (yellow-highlighted in Line 190-199, Page 8-9) have been added in the revised manuscript as follows:\n“Parallel DFT/CI-NEB simulations and STEM-EDS mapping of oxygen-vacancy-rich BSCF-RC interface versus the vacancy-poor LN-RC analogue demonstrate that lattice vacancies provide both low-barrier migration channel and the deep B-site trap needed for irreversible Ru capture. Crucially, this thermodynamically driven atomic trapping fundamentally differs from reverse-trapping phenomena reported elsewhere. Unlike acid-base extraction of alkaline-earth ions (e.g., Sr2+ removal by MoO ) which generates surface vacancies and secondary phase, the present strategy harnesses bulk migration of Ru from the fluorite into the perovskite lattice29. Oxygen-vacancy pathways and thermodynamic stabilization at B-sites drive heterointerface reconstruction, thereby enhancing triple conductivity and interfacial activity, distinct from the surface effects in conventional reverse trapping techniques.”\nComment 2:\nThe manuscript states RC as a single-atom catalyst. Could you explain the key evidence supporting atomic dispersion? If this hypothesis is valid, using the Ru@CeO labeling method may convey this concept more intuitively.\nResponse to C2:\nWe sincerely thank the reviewer for their thoughtful comment and for raising an important point regarding the atomic dispersion of Ru in the RC phase. We fully agree that using the Ru@CeO labeling would provide a clearer and more intuitive presentation of the concept, and we will incorporate this term into the revised manuscript.\nTo address the reviewer's concern, we have gathered several key lines of evidence supporting the atomic dispersion of Ru in the CeO matrix. Firstly, the Extended X-ray Absorption Fine Structure (EXAFS) analysis reveals the dominant scattering peak at 1.5 Å, which corresponds to Ru-O coordination, with an average coordination number of 3.9 (see Fig. 2b and Supplementary Table 2). This is consistent with the presence of Ru atoms dispersed in the CeO lattice, rather than forming clusters or larger particles, as no distinct Ru-Ru scattering is observed in the EXAFS spectra [Sun et al., Adv. Sci. 2022, 9, e2104749; Yan et al., Chem Catalysis 2022, 2, 1594-1623]. This atomic-scale distribution of Ru is further confirmed by High-Angle Annular Dark-Field Scanning Transmission Electron Microscopy (HAADF-STEM), which shows a uniform distribution of Ru atoms within the CeO matrix, without any observable Ru clusters (see Fig. 2d). Additionally, Energy-Dispersive X-ray Spectroscopy (EDS) mapping in conjunction with STEM images corroborates this uniform distribution, where Ru is homogeneously dispersed across the sample, confirming the single-atom nature of the catalyst [Zhang et al., Nat. Commun. 2023, 14, 2664]. Furthermore, our Rietveld refinement of X-ray diffraction (XRD) data also supports this hypothesis, where the lattice contraction in CeO upon Ru doping (Δr = 0.08 Å) suggests the incorporation of Ru into the lattice, rather than the formation of bulk RuO particles (see Supplementary Fig. 2 and Table 1). This lattice strain also contributes to the formation of oxygen vacancies, which are crucial for the catalytic activity. Based on these findings, we are confident that the Ru@CeO labeling is a more suitable and intuitive way to describe the atomic dispersion of Ru in the RC catalyst.\nRelated discussions (yellow-highlighted in Page 2, Line 28; Page 5, Line 104; and Page 25, Line 595) have been added in the revised manuscript as follows:\n“Here, we introduce an atomic trapping strategy that restructures the interfacial chemistry of Ba Sr Co Fe O perovskite and Ru@CeO fluorite heteroelectrodes, achieving catalytic synergy through manipulating Ru coordination environment.\nHere, a Ru-doped CeO catalyst (Ru@CeO , denoted as RC) as a second phase for composite electrode is synthesized through a co-precipitation method followed by calcination at 800 °C for 2 h.\nThe Ru@CeO catalyst was obtained by a co‐precipitation procedure.”\nComment 3:\nTo gain a more comprehensive understanding of the influence of atomic-trapping on crystal structure, it is useful to provide Rietveld-refined XRD patterns and lattice parameters for BSCF and BSCF-C30.\nResponse to C3:\nWe sincerely thank the reviewer for the thoughtful suggestion. We agree that providing Rietveld-refined XRD patterns and lattice parameters for BSCF and BSCF-C30 (Ru-free composite) are useful for a clearer understanding of the structural changes induced by atomic trapping. In response, we have included these details in the revised manuscript (Fig. R1).\nRegarding the crystal structure, the lattice parameters of the cubic phase in BSCF, BSCF-C30, and BSCF-RC30 show minimal variation. This is primarily due to the small ionic radius difference between the hexacoordinated Ru ions and the Co/Fe ions in the perovskite lattice, which results in only slight changes in the lattice parameters of the cubic phase. For the fluorite phase (CeO ), BSCF-RC30 shows a pronounced lattice contraction (a = 5.4059 Å) relative to both BSCF-C30 (a = 5.4102 Å) and RC (a = 5.4085 Å) (Table R1). This additional shrinkage is attributed to charge-compensating oxygen vacancies generated during Ru migration [Xing et al., ACS Energy Lett. 2019, 4, 2601-2607]. As Ru4+ is captured by the BSCF perovskite, the remaining fluorite phase loses positively charged Ru4+ ions, creating a local charge deficit. To balance this, oxygen vacancies form, which shorten the average Ce-O bond length [Zhuang et al., Nat. Catal. 2022, 5, 300-310]. This bond length reduction compacts the fluorite structure, resulting in the observed lattice contraction. The reduced Ce-O bond length and lattice size in BSCF-RC30’s fluorite phase are critical for proton conduction. Shorter Ce-O bonds decrease the distance between oxygen ions, lowering the energy barrier for proton hopping [Zou et al., J. Mater. Chem. A 2022, 10, 5381-5390]. Combined with the increased oxygen vacancy concentration (from charge compensation), this structural optimization is expected to enhance bulk proton conductivity [Chroneos et al., Energy Environ. Sci. 2011, 4, 2774]. Once again, we thank the reviewer for their valuable suggestion, and we trust that these additions improve the clarity of the manuscript.\nRelated data (Supplementary Fig. 6 and Table 1) and discussions (yellow-highlighted in Line 140-148, Page 7) have been added in the revised manuscript as follows:\n“Structural analysis further reveals minimal variation in perovskite lattice parameters across BSCF, BSCF-C30 (Ru-free BSCF and CeO composite at a 70 : 30 mass ratio), and BSCF-RC30 (Supplementary Fig. 6), attributable to the negligible ionic radius mismatch between Ru4+ and host Co/Fe cations (Δr < 0.03 Å). For the fluorite phase, BSCF-RC30 shows a pronounced lattice contraction (a = 5.4059 Å) relative to both BSCF-C30 (a = 5.4102 Å) and the RC catalyst (a = 5.4085 Å) (Supplementary Table 1). This additional shrinkage is attributed to charge-compensating oxygen vacancies generated during the trapping of Ru, which shorten the average Ce-O bond length, facilitating the lowering of the proton hopping barrier and enhancing bulk proton conduction27.”\nComment 4:\nGiven that many single-atom catalysts excel in low-temperature O activation, it is recommended to consider testing symmetric-cell or half-cell data for a pure RC electrode to justify why the composite route is superior. At the same time, a succinct explanation of the benefits of choosing BSCF as the perovskite matrix will be helpful in highlighting the synergistic effect of the RC and BSCF pairing.\nResponse to C4:\nWe sincerely thank the reviewer for this insightful comment. In response, we conducted additional symmetric-cell measurements using pure RC as the air electrode, and the resulting EIS spectra are shown in Fig. R2. The RC electrode shows a large polarization resistance of ~2500 Ω cm2 at 550 °C. This indicates that despite the atomic dispersion of Ru, the RC lacks sufficient triple conductivity, making it difficult to meet the requirements for high-performance air electrodes in R-PCECs [Zhang et al., J. Power Sources 2023, 564, 232832; Lan et al., J. Eur. Ceram. Soc. 2024, 44, 996-1004]. This highlights the necessity of forming a composite with BSCF. As discussed in the manuscript (see Fig. 2j, k and Supplementary Fig. 21, 27), BSCF offers several critical advantages: (i) high oxygen ion and electronic conductivity; (ii) abundant oxygen vacancies; and (iii) a perovskite lattice capable of accommodating and stabilizing migrated Ru atoms via the formation of Ru-O-B (B = Co, Fe) bonds. This integration promotes favorable interfacial charge transfer and boosts hydration and oxygen activation kinetics (see Fig. 3g, j). The dramatically lower polarization resistance observed for BSCF-RC30 compared to pure BSCF or RC electrodes substantiates the synergistic catalytic effects achieved through atomic trapping at the heterointerface (see Fig. 4a and Fig. R2). This evidence collectively underscores the rationality and superiority of the composite design over the single-phase electrode approach.\nRelated data (Supplementary Fig. 28) and discussions (yellow-highlighted in Line 310-318, Page 13) have been added in the revised manuscript as follows:\n“Although the single fluorite-structured RC contains atomically dispersed Ru active sites, its insufficient oxygen ion/electronic conductivity hinders charge transport, leading to a very large polarization resistance (R ) (~2500 Ω cm2 at 550 ℃) when used as a standalone air electrode, which is not sufficient to meet the performance requirements of R-PCECs (Supplementary Fig. 28). BSCF perovskite has high oxygen ion/electron conductivity, abundant oxygen vacancies, and its lattice can stabilize the migration of Ru atoms by forming Ru-O-B (B=Co/Fe) bonds. Upon compositing, BSCF’s superior conductivity compensates for RC’s charge transport limitations, while the interfacial active sites (generated by Ru trapping) synergize with BSCF’s oxygen vacancies to enhance oxygen reduction/evolution kinetics.”\nComment 5:\nFor clarity and reproducibility, it may be helpful to (a) define abbreviations such as BSCF-RC, C, BSCFR-C, and BSCFR at first mention, and (b) provide key experimental details (e.g., gas composition, heating rate, treatment temperature, etc.) for the IR and H O-TPD measurements in the Methods section.\nResponse to C5:\nWe sincerely thank the reviewer for the constructive suggestions aimed at improving the clarity and reproducibility of our manuscript. We have carefully addressed both aspects as follows:\na) We have now clearly defined the relevant abbreviations at their first occurrence in the manuscript. Specifically, “BSCF-RC” refers to the composite air electrode consisting of the perovskite phase Ba Sr Co Fe O (BSCF) and the single-atom catalyst Ru@CeO (RC); “C” is used as a simplified label for the CeO phase alone; “BSCFR” refers to the BSCF phase after Ru incorporation; and “BSCFR-C” denotes the heterointerface region reconstructed by thermally induced Ru migration from the RC to the BSCF matrix. These definitions have been integrated into the revised manuscript to ensure terminology consistency and improve reader understanding.\nb) We have also added detailed methodological information to the Methods section concerning the IR and H O-TPD characterizations. Specifically, samples were first pretreated in 10% H O-air atmosphere at 500 °C for 5 h, followed by natural cooling to room temperature. IR spectra were acquired at room temperature in static ambient air. For the H O-TPD measurements, samples were heated from room temperature to 1000 °C at a rate of 3 °C min-1 under a flowing nitrogen atmosphere, and the desorbed water species were continuously monitored using a mass spectrometer. These modifications help ensure the transparency and reproducibility of our experimental procedures.\nRelated discussions (yellow-highlighted in Page 7, Line 141-142; Page 8, Line 174-176; and Page 26, Line 629-635) have been added in the revised manuscript as follows:\n“BSCF-C30 (Ru-free BSCF and CeO composite at a 70 : 30 mass ratio)”\n“For better comprehension, the interface before and after BSCF capture of Ru from RC are labelled as BSCF-RC and BSCFR-C (BSCFR: BSCF matrix after capture of Ru, C: pure CeO phase), respectively.”\n“ToF-SIMS (ToF-SIMS 5 iontof), H O-TPD (MS, Hiden, HPR20) and IR (IS50 FTIR) were used to analyze the hydration capacity of the samples. Samples were pretreated in a 10% H O-air atmosphere at 500 °C for 5 h, followed by natural cooling to room temperature. Spectra were acquired at room temperature in static ambient air using a Fourier-transform IR spectrometer. For the H O-TPD test, samples were heated from room temperature to 1000 °C at a rate of 3 °C min-1 under flowing nitrogen (N , 50 mL min-1). Desorbed water species (H O, m/z = 18) were monitored in real time using a mass spectrometer.”\nReviewer #2 (Remarks to the Author):\nThe study presents a detailed investigation into using \"atomic-scale interface engineering\" through Ru migration to enhance the electrochemical performance of protonic ceramic electrochemical cells (R-PCECs). The approach is well-motivated, and the experimental and computational analyses are comprehensive. Also, the manuscript is generally well-written and supported by extensive characterization and theoretical modeling. While the manuscript is suitable for publication, a few points would benefit from clarification and minor revision.\nResponse:\nWe express our sincere gratitude to the reviewer for providing us with detailed and insightful comments. These comments have greatly contributed to enhancing the quality of this work. We have made every effort to address these comments in the revised version of our manuscript. Please find our responses to the comments point by point as follows.\nComment 1:\nThe role of oxygen vacancies is central to the proposed mechanism. If available, quantifying oxygen non-stoichiometry before and after Ru trapping (e.g., through TGA or iodometric titration) would strengthen the interpretation.\nResponse to C1:\nWe sincerely appreciate the reviewer’s insightful suggestion regarding the quantitative evaluation of oxygen non-stoichiometry. We fully acknowledge that oxygen vacancies play a central role in our proposed mechanism of Ru trapping and interfacial reconstruction, and their direct quantification would indeed reinforce our interpretation.\nIn the revised manuscript, we have included both thermogravimetric analysis (TGA), iodometric titration and X-ray photoelectron spectroscopy (XPS) to evaluate the changes in oxygen non-stoichiometry before and after Ru trapping. As shown in Fig. R3a, the TGA curves reveal that the mechanically mixed sample without Ru incorporation (BSCF-C30) exhibits a total weight loss of ~1.40% between room temperature and 1000 °C, while the Ru-trapped sample (BSCF-RC30) shows a higher loss of ~1.65% under the same conditions. Considering that the RC support itself is thermally stable with negligible mass loss, this additional oxygen release in BSCF-RC30 can be attributed to the formation of oxygen vacancies. Complementary iodometric titration results show that the average oxygen non-stoichiometry (δ) increases from 0.4052 in untreated BSCF-RC30 to 0.5186 in treated BSCF-RC30, representing a ~28% increase (Fig. 3Rb). In addition, the XPS O 1s spectra of BSCF-RC30 were measured at room temperature. The results show that compared with untreated BSCF-RC30, the active oxygen (O 2-/O-) species in post-migration BSCF-RC30 increased from 32% to 41% (see Supplementary Fig. 22). These results provide strong evidence that Ru trapping induces extra oxygen vacancy formation [Zhao et al., Adv. Mater. 2024, 36, 2405052]. Furthermore, based on the XANES data in the main text (see Fig. 3a-c), it is confirmed that the oxidation states of the transition metals in BSCF remain largely unchanged during the process, indicating that the oxygen vacancy generation is not due to bulk redox process but rather interfacial charge redistribution and local structural relaxation. This further highlights the critical role of the heterointerface in modulating the defect chemistry.\nThese new data and discussions have been incorporated into the revised manuscript, with full experimental details and raw data provided in the Supplementary Information. We are grateful for the reviewer’s comment, which has significantly strengthened the mechanistic interpretation of oxygen vacancy-mediated Ru migration and interface reconstruction.\nRelated data (Supplementary Fig. 19 and 21) and discussions (yellow-highlighted in Line 255-267, Page 11-12) have been added in the revised manuscript as follows:\n“The oxygen vacancy reservoir directly governs the catalyst’s oxygen handling capacity, so we quantified oxygen vacancy differences using thermogravimetric analysis (TGA) and iodine titration (Supplementary Fig. 19). TGA curves acquired in air (50-1000 °C) show that BSCF-RC30 (Ru-trapped composite) exhibits a higher total weight loss (~1.65%) compared to Ru-free BSCF-C30 (~1.40%). Since the RC phase itself is thermally stable (negligible mass loss), this additional weight loss in BSCF-RC30 directly reflects enhanced oxygen vacancy formation induced by Ru migration. In addition, O temperature-programmed desorption (O -TPD) also verifies that BSCF-RC30 hybrid has a superior oxygen dynamic exchange capacity than the BSCF-C30 composite (Supplementary Fig. 20). Iodometric titration shows that the average oxygen non-stoichiometry (δ) of BSCF-RC30 increases by ~28% before and after thermal treatment (Supplementary Fig. 21). This increase is further confirmed by deconvoluted XPS O 1s spectra (Supplementary Fig. 22)41. The results show that the treated BSCF-RC30 exhibits an increase in the ratio of active oxygen species (O 2-/O-) from 32% to 41% compared with the untreated BSCF-RC30, which is consistent with more oxygen vacancies at the heterointerface.”\nComment 2:\nA brief discussion on the long-term stability of the trapped Ru species under operational conditions would be helpful. In particular, clarifying whether any redistribution, segregation, or structural changes of Ru were observed after extended electrochemical testing would be helpful.\nResponse to C2:\nWe thank the reviewer for the insightful question regarding the long-term stability of the trapped Ru species under operational conditions. To evaluate whether any redistribution, segregation, or changes in electronic structure occur during extended electrochemical operation, we performed comprehensive structural and spectroscopic analyses after prolonged treatment. First, XRD analysis was conducted on the BSCF-RC30 composite after thermal aging at 600 °C for 200 h (Fig R4a). The results confirmed that BSCF-RC30 still maintained the cubic perovskite phase and the fluorite phase after long-term heat treatment, with no detectable formation of secondary phases such as RuO , BaO, or SrO. This indicates that the overall crystal structure remains stable over extended operation. In order to further examine the electronic structure and distribution of Ru, we performed XPS and TEM-EDS analyses on the treated samples. The results showed that the binding energy of Ru 3p did not change significantly before and after long-term heat treatment, which indicated that the oxidation state and local electronic environment of Ru remained unchanged, with no evidence of electronic remodeling or redox instability (Fig R4b). In addition, we carried out EDS mapping on the post-treated sample. The data showed that Ru remained predominantly located at the heterointerface and within the BSCF bulk phase, while its presence in the CeO phase was negligible (Fig R4c). This suggests that Ru did not exhibit significant redistribution or reverse migration during high-temperature aging, and remains effectively anchored within the perovskite lattice, supporting the structural robustness of the atomic trapping strategy.\nIn summary, the consistent phase structure, spatial Ru localization, and stable electronic state after 200 hours of operation collectively demonstrate the excellent thermal and electrochemical stability of the trapped Ru species. These findings have been incorporated into the revised manuscript to address the reviewer’s comment.\nRelated data (Supplementary Fig.18, 43, and 44) and discussions (yellow-highlighted in Line 427-437, Page 18) have been added in the revised manuscript as follows:\n“This improved performance is primarily attributed to the BSCF-RC30 electrode’s stable phase structure, lower thermal expansion coefficient (TEC), and sustained catalytic activity (Supplementary Fig. 42). XRD patterns show that the BSCF-RC30 hybrid remains pure cubic and fluorite after sintering at 600 °C for 200 h, without the formation of RuO or alkaline earth metal second phases (Supplementary Fig. 43). Complementarily, XPS shows no shift in the Ru 3p binding energy, indicating an unchanged oxidation state, while TEM-EDS mapping reveals Ru confined to the BSCF lattice/heterointerface with negligible signal in the CeO domain (Supplementary Fig. 44, 18). The unchanged crystal structure, electronic state and spatial distribution collectively demonstrate that Ru remains securely anchored, with no detectable segregation or reverse migration during extended operation, underscoring the long-term robustness of the atomic-trapping strategy.”\nComment 3:\nThe hydration performance is well-discussed. However, it would add value to mention whether any surface species (e.g., hydroxyls, carbonates) form under operating conditions and how they might influence electrode behavior, potentially supported by relevant literature.\nResponse to C3:\nWe sincerely thank the reviewer for encouraging a deeper discussion of surface species that may form under operating conditions and their impact on electrode behaviour. To address this point, BSCF and BSCF-RC30 powders were treated in 10% H O-air at 600 °C for 24 h, followed by SEM (Fig. R5) and XRD (Fig. R6) analyses. SEM images reveal that pure BSCF has a small amount of flaky carbonate deposition after air calcination; after water treatment, the surface is covered with dense hydroxide nanoparticles [Park et al., Energy Environ. Sci. 2024, 17, 1175-1188; Fabbri et al., Adv. Mater. 2012, 24, 195-208; Ren et al., J. Mater. Chem. A, 2019, 7, 18365-18372]. The formation of carbonates and hydroxides is a common degradation mechanism for Ba-based perovskite air electrodes under H O/CO -containing atmospheres, and their detrimental effects on electrode behavior have been widely reported [Liu et al., ACS Appl. Mater. Interfaces, 2020, 12, 50433-50445; Zheng et al., ACS Energy Lett. 2024, 9, 3962-3969]. For example, carbonate deposition covers active sites on the electrode surface, shortens the triple-phase boundary (TPB) length for the oxygen reduction reaction (ORR), and increases oxygen diffusion resistance [Luo et al., ACS Energy Lett. 2022, 7, 2970-2978; Chen et al., ACS Catal. 2019, 9, 7137-7142]. The generation of hydroxide nanoparticles will limit the adsorption and desorption of water, affect the hydration reaction, and block the proton transfer path [Fabbri et al., Adv. Mater. 2012, 24, 195-208].\nIn contrast, no obvious carbonates and hydroxides were observed in surface of BSCF-RC30 due to the uniform encapsulation of high-surface-area CeO on the surface and the formation of a strongly interacting heterointerface. The core mechanism of this difference is that the heterogeneous interface constructed by the Ru atom capture strategy effectively inhibits the migration of alkaline earth metals in the BSCF bulk phase and avoids contact with H O/CO . Meanwhile, the excellent hydration ability of CeO further reduces the probability of hydroxide formation [Xing et al., ACS Energy Lett. 2019, 4, 2601-2607]. The stable phase structure under high water partial pressure and the unique surface configuration enable the BSCF-RC30 cell to maintain good stability in the electrolysis mode under 10%/50% H O-air (see Fig. 6i).\nNotably, no distinct diffraction peaks of BaCO or Ba(OH) were detected in XRD patterns of either sample (Fig. R5), which is in agreement with related studies [Chen et al., Adv. Funct. Mater., 2018, 28, 1704907; Liu et al., Adv. Funct. Mater., 2023, 34, 2311140]. This is mainly because the surface deposition exists in the form of nanoscale amorphous or low-crystallinity phases, which is below the detection limit of conventional XRD [Wang et al., ACS Appl. Mater. Interfaces 2025, 17, 39053−39063].\nRelated data (Supplementary Fig. 50) and discussions (yellow-highlighted in Line 557-565, Page 24) have been added in the revised manuscript as follows:\n“To clarify why the BSCF-RC30 electrode withstands high steam pressures so well, we examined its surface after 24 h in 10% H O-air at 600 °C. BSCF-RC30 retains a clean morphology, showing no detectable carbonate or hydroxide deposits, which is attributed to the strong interaction of the heterogeneous interface that inhibits the outward migration of alkaline earth metals from the BSCF bulk phase and the high-surface-area CeO wrapped on the surface to avoid contact with H O/CO (Supplementary Fig. 50). In contrast, pristine BSCF forms flaky carbonates after air calcination and densely distributed hydroxide nanoparticles after prolonged wet treatment, which can shorten the triple-phase boundary and slow down proton transfer58.”\nComment 4:\nIncluding standard deviations or a brief discussion on the reproducibility of key electrochemical measurements (e.g., ASR, PPD) across multiple samples would help evaluate data reliability.\nResponse to C4:\nThank you for your thorough evaluation of our manuscript and the insightful suggestion regarding the reproducibility of key electrochemical measurements. We sincerely appreciate your valuable feedback, which has greatly enhanced the rigor and credibility of our study. In response to your concern, we performed parallel tests on BSCF-RC30 symmetric and fuel cell (Cell#2), while the data in the initial manuscript was labeled Cell#1, and systematically verified the reproducibility of the area specific resistance (ASR) and the peak power density (PPD).\nThe supplementary data demonstrate excellent consistency in ASR values across temperatures. Under 5% H O-air conditions, the ASR deviations of Cell#1 and Cell#2 are less than 6% in the range of 500-700 °C (e.g., at 600 °C, Cell#1 is 0.085 Ω cm2, Cell#2 is 0.081 Ω cm2, and the relative standard deviation (RSD) ~5%) (Fig. R7a, b), further confirming the stability of the atomic capture strategy in regulating the interface structure. For single-cell PPD, the difference between Cell#1 (1.511 W cm-2) and Cell#2 (1.478 W cm-2) at 650 °C is only 2.2%, and their PPD values at 550 °C (0.816 vs. 0.813 W cm-2) are almost identical, with RSDs all below 3% (Fig. R7c, d). These results indicate excellent repeatability in cell assembly and multi-interface synergy. All raw data (including impedance spectra and IV curves) have been uploaded to the supplementary materials as per the journal’s requirements for verification.\nIn summary, the parallel sample tests fully validate the repeatability of key electrochemical parameters, and the public availability of raw data further ensures the authenticity of our results. We deeply thank you for highlighting the importance of data reliability; your feedback has significantly strengthened the rigor of this work.\nRelated data (Supplementary Fig. 31 and 47) and discussions (yellow-highlighted in Page 14, Line 329-333 and Page 19, Line 449-452) have been modified in the revised manuscript as follows:\n“For completeness, we repeated the symmetric-cell test with an independently fabricated specimen (Cell #2). As shown in Supplementary Fig. 31, the resulting ASR values deviate by no more than ±6% from the previous ones over 500-700 °C, confirming that the low-resistance trend is highly reproducible and intrinsic to the atomic-trapping interface.”\n“The single-cell performance was likewise reproduced on a second device prepared under identical conditions. Peak power densities differ by only 2.2% at 650 °C and < 3% at 550 °C (Supplementary Fig. 47), underscoring the robustness and reliability of the composite-electrode design.”\nComment 5:\nIn Figure 3h, including an iso-value label and plane orientation in the differential charge density plot would make it clearer.\nResponse to C5:\nThanks for pointing out this. The isosurface value of 0.005 is added into the caption of Fig. 3h (Fig. R8). We also add the orthogonal cartesian axes with clearly labelled the Miller index of crystal plane to each Figures containing the interfacial structure, including Figure 2j, 3h, S11, and S12. By doing these, we sincerely hope that your concern can be addressed in the revised version.\nRelated data (yellow-highlighted in Line 226-228, Page 10) have been modified in the revised manuscript as follows:\nFig. R8 (Fig. 3h) | The differential charge density plot of BSCFR-C heterogeneous interface. Electron accumulation and depletion regions are highlighted in yellow and blue, with an isosurface level of 0.005.\nComment 6:\nFor CI-NEB calculations, such as Figures 2j, 3k, and Supplementary Figure S12b, the x-axis label “transport index” appears unclear. It would be more appropriate to use “reaction coordinate” or “image number.” Furthermore, for clarity, please explicitly indicate the reactant, transition state, and product on each plot and ensure a consistent number of intermediate images across cases. The curve smoothing in Figure S11b may introduce artifacts.\nResponse to C6:\nThanks for your suggestions. We now use x-axis labels of “reaction coordinate” in all CINEB determined energy profiles, including Fig. 2j, 3k, and S12 (Fig. R9a, d-f). We also highlight the reactant, transition state, and product using blue circles with solid, dotted, and dash lines in the Ru transport pathways across the interface, including Fig. 2j, S11, and S12 (Fig. R9a-c).\nRegarding the intermediate image number, a common and efficient strategy is to set it as 0.8 to 1 times of root squared sum distance between the initial and final configuration, namely, the reactant and product. As different transport pathways may share totally different transport distance and sum distance, we decide to insert 1-3 intermediate images for proton rotation and hopping, and more intermediate images for Ru transport across the interface (Fig. R9f).\nFor your concerns regarding Fig. S12b, we apologize that we need to use more intermediate images than those in other figures, due to the long transport distance and high root squared sum distance. In the revised version, you can also see that it is a complicate transport pathway with two transition states. Given you suggests, we change the curve smoothing style and recalculate the transport energy barrier only with the data points, namely, without any data from the smoothing curve (Fig. R9e). We hope that the artifacts can be largely avoided in the revised version.\nRelated data (Fig. 2j and 3k, Supplementary Fig. 11 and 12) and discussions (yellow-highlighted in Page 6, Line 130-132; Page 10, Line 231; and Page 8, Line 181-186) have been modified in the revised manuscript as follows:\n“To validate the critical role of oxygen vacancy concentration in atomic trapping dynamics, we systematically investigate a counterexample of La NiO (LN)-RC interface33,34 with almost no oxygen vacancy, and reveals drastically different migration kinetics: Ru atom face a prohibitive energy barrier of 2.00 eV to migrate from RC side to LN side (Supplementary Fig. 12), higher than that in oxygen-deficient BSCFR-C.”\nReviewer #3 (Remarks to the Author):\nThis manuscript from Liu et al presents a compelling strategy for enhancing the performance of air electrodes for reversible protonic ceramic electrochemical cells (R-PCECs). The authors describe a phenomenon termed “atomic trapping,” where Ru species, initially doped into a CeO2 support, are shown to migrate into the lattice of a BSCF perovskite during thermal processing. This process is reported to create a highly active and stable heterointerface, leading to remarkable bifunctional performance in both fuel cell and electrolysis modes. The work combines detailed electrochemical characterization with advanced microscopy, spectroscopy, and DFT calculations to elucidate the underlying mechanisms, which are attributed to optimized charge redistribution, interfacial oxygen vacancy formation, and enhanced triple-phase boundary kinetics.\nThe work is of high quality and the results are impressive however, there are several key points that should be addressed to strengthen the manuscripts conclusions before it can be considered for the publication.\nResponse:\nWe extend our sincere appreciation to the reviewer for positive feedback and recognition of the value of our research. We are also deeply grateful for the numerous insightful and critically important comments raised by the reviewer, as they have significantly contributed to improving the quality and rigor of our manuscript. We have taken every comment into careful consideration and made revisions to our manuscript accordingly. To address all the comments in a comprehensive manner, we have prepared point-to-point responses, which will be provided below.\nComment 1:\nThe sentence starting from line 34 in the abstract section is lengthy and missing key vital points. For instance, the authors did not indicate the temperature at which the peak power density in the FC mode and current density in the EC mode were achieved. It is recommended that the sentence be simplified and clarified to avoid misleading readers into assuming it was obtained at 600 ℃.\nResponse to C1:\nWe sincerely thank the reviewer for pointing out the overly long sentence and the missing temperature information in the abstract. We realized that explicitly stating the test temperatures is essential for an accurate presentation of the key performance metrics. Following the recommendation, we have reorganized the sentence and added the specific temperatures for the fuel cell mode (650 °C), electrolysis mode (650 °C), and durability (600 °C) measurements. The revised sentence now reads:“The electrode with low Ru loading exhibits exceptional bifunctionality, delivering a peak power density of 1.51 W cm-2 and an electrolysis current density of -2.21 A cm-2 at 650 °C. It demonstrates remarkable durability with minimal degradation (0.09 mV h-1) over 400 h at 600 °C, providing a universal strategy for next-generation solid-state energy devices.”\nWe believe this revision eliminates any potential confusion regarding the temperature conditions and improves the clarity and completeness of the abstract. We appreciate the reviewer’s constructive comment, which has helped us refine the manuscript.\nRelated discussions (yellow-highlighted in Line 33-36, Page 2) have been added in the revised manuscript as follows:\n“The electrode with low Ru loading exhibits exceptional bifunctionality, delivering a peak power density of 1.51 W cm-2 and an electrolysis current density of -2.21 A cm-2 at 650 °C. It demonstrates remarkable durability with minimal degradation (0.09 mV h-1) over 400 h at 600 °C, providing a universal strategy for next-generation solid-state energy devices.”\nComment 2:\nThe central claim revolves around “atomic trapping”. While the evidence for Ru migration is strong, the term “trapping” implies a specific dynamic mechanism that may not be fully supported by the data, which could also be interpreted as standard thermodynamically driven diffusion. The authors must provide a more robust justification for using this specific term over a simpler description or provide more direct evidences of a “trapping” phenomenon.\nResponse to C2:\nThank you for your thoughtful and constructive comment regarding the use of the term “atomic trapping”. We understand your concern that this term may imply a specific dynamic mechanism, while the current data might also be interpreted under conventional thermally driven diffusion. In response, we would like to clarify the distinction between the two concepts and explain, based on the data and literature, why the use of “atomic trapping” in our work is appropriate and well-supported.\nFundamentally, “atomic trapping” differs from standard thermal diffusion in both directionality and reversibility. While thermal diffusion involves random, reversible atomic migration under chemical potential gradients, our use of “atomic trapping” refers to a directional and irreversible transfer of Ru from CeO to the BSCF phase, where it becomes structurally anchored and no longer exhibits observable reverse migration [Jones et al., Science 2016, 353, 150-154; Li et al., Angew. Chem. Int. Ed. 2023, 62, e202309341; Zhang et al., Nat. Commun. 2024, 15, 1234]. This process is mediated by the presence of thermodynamically favorable and kinetically stabilizing sites such as oxygen vacancies and strong covalent bonding sites, forming an asymmetric “migration-then-lock” behavior rather than bidirectional equilibrium diffusion [Zhuang et al., Nat. Catal. 2022, 5, 300-310].\nOur experimental and theoretical results support this interpretation. The post-annealing TEM-EDS mapping (Fig. R10) clearly show that Ru remains concentrated in the BSCF side even after long-term calcination, with no detectable back-diffusion into the CeO region. This spatially confined profile, persistent under high temperature, suggests a trapping behavior inconsistent with purely reversible diffusion. Additionally, EXAFS analysis (see Fig. 2h, i) reveals that Ru undergoes a coordination transformation into stable BO octahedra, indicating that it is incorporated into the lattice framework rather than loosely residing in a mobile state. The CI-NEB calculations (see Fig. 2j, k) further confirm that Ru experiences a significant migration barrier and binds strongly at BSCF sites, making spontaneous back-diffusion energetically unfavorable. EELS and XANES data support this with evidence of interfacial charge redistribution and oxygen vacancy stabilization, reinforcing the chemical and structural locking of Ru within the BSCF matrix (see Fig. 3c-f).\nTo further understand our terminology, we note that similar concepts have been reported in recent literature. The term “reverse atomic capture” has been employed to describe analogous behaviors of element sequestration at engineered interfaces [Zhao et al., Adv. Mater. 2024, 36, 2405052; Zhuang et al., Nat. Catal. 2022, 5, 300-310], and “trapping of exsolved atoms” has also been discussed in the context of stable anchoring of mobile species [Song et al., Angew. Chem. Int. Ed. 2023, 63, e202313361]. These works confirm that such directional anchoring behavior is well-established and conceptually aligned with our system, further validating our use of “atomic trapping.”\nIn response to your helpful suggestions, we have made the following revisions: we now emphasize the spatial retention of Ru in the BSCF region in the Results section, clarify the distinction between “trapping” and “diffusion”, and reference relevant literature to justify our terminology. Additionally, we have added Fig. R10 in the Supplementary Information to provide supporting EDS scan data that visually confirm the Ru localization after long-term calcination. We believe these additions will strengthen the precision and credibility of our manuscript.\nThank you again for your critical insights, which have helped us improve the clarity and rigor of our work.\nRelated data (Supplementary Fig. 18) and discussions (yellow-highlighted in Line 249-254, Page 11) have been added in the revised manuscript as follows:\n“Furthermore, TEM-EDS mapping of BSCF-RC30 collected after calcination at 600 °C for 200 h still detected Ru only in the BSCF domains (Supplementary Fig. 18). These results verify that atomic trapping arises from the combined thermodynamic driving force for Ru migration and the presence of robust anchoring sites (oxygen vacancies and strong Ru-O-Co/Fe covalent bonds) which lock Ru irreversibly after migration, rather than permitting the random, reversible diffusion expected from a purely chemical-potential-driven process40.”\nComment 3:\nThe link between Ru migration and presence of oxygen vacancies in the perovskite is critical point. The LN-RC counterexample is illustrative, but the argument would be more convincing if the role of oxygen vacancies as the primary driving force was more definitively established, beyond just correlation.\nResponse to C3:\nWe sincerely thank the reviewer for the insightful comment regarding the causal role of oxygen vacancies in driving Ru migration. Indeed, establishing a clear transition from correlation to causation is essential for reinforcing the mechanistic rigor of our study. In the revised manuscript, we have systematically integrated theoretical modeling and experimental validation to more definitively demonstrate the critical role of oxygen vacancies in facilitating Ru migration and trapping.\nFrom the theoretical perspective, we constructed BSCF-RC and LN-RC interface models, with significantly different oxygen vacancy concentrations. CI-NEB calculations revealed that Ru migration in the oxygen-deficient BSCF interface exhibits lower kinetic barriers and higher binding energies at the target sites; while in LN, due to the complete oxygen octahedron and lack of vacancies, the migration barrier of Ru is significantly increased and difficult to stabilize (see Fig. 2j and Supplementary Fig. 12). Moreover, structural optimization of Ru insertion into BSCF with and without oxygen vacancies further confirmed that the presence of vacancies dramatically improves the incorporation of Ru into the BO octahedral lattice, enabling stable coordination and anchoring (see Fig. 2k). These results clearly establish the direct thermodynamic and kinetic relevance of oxygen vacancies to the trapping process.\nThis theoretical trend is consistently reflected in our experimental observations. STEM-EDS analyses in the BSCF-RC30 and LN-RC30 systems showed clear Ru enrichment toward the BSCF side, with no discernible accumulation in the LN case. This spatial contrast further substantiates the critical role of oxygen vacancies in promoting directional Ru migration and trapping. Additionally, we performed EDS mapping on a newly prepared BaCo Fe Zr Y O (BCFZY)-RC30 composite (See response to Comment 5). We observed that Ru was again captured at the BCFZY side of the interface. Given that BCFZY has been independently reported to exhibit a high concentration of oxygen vacancies at elevated temperatures [Duan et al., Science 2015, 349, 1321-1326], this result supports the generality and robustness of the oxygen-vacancy-assisted Ru trapping mechanism across different perovskite systems.\nBased on these findings, we also revisited the microscopic mechanisms by which oxygen vacancies facilitate Ru trapping. Locally distorted octahedra created by oxygen deficiency offer accessible structural space, while the surrounding electrostatic and electronic environments are favorable for Ru incorporation and charge compensation [Lee et al., ACS Catal. 2019, 9, 7099-7108; Sun et al., J. Am. Chem. Soc. 2022, 144, 1174-1186; Song et al., Angew. Chem. Int. Ed. 2023, 63, e202313361]. This dual functionality of structural regulation and electronic stabilization constitutes the underlying basis for irreversible Ru trapping in perovskite lattices.\nIn summary, we not only observed preferential Ru accumulation near oxygen-vacant regions at the macroscopic level but also elucidated, through first-principles calculations, the dual role of oxygen vacancies as both migration pathways and anchoring sites. In the revised manuscript, we have clarified these mechanistic insights and provided more context regarding the rationale behind our model systems and comparative analyses, in order to more directly address the reviewer’s concern. We believe these additions substantially enhance the clarity and strength of our conclusions.\nRelated discussions (yellow-highlighted in Page 8, Line 181-193 and Page 16, Line 337-382) have been added in the revised manuscript as follows:\n“To validate the critical role of oxygen vacancy concentration in atomic trapping dynamics, we systematically investigate a counterexample of La NiO (LN)-RC interface33,34 with almost no oxygen vacancy, and reveals drastically different migration kinetics: Ru atom face a prohibitive energy barrier of 2.00 eV to migrate from RC side to LN side (Supplementary Fig. 12), higher than that in oxygen-deficient BSCFR-C. This is consistent with STEM-EDS mapping observation of Ru retainment in the CeO matrix after the same heat treatment (Supplementary Fig. 13). Locally distorted octahedra created by oxygen deficiency offer accessible structural space, while the surrounding electrostatic and electronic environments are favorable for Ru incorporation and charge compensation30. Parallel DFT/CI-NEB simulations and STEM-EDS mapping of the oxygen-vacancy-rich BSCF-RC interface versus the vacancy-poor LN-RC analogue demonstrate that lattice vacancies provide both the low-barrier migration channel and the deep B-site trap needed for irreversible Ru capture. Crucially, this thermodynamically driven atomic trapping fundamentally differs from reverse-trapping phenomena reported elsewhere. Unlike acid-base extraction of alkaline-earth ions (e.g., Sr2+ removal by MoO ) which generates surface vacancies and secondary phase, the present strategy harnesses bulk migration of Ru from the fluorite into the perovskite lattice29. Oxygen-vacancy pathways and thermodynamic stabilization at B-sites drive heterointerface reconstruction, thereby enhancing triple conductivity and interfacial activity, distinct from the surface effects in conventional reverse trapping techniques.”\n“Indeed, elemental mapping of a newly prepared BCFZY-RC30 composite (Supplementary Fig. 33) likewise shows Ru exclusively anchored on the BCFZY side, confirming that oxygen-vacancy-assisted Ru trapping is both universal and robust across diverse perovskite hosts.”\nComment 4:\n“Theses” in line 189 is a typo.\nResponse to C4:\nWe thank the reviewer for the careful reading of the manuscript. As correctly pointed out, “Theses” in line 189 is a typographical error and should be “these.” We have corrected this in the revised manuscript and carefully rechecked the text to eliminate similar typos. We appreciate your attention to detail.\nRelated discussions (yellow-highlighted in Line 205, Page 9) have been added in the revised manuscript as follows:\n“These collective evidences delineate an activation mechanism at the BSCF-RC30 heterointerface, mediated by localized electron redistribution, vacancy-coupled oxygen dynamics, and proton transfer modulation.”\nComment 5:\nThe authors claim this strategy is a “universal paradigm” for Co-based perovskites rich in oxygen vacancies. To support this, they demonstrate enhanced performance by adding RC to PBC, BCFZY, SCFN, and LSCF electrodes (Figure 4c). However, the manuscript lacks direct evidence that the same Ru migration mechanism occurs in these other systems. Without supporting STEM-EDS or spectroscopic data for these other composites, the assertion that performance enhancement stems from the same “atom trapping” mecanism remains assumption. To validate such a strong claim of universality, direct evidence of Ru migration should be provided for at least one other perovskite system.\nResponse to C5:\nWe sincerely thank the reviewer for the insightful comment regarding the universality of the proposed “atomic trapping” mechanism. As you correctly pointed out, although Fig. 4c demonstrates performance enhancement in several Co-based perovskite systems (e.g., PBC, BCFZY, SCFN, and LSCF) upon RC addition, the original manuscript lacked direct evidence that these improvements are driven by the same Ru migration and trapping mechanism. To address this, we have now included STEM-EDS characterization of the BCFZY-RC30 composite electrode after thermal treatment in the revised manuscript (Fig. R11). The results clearly show significant Ru accumulation on the BCFZY side, indicating successful migration and anchoring of Ru from RC into the BCFZY phase. This observation closely mirrors the behavior seen in the BSCF-RC30 system, thereby providing direct experimental support for the same trapping mechanism in a second distinct perovskite host. While a full spectroscopic or microscopic investigation of all listed systems is beyond the scope of the current study, this additional direct evidence in the BCFZY system offers strong support for the general applicability of the proposed strategy. In the revised manuscript, we have explicitly presented and discussed these new data. We appreciate the reviewer’s valuable suggestion, which has helped us strengthen the rigor and clarity of our conclusions.\nRelated data (Supplementary Fig. 33) and discussions (yellow-highlighted in Line 379-382, Page 16) have been added in the revised manuscript as follows:\n“Indeed, elemental mapping of a newly prepared BCFZY-RC30 composite (Supplementary Fig. 33) likewise shows Ru exclusively anchored on the BCFZY side, confirming that oxygen-vacancy-assisted Ru trapping is both universal and robust across diverse perovskite hosts.”\nComment 6:\nIs the total Ru loading in the BSCF-RC30 electrode comparable to the Ru content in the BSCFR0.02 electrode?\nResponse to C6:\nWe appreciate the reviewer’s question regarding the comparability of Ru loading between the BSCF-RC30 and BSCFR0.02 electrodes. To ensure meaningful comparison and minimize compositional bias, we deliberately designed the Ru doping level in BSCFR0.02 (2 at.%) to closely match the estimated amount of Ru incorporated into the perovskite structure in the BSCF-RC30 composite. The total Ru content in BSCF-RC30 is ≈0.89 wt.%. This closely matches the calculated Ru content of ≈0.92 wt.% in BSCFR0.02, confirming that the two samples have nearly identical Ru loadings by weight. This approach was intended to isolate the influence of the trapping mechanism from simple Ru concentration effects.\nComment 7:\nWhy the in-situ migration of Ru from CeO creates a more active interface than simply pre-doping Ru into BSCF lattice? Is it the formation of a specific interfacial structure that cannot be achieved by other means? Elaborating on this point would significantly improve impact of this great work.\nResponse to C7:\nWe sincerely thank the reviewer for raising this important point. Compared to conventional pre-doping strategies, the in-situ migration of Ru from CeO enables the formation of a functionally synergistic heterointerface, which cannot be achieved through direct lattice doping alone. This unique interface structure is central to the superior catalytic performance of the BSCF-RC30 system.\nThe heterointerface formed via Ru migration exhibits strong atomic-level coupling and localized electronic reconstruction. TEM-EDS mapping reveals that Ru is predominantly anchored at the BSCF side of the interface (see Fig. 2g), and EXAFS confirms its incorporation into stable RuO octahedra (see Fig. 2h, i). Complementary EELS and XANES results further show pronounced Co valence increase and electronic restructuring at the interface, indicating localized charge redistribution and enhanced redox activity (see Fig. 3d-f). This atomic-scale local regulation effect cannot be obtained by uniform doping.\nMoreover, the in-situ constructed CeO modified layer contributes significantly to interfacial hydration and proton conduction. CeO is known for its excellent proton transport properties and hydration capacity (see Fig. 3g, j and Supplementary Fig. 24, 25), and its intimate contact with BSCF at the heterointerface facilitates the formation of an extended triple-phase boundary under humid conditions [Wang et al., J. Am. Chem. Soc. 2016, 138, 6298-6305; Xing et al., ACS Energy Lett. 2019, 4, 2601-2607]. This leads to markedly lower interfacial polarization resistance and higher current density in the BSCF-RC30 electrode, while such benefits are absent in the BSCFR0.02 system with uniformly distributed Ru. Additionally, the high-temperature migration process promotes adaptive structural integration, making the formed heterogeneous interface have good structural and thermodynamic stability (see Fig. 4h, i and Supplementary Fig. 7). The CeO shell on BSCF acts as a protective layer, mitigating surface degradation, element leaching, and microstructural coarsening under hydrothermal conditions, thereby improving the long-term durability of the electrode (see Fig. 5e, g, i). In contrast, pre-doped Ru remains embedded in the bulk lattice without contributing to interfacial stability or surface robustness.\nIn summary, the Ru migration strategy synergistically improves the oxygen activation performance of the fluorite phase and the perovskite phase, and also has irreplaceable advantages in constructing a highly active, strongly coupled and stable multifunctional interface. This mechanism not only enhances the oxygen vacancy-related reaction kinetics, but also expands the design ideas of perovskite electrode materials. These points have been further elaborated and emphasized in the revised manuscript to strengthen the mechanistic depth and overall impact of our work.\nRelated discussions (yellow-highlighted in Line 365-371, Page 16) have been added in the revised manuscript as follows:\n“The BSCF-RC30 electrode exhibits an ASR of only 0.57 Ω cm2, markedly lower than BSCF-C30 (1.22 Ω cm2), BSCFR0.02 (1.31 Ω cm2) and the combined BSCFR0.02-C30 (1.19 Ω cm2). This pronounced advantage arises from Ru trapping, which establishes a highly active CeO /BSCF heterointerface, while the efficient charge transfer and abundant vacancies at the interface not only provide active sites for hydration and oxygen activation, but also promote ion exchange between the two phases, which is an interfacial synergistic effect that neither simple CeO modification nor Ru pre-doping can provide.”\nComment 8:\nThere appear to be a contradiction in the E a trend reported. For the BSCF-RCx series, the BSCF-RC30 electrode shows the lowest E (0.94 eV) in dry air (lines 334-337). However, for the other perovskite systems (PBC, BCFZY, etc.), the text state that the hybrid electrodes exhibit higher activation energies in dry air compared to their pristine counterparts (lines 404-406 and Supp. Fig. 38). The explanation provided, that limited oxygen vacancy mobility in the RC phase increases E , does not account for why the BSCF-RC30 system behaves differently. This inconsistency should be clarified. Is there some unique interaction at the BSCF-RC interface that is not present in the others?\nResponse to C8:\nWe sincerely thank the reviewer for the thoughtful comment regarding the reported trends in activation energy (E ). We would like to clarify the apparent discrepancy. The E value of 0.94 eV for the BSCF-RC30 electrode, as mentioned in lines 335-337, corresponds to the measurement under humidified 5% H O-air conditions, as shown in Fig. 4a. In contrast, the E values under dry air conditions for the BSCF-RC series are presented in Supplementary Fig. 29. Under these dry conditions, BSCF and BSCF-RC30 exhibit activation energies of 1.03 eV and 1.14 eV, respectively, which is consistent with the trend observed in other perovskite systems (e.g., PBC, BCFZY) where RC addition slightly increases E in dry air due to reduced oxygen vacancy mobility in the RC phase. The confusion likely arose from interpreting the humidified-condition data in Fig. 4a as dry-air values. To prevent this misunderstanding, we have revised the manuscript to clearly indicate the specific testing atmosphere associated with each activation energy measurement in the corresponding figure caption and main text.\nWe appreciate the reviewer’s careful attention to this point, which has allowed us to improve the clarity and accuracy of our presentation.\nRelated discussions (yellow-highlighted in Page 14, Line 333-335 and Page 15, Line 349-350) have been added in the revised manuscript as follows:\n“a Arrhenius plots of the ASRs of the BSCF and BSCF-RCx (x = 20, 30 and 40 wt.%) air electrodes under humidified 5% H O-air conditions.”\n“Activation energy analysis derived from ASR-temperature dependence (500-700 °C) reveals values of 1.04, 1.02, 0.94, and 0.96 eV for bare BSCF, BSCF-RC20, BSCF-RC30, and BSCF-RC40 under 5% H O-2 conditions, respectively (Fig. 4a).”\nComment 9:\nFor Fig 5., please fix the error with the captioning of the a, b, and c parts.\nResponse to C9:\nWe thank the reviewer for the careful observation regarding the caption of Fig. 5. Indeed, the caption in the original manuscript was not sufficiently accurate in describing the content of Fig. 5a, b, and c, which may cause confusion for readers. In the revised manuscript, we have corrected and clarified the caption as follows:\n“Fig. 5 | Performance and stability of BSCF-RCx cells in PCFC mode. a Current-voltage polarization and power density curves at different temperatures for the BSCF-RC30 air electrode. b Corresponding electrochemical impedance spectra at open circuit voltage. c Peak power density comparison for BSCF-RCx series cells operated at 650 and 500 °C.”\nWe have also reviewed other figure captions throughout the manuscript to ensure consistency and accuracy. We appreciate the reviewer’s suggestion, which has helped us improve the clarity and quality of the manuscript.\nRelated discussions (yellow-highlighted in Line 468-471, Page 20) have been added in the revised manuscript as follows:\n“Fig. 5 | Performance and stability of BSCF-RCx cells in PCFC mode. a Current-voltage polarization and power density curves at different temperatures for the BSCF-RC30 air electrode. b Corresponding electrochemical impedance spectra at open circuit voltage. c Peak power density comparison for BSCF-RCx series cells operated at 650 and 500 °C.”"}, {"round": 2, "reviewer_comments": {"reviewer_1": "All the issuses have been well addressed.", "reviewer_2": "The revised manuscript has addressed and improved upon more than the suggested changes. Accordingly, it should be accepted as it is.", "reviewer_3": "The authors have well addressed the technical comments, and the current version is acceptable to move forward for publication now."}, "author_response": ""}]

Decision letter