Highly ionic-dispersed oxygen electrode for reversible proton ceramic electrochemical cells

Domain: biology
Published: 2026-03-14
Editorial outcome: major_revision

Abstract
The key challenges for commercializing reversible proton ceramic electrochemical cells (R-PCECs) are the insufficient proton conductivity and inferior thermomechanical stability of oxygen electrodes in air with water vapor. We report a multielement micro-doped BaCoO3-δ-based perovskite material, in which disorder is induced in the ionic substructure to maximize the oxygen-water reaction activity. Atom probe tomography and density functional theory calculations reveal that reduced proton adsorption/diffusion energy barriers are triggered by homogeneous ion distributions in the perovskite oxide. Moreover, the thermally driven mild oxygen release can be further offset by beneficial proton uptake, thereby increasing the thermomechanical durability of the oxygen electrode. The resulting R-PCECs obtain a peak power density of 1.56 W cm-2 and an electrolysis current density of 2.0 A cm-2@1.3 V at 600 °C while demonstrating long-term stability exceeding 780 hours, with degradation rates of 19.3 and 16.9 μV h-1 in fuel cell and electrolysis modes, respectively.

Introduction
Solid oxide cells, offering high-efficiency power generation and green hydrogen production, represent a promising sustainable energy conversion technology1, 2. Reversible proton ceramic electrochemical cells (R-PCECs) constitute a substantial advancement by reducing the operating temperature from 750-1000 °C to the intermediate range of 350-600 °C, thus mitigating material degradation and sealing challenges associated with high-temperature operation3, 4, 5, 6, 7, 8, 9. However, the commercialization of R-PCECs is hindered by oxygen electrodes lacking simultaneous high catalytic activity and reliable operational stability10, 11, 12. This challenge is particularly demanding given the requirements for H+/O2-/e- triple-conductivity while preserving the microstructural integrity in harsh humidified oxidizing environments13, 14.

The typical oxygen electrodes for R-PCECs predominantly have ABO3-type cubic perovskite structures. The highly symmetric lattice generally engenders isotropic ion migration and displays a greater tendency for atomic orbital overlap, which leads to promising potential of proton conduction15. Nevertheless, investigations have revealed a marked discrepancy between theoretical predictions and practical applications, such as Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF), which has a suboptimal protonic conductivity16. A plausible explanation for this discrepancy lies in the nanoscale heterogeneity of oxygen sublattice ordering, despite apparent macroscopic crystallographic regularity. Localized strain in B-site cations and the spatial distribution of redox-active transition metal species with different valences engender microscopic domains of various degrees of order/disorder. This nanoscale heterogeneity subsequently has profound implications for both the electrochemical kinetics and thermodynamic stability of materials.

Previous investigations revealed that B-site cation substitution strategies could modulate both the electrocatalytic activity and structural resilience. Zhou et al.'s investigation indicated that the improved oxygen reduction reaction (ORR) activity of SrCo0.8Nb0.1Ta0.1O3-δ can be attributed to the enhanced disorder in its ionic arrangement, thus establishing correlations between local structural disorder and enhanced catalytic kinetics17, 18. Conversely, the heterogeneous distribution of transition-metal elements induces pronounced substantial thermal stress gradients during thermal cycling, resulting in electrode delamination. Consequently, optimization of the cationic distribution homogeneity represents a multifunctional approach for concurrently optimizing the catalytic activity and operational durability of oxygen electrodes materials19.

Herein, we report a multielement micro-doped BaCoO3-δ-based perovskite engineered to induce disorder in its ionic substructure, maximizing the accessibility and utilization efficiency of catalytically active sites. Through atom probe tomography (APT) and first-principles density functional theory (DFT) calculations, BaCo0.8(Zr1/6Ti1/6Zn1/6In1/6Cu1/6Mo1/6)0.2O3-δ (BCZTZICM) exhibits dispersed elemental distribution, facilitating low proton diffusion energy barriers. Compositional homogenization at the atomic scale also results in diminished internal thermal stresses, resulting in a low thermal expansion coefficient (TEC) of 17.67×10-6 K-1. Furthermore, X-ray absorption near-edge structure (XANES) and in-situ Fourier transform infrared (FT-IR) spectroscopy confirms an enhanced proton uptake capability, which synergistically mitigates the microstructural degradation induced by electrode cracking. A Ni-BaZr0.1Ce0.7Y0.1Yb0.1O3-δ (BZCYYb)-supported single cell achieves a peak power density (PPD) of 1.56 W cm-2 and an electrolysis current density of 2.0 A cm-2@1.3 V at 600 °C, demonstrating the promising performance of this electrode design.

Results
Materials design and distribution characterization

Given the significant impact of B-site ionic composition on perovskite activity and stability, we categorized B-site cations into active elements (Fe, Co, Ni, etc.) and dopant elements (Zr, Y, Zn, Cu, etc.) to further refine the disorder characteristics within the entropy concept20, 21. The mixed ionic-electronic conductor BSCF and the H+/O2-/e- triple conductor BCFZY was selected, both exhibiting high activity and serving as benchmarks. The B-site composition of BSCF comprises entirely active elements, whereas BCFZY represents a structure regulated by 10% Zr and 10% Y co-doping. Furtherly, the multi-element micro-doped materials with more diverse elemental regulation was designed, where B-site Co ions are modulated by 20% co-doping of (Zr, Ti, Zn, In, Cu, Mo). This configuration includes low valence of Zn2+ and Cu2+, isovalent In3+, Zr4+, and Ti4+, and high valence Mo6+, producing an overall valence effect consistent with 10% Zr and 10% Y co-doping. The design incorporates a gradient ionic radius distribution from 0.6 Å (Cu2+) to 0.72 Å (Zr4+), inducing controllable lattice strain without destabilization, while balancing covalency and oxygen vacancy formation through transition metal elements (Zn, Cu, Mo) and non-transition metal elements (Zr, Ti, In). This yielded the multi-element micro-doped BCZTZICM, which together with BSCF and BCFZY establishes a B-site elemental dispersion gradient design concept. Differs fundamentally from conventional high-entropy or single-dopant approaches. Instead of maximizing configurational entropy, multi-element micro-doping strategy distributes a low total doping concentration (20 % of B-sites) across six distinct elements (Zr, Ti, Zn, In, Cu, Mo) to prioritize ionic dispersion. This design suppresses ion aggregation and disrupts defect clusters while preserving 80 % active Co for high catalytic activity. The resulting homogeneous local environment, combined with moderate lattice strain from valence and ionic-radius diversity, enables uniform proton-transport pathways with low energy barriers and yields a low, near-linear thermal expansion. Ionic dispersion is the central mechanism, synergistically enhanced by valence diversity and controlled strain.

X-ray diffraction (XRD) Rietveld refinement reveals consistent Pm-3m space group symmetry across all the compositions, with Fig. 1a and Supplementary Fig. 1a-b demonstrating comparable lattice parameters: BSCF (a=3.986 Å), BCFZY (a=4.116 Å), and BCZTZICM (a=4.0944 Å). High-resolution transmission electron microscopy imaging of BCZTZICM (Fig. 1b) demonstrates lattice fringes of 0.288, 0.183, and 0.236 nm corresponding to the (011), (120), and (111) planes, respectively, which is consistent with the refinement results. Selected area electron diffraction (Supplementary Fig. 2) reveals the cubic superlattice structure of BCZTZICM along the [1̅ 10] zone axis. Energy-dispersive spectroscopy (Supplementary Fig. 3) confirms the successful incorporation of dopants into the BCZTZICM lattice, with a homogeneous distribution.

Proton conduction bottlenecks in oxygen electrodes critically limit PCEC performance. State-of-the-art electrolytes (BaZr0.8Y0.2O3-δ, BZCYIYb) exhibit proton conductivity of ~0.01-0.02 S cm-1 at 600 °C under humidified conditions5, 22, 23. The proton conductivity of conventional oxygen electrode materials, as evidenced by measurements utilizing diverse techniques-such as the bulk diffusion coefficient testing24 for (Ba/Sr)(Co/Fe/W)O3-δ@PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (~10-5 S cm-1 at 800 °C), hydrogen permeation25 for La0.8Sr0.2Sc0.5Fe0.5O3-δ (1.1×10-4 S cm-1 at 600 °C), and the Hebb-Wagner DC polarization method26 for Ba0.3Gd0.7LaCo2O6-δ (4.4×10-5 S cm-1 at 600 °C) -is consistently and significantly lower than that of state-of-the-art proton-conducting electrolytes. This severe conductivity mismatch confines electrochemical reactions to narrow triple-phase boundaries, restricting the effective reaction zone to approximately 1-2 μm near the electrode-electrolyte interface, limiting electrode utilization and overall cell performance27. Insufficient proton conductivity can be partially attributed to local ionic arrangement and valence inhomogeneity, as illustrated by the clustering and dispersion patterns of B-site ion arrangements in the PCEC structure shown in Fig. 1c-d. Proton conduction in perovskite oxide mechanisms exhibit temperature-dependent dominance. At elevated temperatures, protons can migrate via the vehicle mechanism as mobile OH- species through oxygen vacancy sites. At intermediate temperatures characteristic of PCEC operation (450-600 °C), the prohibitively high activation energy (>1 eV) for OH- transport renders this pathway unfavorable, and the Grotthuss mechanism predominates. Grotthuss mechanism occurs through thermally activated hopping between adjacent oxygen sites, with transport rates being critically dependent on local electronic environments and geometric constraints7, 28. Compositional heterogeneity creates energy barriers limiting proton conductivity, while uniformly dispersed oxygen vacancies enable facile migration pathways (Fig. 1e-f).

FT-IR spectroscopy reveals symmetric B-O-B stretching (610 cm-1) and asymmetric B-O-B’ modes (625 cm-1), showing increasing B-site cation asymmetric arrangement from BSCF through BCFZY to BCZTZICM (Supplementary Fig. 4)29. APT, which offers atomic-scale precision with high-spatial resolution compositional analysis30. Through APT characterization of dense needle-shaped samples prepared from the three materials, detailed ionic distribution information was obtained. Supplementary Fig. 5a-b reveals a modest increase in the oxygen content of BSCF, accompanied by pronounced Co and Fe depletion at a depth of 35 nm. BCFZY exhibits relatively large fluctuations in Co, Fe, and O concentrations, suggesting potential ionic clustering and heterogeneous valence distributions (Supplementary Fig. 6). Examination of the BCZTZICM ionic arrangement along the y- and z-axes (Fig. 1g and Supplementary Fig. 7) reveals a more homogeneous distribution without regional ionic enrichment. To gain deeper insights into the local chemical environment and potential nanoscale segregation tendencies, radial distribution function (RDF) analysis derived from APT was performed30, 31. The key metric is the proximity of the normalized concentration profiles to unity across various radial distances, which indicates similar association tendency between different ion species. As illustrated in Supplementary Fig. 8, with Co ions as the center, the normalized concentrations of high-abundance Co cations and oxygen anions approach unity, while trace-doped In3+ and Zr4+ ions display vibrant concentration profiles within the analyzed region (constrained by limited abundance), collectively indicating a remarkably homogeneous spatial distribution of Co ions, O2-, and dopant ions. In contrast, BSCF suggests a slight tendency for Co ions to disperse from Fe, implying a preferential association among like ions (Co-Co and Fe-Fe) within the analyzed region, and O2- demonstrate a stronger association with Co than with Fe. For BCFZY, Co exhibits a preferential association with Y3+, Fe ions and O2-, while displaying a relative exclusion toward Zr4+.

DFT calculations based on optimized structural models (optimized structures: Supplementary Fig. 9a-c, Supplementary Data 1) reveal that BCZTZIRCM exhibits the lowest proton diffusion barrier (Fig. 1h) along bulk transport pathways (Supplementary Fig. 10a-c). While BSCF contains low-barrier sites, high-barrier regions obstruct migration, limiting conductive pathways and the overall proton conductivity. Notably, BCZTZICM maintains highly-uniformity, with lower difference between maximum and minimum barrier sites, eliminating bottleneck effects and creating extensive networks of facile transport pathways that enable rapid proton diffusion in three dimensions.

Metal-oxygen bond length distributions analysis (Fig. 2a) reveals minor Co-O delocalization in BCZTZICM versus significant deviations in BSCF. Electron localization function (ELF) calculations (Fig. 2b-c, Supplementary Fig. 11) were used to assess the internal homogeneity. The standard deviations measurements confirm that the multi-doped systems exhibit increased uniformity of Co-O bonds and oxygen ions. Additionally, visualization diagrams along the (100) plane and bird's-eye view ELF maps are presented in Supplementary Fig. 12. These results demonstrate that the presence of dopant ions on the (100) plane, no significant polarization effects on oxygen ELF, indicating enhanced structural homogeneity in BCZTZICM and the mitigation of potential well effects from doping ions on carrier migration processes. In contrast, BSCF exhibits pronounced O-ELF polarization disparities (Supplementary Fig. 12b and e), resulting in the high ionic migration barrier sites.

BCZTZICM exhibits the highest average O-ELF value compared with BCFZY and BSCF, indicating enhanced oxygen locality that suppresses high-temperature oxygen desorption while preserving proton pathways (Fig. 2d). The oxygen vacancy formation energies demonstrate BCZTZICM’s minimal energy variation between sites (Supplementary Fig. 9), confirming structural homogeneity (Fig. 2e). Thermogravimetric analysis demonstrates a mass loss of 1.12% for BCZTZICM and higher losses for BSCF and BCFZY, indicating an enhanced metal-oxygen bond energy of BCZTZICM. Based on room-temperature XANES, the oxygen non-stoichiometric ratios were obtained. BCZTZICM demonstrates a higher concentration of inherent oxygen vacancies compared to BCFZY, which is beneficial to surface activity (Fig. 2f)28. Oxygen temperature-programmed desorption (O2-TPD) reveals minimal O2 desorption peaks for BCZTZICM (Supplementary Fig. 13).

Thermal-mechanical stability analysis

Thermal-mechanical compatibility represents a critical requirement for practical PCEC implementation, as thermal expansion mismatch between oxygen electrode (20-25×10-6 K-1) and electrolyte (10-12×10-6 K-1) components generates destructive stresses during operational temperature cycling10. Elevated Co/Fe content generates excessive intergranular thermal stress, resulting in electrode fragmentation. Integration of APT results with theoretical calculations reveals that materials exhibiting severe ionic localization (e.g., BSCF) readily form excessive thermal stress ultimately inducing electrode cracking (Fig. 3a).

Thermal expansion tests in Fig. 3b presents TECs in dry air (300-800 °C): BSCF (26.08×10-6 K-1), BCFZY (22.56×10-6 K-1), and BZCYYb (11.98×10-6 K-1). BCZTZICM achieves a remarkably low TEC of 17.67×10-6 K-1 (300-800 °C), lower than those of most reported oxygen electrodes (Supplementary Table 1). BCZTZICM exhibits near-linearity thermal expansion behavior, contrasting with BSCF and BCFZY deviations from linearity above 400 °C. This deviation results from thermally induced lattice oxygen loss, reducing cation coordination numbers and increasing cation electrostatic repulsion, leading to anomalous expansion behavior and potential mechanical instability19. Multielement B-site doping in BCZTZICM effectively suppresses lattice oxygen loss through uniform local ion coordination environment, yielding near-linear behavior.

XANES provides quantitative analysis of valence changes during thermal treatment under different atmospheric conditions. Co K-edge and Fe K-edge XANES spectra for BCZTZICM, BCFZY, and BSCF reveal temperature- and atmosphere-dependent valence transitions (Fig. 3c and Supplementary Fig. 14a-b and 15b-c). Calibration using CoO, Co3O4, and LaCoO3 for Co and FeNbO4 and SrFeO3 for Fe (Supplementary Fig. 15a) established precise relationships between the absorption edge positions and average oxidation states. As shown in Supplementary Fig. 16a-c and 17a-b, the changes in the lattice oxygen content can be accurately determined.

Under dry air conditions, BCZTZICM demonstrates minimal lattice oxygen content change (-0.075) compared with substantial losses for BSCF (-0.247) during heating to 600 °C (Fig. 3d), indicating lower thermal expansion, which were calculated using the equation in Supplementary Note 1. Notably, steam-treated samples exhibit elevated B-site cation valences. This oxidative response upon proton uptake originates from a distinct mechanism: the lowering of antibonding O 2p states hybridized with transition metal 3d orbitals, coupled with modifications to the local coordination environment induced by protonic defect incorporation (Supplementary Note 2)32. This mechanism has been experimentally validated through operando XAS studies on analogous perovskite systems, depending on the specific electronic structure characteristics (Supplementary Note 2)33, 34. After normalization for elemental composition, the lattice oxygen content variations derived from XANES analysis quantify the proton uptake capacity (Fig. 3d). BCZTZICM exhibits a competitive proton uptake capability under humid conditions, with a lattice oxygen content increase of 0.056, compared with only 0.042 for BSCF, when switching from dry to humidified environments at 600 °C. The multi-element micro-doping, by maximizing ionic dispersion, creates this enhanced structural homogeneity, which in turn leads to a greater availability of uniform active sites, resulting in significantly improved proton uptake capability.

Rigorous thermal cycling tests assessed BCZTZICM electrodes durability under conditions simulating realistic PCEC operation. Testing protocols involved 40 rapid thermal cycles between 400 °C and 600 °C with a heating rate of 20 °C min-1 and a cooling rate of 10 °C min-1, representative of aggressive startup/shutdown scenarios in grid-responsive applications35. Chemical compatibility assessment between BCZTZICM and BZCYYb (Supplementary Fig. 18) reveals no emergent peaks or shifts following 1000 °C co-sintering, confirming compatibility. High-temperature XRD (Supplementary Fig. 19) and patterns obtained 200 hours post-calcination at 600 °C (Supplementary Fig. 20) demonstrate structural robustness. The steam tolerance of BCZTZICM was evaluated by high-temperature XRD during cyclic transitions between dry and humid air atmospheres. As shown in Supplementary Fig. 21, the material exhibits only reversible lattice expansion/contraction, without emergence of secondary phases.

ASR increase over 100 hours and 40 complete cycles (Fig. 3e) demonstrated competitive electrode stability: BSCF showed polarization resistance (Rp) 256% degradation (6.4% per cycle), BCFZY 85% (2.125% per cycle), while BCZTZICM exhibited only 28% (0.58% per cycle). The initial, 20 and 40 cycle polarization impedances are shown in the Supplementary Fig. 22. This substantially lower degradation rate confirms the promising long term and multi-thermal cycle stability of the BCZTZICM, enabling reliable operation under demanding conditions. These results demonstrate competitive mechanical durability enabling reliable operation under demanding conditions. Post-cycling scanning electron microscopy (SEM) cross-sectional analysis reveals structural origins of the enhanced durability. BSCF and BCFZY electrodes (Fig. 3f) exhibit extensive cracking networks with 1-2 μm crack widths, complete delamination from electrolyte interfaces, and significant porosity changes, indicating mechanical failure. Surface laminar cracks and internal fissures correlate with elevated thermal stress. In contrast, Fig. 3g shows that BCZTZICM maintains structural integrity with minimal visible defects, confirming its competitive thermomechanical compatibility. The thermal cycling conditions and the corresponding crack widths observed in post-test microstructural analysis for reference electrodes are summarized in Supplementary Table 2. Enhanced thermomechanical stability originates from two synergistic mechanisms: the homogeneous ionic distribution enables uniform oxygen release, yielding a reduced and near-linear thermal expansion coefficient; concurrently, proton uptake capability suppresses Co reduction while the accompanying chemical expansion compensates for thermal contraction during cooling.

Proton uptake conduction dynamics

Proton transport in perovskite-based oxygen electrodes follows the Grotthuss mechanism in both operational modes. Surface proton transport, in contrast, originates from the proton uptake process through H2O dissociative chemisorption to form protonic defects. Fundamental distinctions arise in proton sources between operational modes: in fuel cell mode, protons originate from hydrogen dissociation at the hydrogen electrode, transporting through electrolyte, and forming water at the oxygen electrodes, while water vapor generated as reaction product undergoes subsequent proton uptake, augmenting electrode protonic conductivity. Electrolysis mode generates protons exclusively through water dissociation at the oxygen electrode, with subsequent transport to the hydrogen electrode. The overall efficiency critically depends on electrode ability to rapidly adsorb water vapor, facilitate dissociative chemisorption, and transport protons into the bulk. Physical absorption assessment through TPD-mass spectrometry following water vapor exposure (250 °C, 20% H2O, 2 hours). BCZTZICM exhibits enhanced surface physisorption and water storage capacity, shown in Fig. 4a. DFT calculations provide molecular-level insights into the enhanced water interaction mechanisms. BCZTZICM surfaces exhibit favorable adsorption energetics compared with those of BSCF and BCFZY, with lower hydration reaction energetics, as shown in Fig. 4b.

In-situ FT-IR spectroscopy during high-temperature water exposure provides real-time hydration kinetics and intermediate species formation monitoring. Hydroxyl formation was monitored via O-H stretching vibrations (3200-3700 cm-1) following 10% water vapor introduction at 600 °C21. BCZTZICM demonstrates rapid chemisorption kinetics, achieving steady-state hydroxyl formation within 5 minutes, like BCFZY, versus more than 45 minutes for BSCF. Upon dry air purging, the hydroxyl elimination kinetics reveal proton mobility. BCZTZICM shows the rapid disappearance of the hydroxyl peak (complete elimination within 15 minutes), indicating efficient transport of protons from bulk positions to surface sites, contributing to the ionic conductivity. BSCF retains significant hydroxyl signatures after 45 minutes, suggesting severely limited proton mobility.

Electrochemical Performance Evaluation

Electronic conductivity measurements (Supplementary Fig. 23) reveal the characteristic of BSCF non-monotonic temperature dependence, directly correlated with severe oxygen loss. Excessive oxygen vacancies annihilate electron holes, reducing carrier concentration and conductivity at high temperatures. In contrast, BCZTZICM maintains stable oxygen content at operating temperatures, exhibiting near-linear temperature-dependent conductivity dominated by thermally activated electron hopping. This demonstrate that lattice oxygen desorption in BCZTZICM exhibits a stochastic desorption characteristic. In contrast, the non-linear conductivity responses about temperature in BSCF and BCFZY are associated with the inhomogeneity of local structures.

Electrical conductivity relaxation (ECR) measurements were conducted to extract oxygen surface exchange coefficients (kchem) and bulk diffusion coefficients (Dchem) (Supplementary Fig. 24). Notably, BCZTZICM shows inferior oxygen activity versus BSCF under dry air. Electrochemical impedance spectroscopy (EIS) of symmetric cells determines of electrode-specific contributions. Dry-air testing of Gd0.2Ce0.8O1.9 (GDC) based cells eliminated the confounding effects of proton conduction. An Arrhenius plot (Supplementary Fig. 25a) reveals unfavorable BCZTZICM ORR kinetics under dry conditions due to enhanced metal-oxygen bonding from multi-doping, consistent with DFT calculations. BSCF and BCFZY exhibits better intrinsic ORR activity than BCZTZICM (Supplementary Fig. 25b-d). However, in controlled humidified environments (5% H2O in air), BCZTZICM achieves remarkably low area-specific resistances (ASRs) of 0.045, 0.078, 0.163, 0.354, and 0.998 Ω cm2 across the complete operational temperature range of 700-500 °C (Fig. 4d-e), substantial outperforming BCFZY (0.337 Ω cm2) and BSCF (0.311 Ω cm2) at 600 °C based on the BZCYYb electrolyte (Supplementary Fig. 26a-b). The enhanced proton transport kinetics becomes pronounced at lower temperatures (500 °C) where proton conduction dominates. Disparities between GDC and BZCYYb symmetric cells testing demonstrate BCZTZICM’s advantages in protonic conduction. In addition, steam introduction reduces the ASRs versus dry air (Supplementary Fig. 27), through accelerated surface exchange and bulk conduction.

Distribution of relaxation times (DRT) analysis deconvolves impedance spectra into distinct processes contributions. High-frequency regions (>1000 Hz) reflect charge transfer processes at electrode-electrolyte interfaces, intermediate-frequencies (1-1000 Hz) correspond to bulk ion transport within electrode, and low-frequencies (<1 Hz) indicate gas-phase diffusion limitations36, 37. As shown in Fig. 4f, BCZTZICM exhibits substantially reduced contributions across all frequency ranges, with dramatic improvements in the mid-frequency regions corresponding to bulk proton transport. The weaker low-frequency response reflects enhanced surface proton uptake, aligning with the in-situ FT-IR and H2O-TPD results. This confirms atomic-scale compositional homogenization creates extensive facile proton transport pathways optimizes proton uptake, and retains sufficient intrinsic oxygen activity. Comparison with state-of-the-art electrode materials confirm the promising performance of BCZTZICM (Supplementary Fig. 28).

Single cells with Ni-BZCYYb hydrogen electrodes, BZCYYb electrolytes, and BCZTZICM oxygen electrodes underwent I-V and EIS testing. BCZTZICM oxygen electrodes achieve a 1.3 W cm-2 PPD at 650 °C with a 0.07 Ω cm2 polarization resistance (Supplementary Fig. 29) in fuel cell mode, versus BCFZY (0.9 W cm-2) and BSCF (0.81 W cm-2), as shown in Supplementary Fig. 30-31. The current density in electrolysis mode is 1.8 A cm-2 at 650 °C (Supplementary Fig. 32a), versus BSCF (1.32 A cm-2) and BCFZY (1.29 A cm-2) (Supplementary Fig. 33-34), respectively. Supplementary Fig. 35 presents cross-sectional SEM images of the comparable electrolyte (13-14 μm) and electrode thickness (10-11 μm).

R-PCECs incorporating thin film electrolytes were fabricated following the methodology described by Duan et al.4, 38 The measurements in Fig. 5a yield PPDs of 1.56, 1.24, 0.93, and 0.64 W cm-2 at 600, 550, 500, and 450 °C, respectively. As shown in Fig. 5c, the total cell polarization resistance reaches only 0.11 Ω cm2 at 600 °C. Electrolysis measurements (Fig. 5b) at a thermoneutral voltage (1.3 V) yield results of 2.0-0.6 A cm-2 at 600-450 °C. The enhanced oxygen evolution reaction (OER) activity reflects the enhanced proton uptake capabilities and the optimized proton transport properties.

Faradaic efficiency measurements confirm hydrogen production selectivity with minimal parasitic losses. The efficiencies reach 85.4±0.55% at optimal operating conditions (1.2 A cm-2, 550 °C) and remain above 83% even at aggressive current densities (1.6 A cm-2). The corresponding hydrogen production rate of 9.264 mL cm-2 min-1 and electrical energy-to-chemical energy conversion efficiency reached 58-68%, demonstrating capability for high-purity hydrogen production applications (Supplementary Fig. 36)39. Supplementary Fig. 37 presents cross-sectional SEM images of the tested cell featuring electrolyte (~7 μm) and electrodes (9-10 μm). Performance and impedance comparison establishes BCZTZICM among higher-performing PCEC systems (Supplementary Fig. 38 and Supplementary Table 3). A comparison of the low-temperature (450 °C) performances is shown in Supplementary Table 4. Temperature-dependent performance analysis reveals competitive kinetics across the complete operational window, with particularly competitive low-temperature capabilities that enable operation below 500 °C, providing advantages for thermal management requirements in distributed energy conversion. Fig. 5d demonstrates reversible operation between fuel cell mode (0.5 A cm-2) and electrolysis mode (-0.5 A cm-2) over 120 hours, showing competitive flexibility with minor hysteresis, confirming suitability for grid-responsive energy storage40. The long-term galvanostatic testing at 550 °C (Fig. 5e) demonstrates stable performance over 780 hours, with degradation rates of 19.3 and 16.9 μV h-1 in fuel cell mode and electrolysis mode, respectively. The degradation originates from seal failure induced by silver sealant oxidation, creep, thermal stress mismatch, and mechanical deterioration under high-temperature and humid atmospheres. Recent studies on oxygen electrode materials and corresponding long-term stability in PCEC summarized in Supplementary Table 5. The reliable stability observed at reduced temperatures, coupled with the competitive activity of BCZTZICM, highlights the promising application prospects of PCEC for efficient and sustainable energy conversion.

Methods
Materials Synthesis:

The mentioned materials BCFZY, BSCF, and BCZTZICM were synthesized using sol-gel methods, consistent with literature description41. In the case of BCZTZICM, the ionic sources employed for dissolution and thorough mixing in the solution include Ba(NO3)2-99%, Co(NO3)2·6H2O-98.5%, Zn(NO3)2·6H2O-99%, Cu(NO3)2·3H2O-99% (above purchased from Sinopharm Chemical Reagent Co., Ltd., analytical grade), Zr(NO3)4·5H2O-99.99% (purchased from Shanghai Macklin Biochemical Co., Ltd, analytical grade), C16H36O4Ti-99%, In(NO3)3·2H2O-99.9% (purchased from Shanghai Aladdin Biochemical Co., Ltd., analytical grade), and H8MoN2O4-99.99% (purchased from Alfa Aesar (China) Chemical Co., Ltd.). Ethylenediaminetetraacetic acid (EDTA) and citric acid monohydrate (CA) were used as complexing agents, and ammonia hydroxide adjusted the pH to about 7. The overall molar ratio of CA: EDTA: total metal ions was maintained at 2:1:1. Sr(NO3)2-99%, Y(NO3)3·6H2O-99%, and Fe(NO3)3·9H2O-98.5% used for synthesizing BSCF and BCFZY were purchased from Sinopharm Chemical Reagent Co., Ltd. The precursor BZCYYb was obtained via solid-state reaction method using the oxide as a cation source and ethanol as a solvent. BaCO3-99%, ZrO2-99%, CeO2-99.9% were purchased from Sinopharm Chemical Reagent Co., Ltd., and Y2O3-99.99%, Yb2O3-99.99% were purchased from purchased from Shanghai Aladdin Biochemical Co., Ltd. The precursors were heated at 180 °C for 5 hours, then calcined at 1000 °C for 5 hours in a muffle furnace and annealed to obtain the desired powder. GDC powder was supplied by NexTech Materials, Ltd.

Cell Fabrication:

Hydrogen electrode supported single cells were prepared by co-pressing NiO+BZCYYb (60:40 wt.% ratio) with BZCYYb (0.012 g) at 3 Mpa and calcining at 1475 °C for 10 hours. The perovskite powder was mixed uniformly with isopropyl alcohol-99.7%, ethylene glycol-99.5%, and propylene glycol solvents-99% (purchased from Sinopharm Chemical Reagent Co., Ltd., analytical grade) using ball milling at 400 rpm for 30 minutes. Then the resulting mixture was sprayed on the electrolyte surface and co-sintered at 1000 °C for two hours to obtain full cells with 0.45 cm-2 active area. The symmetrical cells were pressed at 3 Mpa from BZCYYb powder mixed with 1wt.% NiO for sinterability and calcined at 1475 °C for 10 hours, with active area of 1.094 cm-2.

Characterization:

Room-temperature X-ray diffraction (XRD) data were obtained by the Bruker D8 advances, and its high-temperature component (Rigaku D/max 2500 V) was used to test the crystal structure of the sample at high temperatures. The material's microstructure, representative lattice diffraction fringes, and elemental distributions were obtained through high-resolution transmission electron microscopy (HR-TEM, JOEL, JEM-2100) with energy-dispersive spectroscopy (EDS). Specimen tips were fabricated via focused ion beam (FIB, Thermo Fisher Scientific, Helios CX) milling, followed by atom probe tomography (APT, LEAP 6000 XR) analysis under vacuum chamber employing laser-assisted field evaporation to acquire elemental composition and three-dimensional atomic positions. Scanning electron microscopy (SEM, ZEISS, Sigma 560, ZEEPTOOLS, ZEM15C) was employed for post-analysis characterization of the cells cross-sectional morphology. The thermal expansion coefficient (TEC) of samples from 25 °C to 1000 °C was obtained by a 402C dilatometer, NETZSCH. For Fourier Transform Infrared (FT-IR) spectral measurements, powders were mixed with KBr and then analyzed on a Spectrum 3 spectrometer (PerkinElmer) with high-temperature assemblies in the frequency range of 500-4000 cm-1 with a resolution of 2 cm-1. Hard X-ray absorption spectroscopy was performed at the Beamline station BL11B of the Shanghai Synchrotron Radiation Facility. The acquired hard XAS data were processed according to the standard procedures using the ATHENA module implemented in the IFEFFIT software packages42. The photon energy of the Co K-edge and Fe K-edge was calibrated using metal foils as references. A thermogravimetric analysis (TGA, NETZSCH, STA 449 F3) was employed to investigate the weight loss of the materials at the range of 100 and 1000 °C. H2O-temperature programmed desorption (H2O-TPD) was conducted to observe the adsorbed H2O desorption temperature and desorption amount of BCZTZICM, BSCF, and BCFZY through mass spectrum (Hiden Analytica, HPR20). In-situ near-ambient pressure X-ray photoelectron spectroscopy (NAP-XPS, SPECS) was performed in 0.3 mbar air atmosphere. During the in-situ process from 25 °C to 600 °C, continuous data acquisition was performed. The final scan collected after the sample was cooled to 25 °C and fully stabilized was designated as the room-temperature data. Fig. 1c-d and 3a were created using PowerPoint.

Computational methods:

The density functional theory (DFT) calculations were conducted using the Vienna Ab initio Simulation Package (VASP 5.4.4) with two Hygon 7490 64-core processors43, 44. The electronic exchange and correlation were treated using the Perdew-Burke-Ernzerhof (PBE) functional within the generalized gradient approximation (GGA) framework45. Core-valence electron interactions were described by the projector-augmented wave (PAW) pseudopotentials46, 47, and the electronic wavefunctions were expanded in a plane wave basis set with a kinetic energy cutoff of 500 eV. The perovskite structures were constructed with a 3×3×3 supercell containing 27 ABO3 formula units, and lattice parameters of 11.8720 Å×11.8720 Å×25.8934 Å (α=β=γ=90°). All model coordinates related to calculation are shown in Supplementary Data 1. Electronic convergence was achieved when the energy difference between successive iterations was less than 10-5 eV, while geometric optimization was considered complete when the forces on individual atoms were below 0.05 eV/Å. Brillouin zone sampling was performed using the Monkhorst-Pack scheme48 with a k-point mesh of 3×3×1. Spin polarization effects were incorporated in all calculations. Atomic coordinates of BCZTZICM, BSCF, and BCFZY for optimization calculation are shown in Supplementary Data 1.

Electrochemical Measurements:

Electrochemical Impedance Spectroscopy (EIS) of the symmetrical cells and single cells was tested by potentiostat (Solartron 1287) and a galvanostat combined with frequency response analyzer (Solartron 1260A) from 1 MHz to 0.01 Hz under an open circuit voltage, and the impedance of the electrolysis cell was measured at a nominal voltage of 1.3 V. The I-V and I-P curves were collected by source meter (Keithey 2440). During the test, dry H2 was fed into the fuel side at a flow rate of 80 ml min-1 (STP), while the oxygen electrodes were fed with flowing air at a flow rate of 80 ml min-1 (STP). The Faradaic efficiency for electrolysis was quantified using online gas chromatography (GC, PANNA, A60 pro). During the measurement, the oxygen electrode side was supplied with 60 mL min-1 flowing air (20% H2O), while the hydrogen electrode side was purged with a flowing argon stream, which acted as a carrier gas directed to the GC for quantitative analysis of the produced hydrogen. The cell was stabilized for 20 minutes at each current density (0.8, 1.2, and 1.6 A cm-2) before data collection. The GC chromatograms showed a distinct hydrogen peak, without observable nitrogen or oxygen peaks, confirming the reliable sealing of the cell and the absence of gas cross-leakage. To ensure statistical reliability, the measurement at each current density was repeated five times. The FE and energy efficiency was calculated referring following formula:

FE(%) = nH2,measured / nH2,theoretical × 100% = nH2,measured × (n×F) / I × t × 100%

EEE(%) = nH2,measured × ΔHH2,LHV / I × V × 100%

Where n is the measured hydrogen production rate (mol s-1), n is 2, F is the Faradaic constant (96,485 C mol-1) and I is the applied current (A), ΔHH2,LHV is the lower heating value reaction enthalpy for steam electrolysis, with a value of 241.8 kJ mol-1, and V is the corresponding voltage (V)4, 13, 39.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors demonstrated a multi-element micro-doping strategy to develop BaCoO3-based perovskite oxygen electrode materials [BaCo0.8(Zr1/6Ti1/6Zn1/6In1/6Cu1/6Mo1/6)0.2O3-δ]. Through six-element synergistic doping, induced atomic-scale uniform ion arrangement and significantly reduced proton diffusion energy barriers, with proton diffusion barriers at various sites differing by 0.11 eV. This demonstrates a novel oxygen ion-proton conduction balancing strategy for PCEC oxygen electrodes. The material experienced rigorous thermal cycling stability tests, exhibiting enhanced thermomechanical stability, and achieved over 780 hours of stable fuel cell power generation and water electrolysis performance. The XANES quantification method for thermal reduction and hydration reactions employed in this manuscript is interesting. The electrode achieved exceptional performance at 450 °C with a peak power density of 0.64 W cm-2 and electrolysis current density of 0.6 A cm-2. The manuscript presents comprehensive content with rigorous argumentation and can be published in Nature Communications after minor revision. The following issues require further clarification to complete this research:\n1. Clarify the distinction and relationship between the proposed multi-element micro-doped BCZTZICM and high-entropy perovskite materials to more precisely position the innovation of this work.\n2. Provide quantitative description of the \"proton conduction bottleneck\" with literature citations and typical proton conductivity values for comparison.\n3. Why do the authors believe protons transport via lattice oxygen hopping rather than combining with oxygen ions as hydroxyl radicals in perovskites?\n4. Do different proton conduction mechanisms exist during fuel cell mode and water electrolysis in PCEC operation?\n5. The discussion of hydration reactions appears abrupt. Please elaborate on the impact of hydration reactions on PCEC reaction kinetics to improve the article logic.\n6. Label the elements corresponding to different colors in APT characterization.\n7. Annotate crack dimensions in Fig. 3f-g.", "reviewer_2": "Overall, this paper convincingly demonstrates the promise of a multielement micro-doping strategy for BCZTZICM oxygen electrodes. The authors present a compelling set of results, showing not only excellent ORR/OER activity but also strong thermal-mechanical stability and durability across repeated thermal cycling. Therefore, I recommend the manuscript undergo major revision before being accepted by high-impact Nature Communications.\n1. While the authors mention the basic rationale for employing a multielement doping strategy to enhance electrocatalytic activity and structural resilience, the criteria for selecting the specific dopants remain unclear. Were these elements chosen primarily because they are transition metals or rare-earths, or was the intention to collectively approach a high-entropy configuration? Co/multi-doping is often used to suppress phase transitions, maintain low TEC, and thereby enhance ion mobility—expected outcomes in this work. To better support the novelty and rigor of the study, additional background or theoretical justification should be provided, either in the introduction or in the initial section on material design (lines 88–91).\n2. In lines 242–245, the discussion on ASR variations raises some ambiguity. Since ASR is a key metric to assess degradation during stability or cycling tests, reporting the changes across the three electrode types is valid. However, the method of obtaining these values should be clarified. Were the reported percentages of resistance increase calculated per cycle, or do they represent cumulative values over the full test? If not calculated per cycle, providing the variation per cycle would be more appropriate to substantiate the superior stability of the designed electrode. Additional explanation here would strengthen the credibility of the results.\n3. In line 334, the authors state that Duan’s method was used to fabricate the ultra-thin electrolyte. However, the exact thickness is not clearly reported—only ~7 μm can be inferred from the SEM images in the supplementary information. If the Nature Energy paper was used as a reference, did the authors attempt to fabricate an even thinner electrolyte (e.g., ~3 μm as reported there)? Demonstrating such a thickness could further enhance the performance and strengthen the impact of this work.\n4. In line 350, the measurement of Faradaic efficiency is indeed essential for evaluating electrolysis performance. However, the methodology should be explicitly described. Was the measurement conducted using gas chromatography or a standard flow meter? How many repetitions were performed under each condition? Given the variability and potential errors associated with Faradaic efficiency, the values should be reported together with error bars to demonstrate data reliability. Furthermore, presenting the corresponding energy efficiency, derived from the Faradaic efficiency, would provide a more comprehensive assessment of the electrolysis performance.\n5. In line 363, regarding the description of Fig. 5e, the single cell demonstrates stability over 120 hours, which is encouraging. However, in the reversible operation test, the voltage response under FC mode appears less stable compared to EC mode. Could the authors provide an explanation for this difference in stability?\n6. For the stability results shown in Fig. 5e and 5f, evaluating the durability of a single cell should go beyond reporting the degradation rate alone. A comparison with degradation rates reported in the literature is equally important to place the results in context. The authors can critically compare the temperature-dependent performance, which is good; however, it is also encouraged to compare the degradation behavior of their electrode with previously reported works. Recent studies by Hanping Ding and Chuancheng Duan provide relevant benchmarks, and including such comparisons would significantly strengthen the discussion and highlight the novelty of this work.", "reviewer_3": "This manuscript is yet another paper on PCFC/PCEC air electrode from Zhou/Shao’s group. This manuscript reports a “micro-doped” version of BaCoO3 perovskite which shows superior performance as an air electrode compared with other “conventional” air electrode. The authors claim that the reason for this improved performance is due to the “homogenous” distribution of oxygen vacancies in the lattice as well as the lowered thermochemical expansion of the new material. However, I am very skeptical about this explanation, since I do not see that the computational and experimental efforts presented in this work can directly lead to these conclusions. Further, I fail to see that the improved performance is indeed from these factors as detailed by the authors (rather than any other factors or other properties of the new perovskite oxide).\nOverall, I think that this manuscript is a perfect example of the “deadlock” situation/conundrum of the current research on air electrode, i.e., if a new material is found to perform well, we do not necessarily know why it is good. Therefore, I suggest that a major revision is needed before this manuscript can be published. More detailed questions below:\n1. The authors claim that the “micro-doping” makes oxygen vacancies distribution more random (note that this is different from the case of doped SrCoO3 cited by the authors, un-doped SrCoO3 is known to form ordered phase, but BSCF does not). However, no experimental evidence can show this. APT is not very good at showing the atomistic arrangement and short-range ordering (I believe what the authors saw is probably due to sample preparation or existence of grain boundaries). The XPS O 1s binding energy cannot be used to prove this, either (see for example Idriss, Surface Science 2021, 712, 121894. https://doi.org/10.1016/j.susc.2021.121894). In fact, the conventional thinking might be the opposite, i.e., doping with multiple cations can lead to more trapping of protons or reduced oxide ion mobility, due to more defect-defect interaction (Coulombic/electrostatic interactions). The computational results are not trustworthy, either. It is well known that performing computations on perovskite lattices with many different dopants (each with a low concentration) is difficult. Since the authors use such a small superlattice, it is impossible to have an exhaustive search of all possible configurations thus statistic errors will be present.\n2. The authors show no appreciable chemical expansion (I suggest that the authors read this paper on chemical expansion: Phys. Chem. Chem. Phys., 2015, 17, 10028-10039) for the BCZTZICM sample. They claim that this is due to the hydration process. However, hydration process is known to be charge neutral (see this in any classic papers on proton conductors), i.e., no B-site cation redox involved. It is very strange to see that Co XANES show a peak shift upon switching to wet atmosphere (the peak shift is very small though, maybe due to formation of a secondary phase in wet atmosphere).\n3. It is very difficult to understand the key contributors to the increased performance of the BCZTZICM sample without knowing some fundamental properties of the perovskite (i.e., oxygen non-stoichiometry, ionic/electronic conductivity). It is entirely possible that the better performance compared with BSCF is NOT from “ionic dispersion”. The authors might want to reflect how to prove this causality."}, "author_response": "Point-to-Point Responses to Reviewers’ Comments and Suggestions\nFirst, we express our sincere gratitude to editor for handling our manuscript and for the insightful\nfeedback provided by the reviewers. Their comments have greatly improved the quality and clarity of\nthe paper.\nBelow, the following are the point-by-point response to reviewers’ comments. All changes in the\nmanuscript have been marked in red for convenience.\nReviewer #1 (Comments to the Author):\nThe authors demonstrated a multi-element micro-doping strategy to develop BaCoO -based perovskite\n3\noxygen electrode materials [BaCo (Zr Ti Zn In Cu Mo ) O ]. Through six-element\n0.8 1/6 1/6 1/6 1/6 1/6 1/6 0.2 3-δ\nsynergistic doping, induced atomic-scale uniform ion arrangement and significantly reduced proton\ndiffusion energy barriers, with proton diffusion barriers at various sites differing by 0.11 eV. This\ndemonstrates a novel oxygen ion-proton conduction balancing strategy for PCEC oxygen electrodes.\nThe material experienced rigorous thermal cycling stability tests, exhibiting enhanced\nthermomechanical stability, and achieved over 780 hours of stable fuel cell power generation and water\nelectrolysis performance. The XANES quantification method for thermal reduction and hydration\nreactions employed in this manuscript is interesting. The electrode achieved exceptional performance\nat 450 °C with a peak power density of 0.64 W cm-2 and electrolysis current density of 0.6 A cm-2. The\nmanuscript presents comprehensive content with rigorous argumentation and can be published in\nNature Communications after minor revision. The following issues require further clarification to\ncomplete this research:\nComment 1:\nClarify the distinction and relationship between the proposed multi-element micro-doped BCZTZICM\nand high-entropy perovskite materials to more precisely position the innovation of this work.\nResponse to C1:\nThanks for the reviewer’s insightful comments. We appreciate the opportunity to clarify our\nmaterial design philosophy and its distinction from conventional high-entropy perovskites.\nAccording to the Boltzmann entropy formula:\n1\n𝑆 = 𝑘 ·lnΩ\n𝑐𝑜𝑛𝑓 𝐵\nWhere k =1.38×10-23 J K-1, is the Boltzmann constant and Ω represents the number of microstates.\nB\nFor molar configurational entropy, k ·N =R, where N is Avogadro's constant and R=8.314 J mol-1 K-\nB A A\n1. Thus, the molar mixing entropy of a single sublattice is:\n𝑛\n𝑆 = −𝑅 ·∑ (𝑥 ·𝑙𝑛 𝑥 )\n𝑐𝑜𝑛𝑓 𝑖 𝑖\n𝑖=1\nBased on the above formula, the calculated configurational entropies of BSCF, BCFZY and\nBCZTZICM as below.\nBSCF (Ba Sr Co Fe O ): S =1.194 R\n0.5 0.5 0.8 0.2 3-δ conf\nBCFZY (BaCo Fe Zr Y O ): S =1.194 R\n0.4 0.4 0.1 0.1 3-δ conf\nBCZTZICM (BaCo (Zr Ti Zn In Cu Mo) O ): S =0.859 R\n0.8 1/6 1/6 1/6 1/6 1/6 0.2 3-δ conf\nThis counterintuitive result reveals a fundamental distinction in our design philosophy.\nConsidering the unique structure of perovskite materials, where A-site and B-site elements occupy\ndistinct spatial positions, we further refined the disorder characteristic inherent in the high-entropy\nconcept by precisely distinguishing B-site components into active site elements (such as Co, Fe, Ni,\netc.) and regulatory dopant elements. Specifically, the B-site of BSCF can be regarded as entirely\ncomposed of active elements, the B-site of BCFZY as a structure regulated by 10% Zr and 10% Y\ndoping. And the B-site of BCZTZICM as Co regulated by 20% mixed dopants (Zr, Ti, Zn, In, Cu, Mo),\ncreating highly dispersed local coordination environments around 80% Co active sites. This establishes\na gradient design concept for B-site dispersion in perovskite materials.\nOur selection of Zr, Ti, Zn, In, Cu, and Mo as dopants was guided by a comprehensive\nconsideration of several key factors.\n1. Valence diversity: In3+, Zr4+, Ti4+ (isovalent); Zn2+, Cu2+ (acceptor doping); Mo6+ (donor doping),\nThe valence is consistent with that of 10% Zr and 10% Y co-doped BCFZY, while simultaneously\naccommodating diverse valence configurations.\n2. Ionic radii gradient: 0.72 Å (Zr4+) to 0.60 Å (Cu2+)-inducing controlled lattice micro strain without\ndestabilization.\n3. Electronic structure modulation: Transition metals (Zn, Cu, Mo) vs. non-transition metals (Zr, Ti,\nIn)-balancing covalency and oxygen vacancy formation energy.\n2\n4. Each element individually intercalates into the ABO sub-lattice, disrupting the continuous ACoO\n3 3-\nsub-lattice while suppressing excessive expansion of adjacent sub-lattices. This process inhibits\nδ\nproton conduction by restraining continuous oxygen vacancy formation and concurrently\ndestabilizes the oxygen electrode structure.\nAction: The following description has been incorporated into the manuscript to present the design logic\nof the material to the readers.\nPage 4-5 Line 77-91:\nMaterials design and distribution characterization\nGiven the significant impact of B-site ionic composition on perovskite activity and stability,\nwe categorized B-site cations into active elements (Fe, Co, Ni, etc.) and dopant elements (Zr, Y, Zn,\nCu, etc.) to further refine the disorder characteristics within the entropy concept. The mixed ionic-\nelectronic conductor BSCF and the H+/O2-/e- triple conductor BCFZY was selected, both exhibiting\nexcellent activity and serving as benchmarks. The B-site composition of BSCF comprises entirely\nactive elements, whereas BCFZY represents a structure regulated by 10% Zr and 10% Y co-doping.\nFurtherly, the multi-element micro-doped materials with more diverse elemental regulation was\ndesigned, where B-site Co ions are modulated by 20% co-doping of (Zr, Ti, Zn, In, Cu, Mo). This\nconfiguration includes low valence of Zn2+ and Cu2+, isovalent In3+, Zr4+, and Ti4+, and high valence\nMo6+, producing an overall valence effect consistent with 10% Zr and 10% Y co-doping. The design\nincorporates a gradient ionic radius distribution from 0.6 Å (Cu2+) to 0.72 Å (Zr4+), inducing\ncontrollable lattice strain without destabilization, while balancing covalency and oxygen vacancy\nformation through transition metal elements (Zn, Cu, Mo) and non-transition metal elements (Zr, Ti,\nIn). This yielded the multi-element micro-doped BCZTZICM, which together with BSCF and BCFZY\nestablishes a B-site elemental dispersion gradient design concept.\nComment 2:\nProvide quantitative description of the \"proton conduction bottleneck\" with literature citations and\ntypical proton conductivity values for comparison.\nWe are grateful to the reviewer for valuable comment regarding the quantitative description of\nthe proton conduction bottleneck.\nThe proton conduction bottleneck in oxygen electrodes represents a critical performance\nlimitation in protonic ceramic electrochemical cells (PCECs). State-of-the-art electrolytes such as\nBaZr Y O (BZY20) and BaZr Ce Y Yb O (BZCYYb) exhibit proton conductivity of\n0.8 0.2 3-δ 0.1 0.7 0.1 0.1 3-δ\napproximately 0.01-0.02 S cm-1 at 600 °C under humidified conditions (Enrico Traversa et al. Nature\nMater. 2010, 9, 846-852., Donglin Han et al. J. Mater. Chem. A, 2024, 12, 5875., J. Mater. Chem. A,\n2018, 6, 18571). In contrast, conventional oxygen electrodes demonstrate significantly lower proton\nconductivity values. The proton conductivity of oxygen electrode materials, as measured by various\nmethods including bulk diffusion coefficient (D ) testing under humid atmospheres (Ling Zhao et\nchem\nal. Adv. Mater. 2024, 2405052), hydrogen permeation membranes (Lei Bi et al. SusMat. 2023, 3, 697-\n708.), and the Hebb-Wagner DC polarization method (Tadeusz Miruszewski et al. J. Mater. Chem. A,\n2024, 12, 13488), is significantly lower than that of the electrolyte. This dramatic conductivity\nmismatch (2-4 orders of magnitude) creates a severe transport bottleneck, confining electrochemical\nreactions to the narrow triple-phase boundaries (TPBs) where electrolyte, electrode, and gas phase\nmeet. Consequently, the effective reaction zone is restricted to approximately 1-2 μm near the\nelectrode-electrolyte interface, severely limiting electrode utilization and overall cell performance.\nThis quantitative analysis underscores the critical need for developing triple-conducting (H+/O2-/e-)\noxygen electrodes to extend the electrochemically active region.\nAction: We have made the following revisions to the manuscript:\nPage 5, Line 101-108:\nProton conduction bottlenecks in oxygen electrodes critically limit PCEC performance. State-of-\nthe-art electrolytes (BaZr Y O , BZCYYb) exhibit proton conductivity of ~0.01-0.02 S cm-1 at\n0.8 0.2 3-δ\n600°C under humidified conditions. Whereas conventional oxygen electrodes demonstrate a lower 10-\nand Hebb-Wagner polarization methods. This severe conductivity mismatch confines electrochemical\nreactions to narrow triple-phase boundaries, restricting the effective reaction zone to approximately 1-\n2 μm near the electrode-electrolyte interface, limiting electrode utilization and overall cell performance.\nWhy do the authors believe protons transport via lattice oxygen hopping rather than combining with\noxygen ions as hydroxyl radicals in perovskites?\nResponse to C3:\nWe appreciate the reviewer for this valuable feedback.\nOur conclusion that protons transport via lattice oxygen hopping (Grotthuss mechanism) rather\nthan bonding with lattice oxygen ions and transporting between vacancies as hydroxyl (Vehicle\nmechanism) is based on several complementary lines of evidence:\nFirst, the Grotthuss-type mechanism, whereby protons hop between adjacent lattice oxygen sites,\nis established in perovskite proton conductors through extensive experimental and computational\nstudies (K.D Kreuer, Solid State Ionics, 1999, 125, 285-302., Yashima M., Nat. Commun. 2023 14,\n7466., Yoshihiro Yamazaki, et al. Nat. Mater., 2025. https://doi.org/10.1038/s41563-025-02311-w). It\nhas been directly observed through quasi-elastic neutron scattering (QENS) and neutron diffraction in\nsimilar perovskite systems. The characteristic activation energies we observe (typically 0.4-0.6 eV)\nare consistent with this proton hopping mechanism rather than the higher barriers (> 1 eV) expected\nfor hydroxyl group migration.\nSecond, electrochemical impedance spectroscopy results and temperature-dependent\nconductivity measurements demonstrate proton transport behavior that aligns with the oxygen-hopping\nmodel (Zongping Shao, et al. Applied Catalysis B: Environmental, 2023, 331, 122682.). The isotope\neffect studies (H/D exchange) further confirm that the rate-limiting step involves proton transfer\nbetween oxygen sites rather than the collective motion of OH groups. Additionally, computational\nstudies using density functional theory (DFT) on our specific material composition indicate that the\nenergy landscape favors sequential proton hopping with lower migration barriers compared to\nhydroxyl group diffusion.\nWe acknowledge that under certain conditions, particularly at oxygen electrode surfaces during\nelectrochemical reactions, transient hydroxyl species may form as reaction intermediates. However,\nfor bulk proton conduction through the perovskite lattice-which is the primary transport pathway we\n5\nare discussing-the lattice oxygen hopping mechanism is energetically favorable and kinetically\ndominant.\nComment 4:\nDo different proton conduction mechanisms exist during fuel cell mode and water electrolysis in PCEC\noperation?\nResponse to C4:\nThank you for your valuable insight and comments. We appreciate it for pointing this out.\nThe bulk proton conduction mechanism in the perovskite-based oxygen electrode is the same in\nboth fuel cell and electrolysis modes, primarily following the Grotthuss mechanism. Protons migrate\nby forming bonds with lattice oxygen ions and hopping from one oxygen site to an adjacent one\nthrough thermal activation. This process involves the sequential breaking and forming of O-H bonds,\naccompanied by cooperative lattice relaxation. However, significant differences exist in the directional\ntransport pathways, electrode reactions, and proton sources between the two operating modes.\nIn fuel cell mode, protons originate from hydrogen dissociation at the hydrogen electrode and are\ntransported through the dense electrolyte membrane to the oxygen electrode, where they react with\noxygen to form water. While moderate water vapor partial pressures can generate protonic defects that\nenhance ionic conductivity, the primary proton supply comes from hydrogen dissociation rather than\nwater uptake. In contrast, during electrolysis mode, protons are generated exclusively through water\ndissociation at the oxygen electrode, creating protonic defects that are subsequently transported\nthrough the electrolyte to the hydrogen electrode for hydrogen evolution. In this mode, proton\ngeneration relies entirely on water vapor uptake from the feed atmosphere.\nThis fundamental distinction in proton source and transport direction constitutes a key operational\ndifference between fuel cell and electrolysis modes in PCECs, with important implications for water\nmanagement strategies and electrode optimization in each operating regime.\nAction: We have made the following revisions to the manuscript:\nPage 13, Line 267-275:\nProton transport in perovskite-based oxygen electrodes follows the Grotthuss mechanism in both\noperational modes, involving thermally-activated hopping between lattice oxygen sites via sequential\nO-H bond cleavage/formation coupled with cooperative lattice relaxation. However, fundamental\ndistinctions arise in transport directionality and proton sources. In fuel cell mode, protons originate\nfrom hydrogen dissociation at the anode, traverse the electrolyte, and react with oxygen at the cathode\nto form water. Conversely, electrolysis mode generates protons exclusively through water dissociation\nat the oxygen electrode, with subsequent transport to the hydrogen electrode. The efficiency depends\non electrode ability to rapidly adsorb water vapour, facilitate dissociative chemisorption, and transport\nprotons into the bulk. Physical absorption assessment through TPD-mass spectrometry following water\nvapour exposure (250 °C, 20% H O, 2 hours).\nThe discussion of hydration reactions appears abrupt. Please elaborate on the impact of hydration\nreactions on PCEC reaction kinetics to improve the article logic.\nResponse to C5:\nThank you very much for your valuable comments on the structure of the article.\nAction: We have added the following parts of the manuscript.\nPage 13, Line 267-275:\nProton uptake conduction dynamics\nProton transport in perovskite-based oxygen electrodes follows the Grotthuss mechanism in both\noperational modes, involving thermally-activated hopping between lattice oxygen sites via sequential\nO-H bond cleavage/formation coupled with cooperative lattice relaxation. However, fundamental\ndistinctions arise in transport directionality and proton sources. In fuel cell mode, protons originate\nfrom hydrogen dissociation at the anode, traverse the electrolyte, and react with oxygen at the cathode\nto form water. Conversely, electrolysis mode generates protons exclusively through water dissociation\nat the oxygen electrode, with subsequent transport to the hydrogen electrode. The efficiency depends\non electrode ability to rapidly adsorb water vapour, facilitate dissociative chemisorption, and transport\n7\nprotons into the bulk. Physical absorption assessment through TPD-mass spectrometry following water\nvapour exposure (250 °C, 20% H O, 2 hours).\nLabel the elements corresponding to different colors in APT characterization.\nResponse to C6:\nThanks to the reviewer for pointing this out, we have presented the distribution of all elements in the\ntext.\nAction: We have completed the following figures of the manuscript.\nSupplementary information\nSupplementary Fig. 5.\n3D reconstruction of the APT data showing mapping of the BSCF sample with Ba, Sr, Co, Fe, and O\nionic distributions. Integrated line profiles show chemical composition at the (a) apex and (b)\nmidsection regions of the analyzed needle.\n3D reconstruction of the APT data showing mapping of the BCFZY sample with Ba, Co, Fe, Zr, Y, and\nO ionic distributions. Integrated line profiles show chemical composition at the apex regions of the\nanalyzed needle.\nSupplementary Fig. 7.\n3D reconstruction of the APT data showing mapping of the BCZTZICM sample with Ba, Co, Zr, Ti,\ny-axis of the analyzed needle.\nComment 7:\nAnnotate crack dimensions in Fig. 3f-g.\nResponse to C7:\nWe appreciate the reviewer for this valuable feedback.\nAction: We have made the following revisions to the manuscript.\nPage 11\nFig. 3 Suppressed thermal expansion and improved cycling stability through compositional homogenization.\n(a) Schematic of electrode tearing driven by increased internal thermal stress resulting from continuous oxygen\nvacancies. (b) Thermal expansion coefficient curves measured in air over 300-800 °C. (c) Co K-edge XANES spectra\n10\nof BCZTZICM samples: as-prepared (room temperature), dried at 600 °C, and treated under 20% H O atmosphere.\n2\n(d) Average B-site valence changes derived from XANES analysis. The Rp response of BCZTZICM, BCFZY, and\nBSCF-based symmetric cell electrodes during 40 thermal cycles between 600 °C and 400 °C. Cross-sectional image\nof cells after cycling (f) BSCF (left), BCFZY (right), and (g) BCZTZICM-based electrode symmetric cell with cracks\ninside the electrode bulk and at the electrode-electrolyte interface.\nPage 13, Line 262-264\nBSCF and BCFZY electrodes (Fig. 3f) exhibit extensive cracking networks with 1-2 μm crack widths,\ncomplete delamination from electrolyte interfaces, and significant porosity changes, indicating\nmechanical failure.\nReviewer #2 (Comments to the Author):\nOverall, this paper convincingly demonstrates the promise of a multielement micro-doping strategy\nfor BCZTZICM oxygen electrodes. The authors present a compelling set of results, showing not only\nexcellent ORR/OER activity but also strong thermal-mechanical stability and durability across\nrepeated thermal cycling. Therefore, I recommend the manuscript undergo major revision before being\naccepted by high-impact Nature Communications.\nComment 1:\nWhile the authors mention the basic rationale for employing a multielement doping strategy to enhance\nelectrocatalytic activity and structural resilience, the criteria for selecting the specific dopants remain\nunclear. Were these elements chosen primarily because they are transition metals or rare-earths, or was\nthe intention to collectively approach a high-entropy configuration? Co/multi-doping is often used to\nsuppress phase transitions, maintain low TEC, and thereby enhance ion mobility—expected outcomes\nin this work. To better support the novelty and rigor of the study, additional background or theoretical\njustification should be provided, either in the introduction or in the initial section on material design\n(lines 88–91).\nResponse to C1:\nWe sincerely thank the reviewer for this insightful comment regarding the rationale behind our\n11\nmultielement doping strategy. The reviewer has raised a critical point about the selection criteria for\nthe specific dopants, which indeed lies at the core of our material design philosophy. We appreciate\nthe opportunity to clarify that our approach is distinct from a conventional high-entropy configuration\nand is instead guided by a refined “B-site dispersion gradient” concept.\nAccording to the formula of configuration entropy formula, the configuration entropy of BSCF\nand BCFZY is 1.194 R, while that of BCZTZICM is 0.859 R. We further refined the disorder\ncharacteristic inherent in the high-entropy concept by precisely distinguishing B-site components into\nactive site elements (such as Co, Fe, Ni, etc.) and regulatory dopant elements. Specifically, the B-site\nof BSCF can be regarded as entirely composed of active elements, the B-site of BCFZY as a structure\nregulated by 10% Zr and 10% Y doping. And the B-site of BCZTZICM as Co regulated by 20% mixed\ndopants (Zr, Ti, Zn, In, Cu, Mo), creating highly dispersed local coordination environments around 80%\nCo active sites.\nThe selection of Zr, Ti, Zn, In, Cu, and Mo was not arbitrary but was guided by a synergistic\nconsideration of multiple physicochemical criteria to collectively enhance electrocatalytic activity and\nstructural resilience:\n1. Valence diversity: In3+, Zr4+, Ti4+ (isovalent); Zn2+, Cu2+ (acceptor doping); Mo6+ (donor doping).\nThe valence is consistent with that of 10% Zr and 10% Y co-doped BCFZY, while simultaneously\naccommodating diverse valence configurations.\n2. Ionic radii gradient: 0.72 Å (Zr4+) to 0.60 Å (Cu2+), inducing controlled lattice micro strain without\ndestabilization.\n3. Electronic structure modulation: Transition metals (Zn, Cu, Mo) vs. non-transition metals (Zr, Ti,\nIn) - balancing covalency and oxygen vacancy formation energy.\n4. Each element individually intercalates into the ABO sub-lattice, disrupting the continuous\n3-δ\ninhibits proton conduction by restraining continuous oxygen vacancy formation and concurrently\ndestabilizes the oxygen electrode structure.\nWe believe this refined “B-site dispersion gradient” concept and the detailed justification\nprovided now offer a clearer and more rigorous foundation for our study's novelty. We have integrated\n12\nthese points into the revised manuscript to strengthen the background and theoretical justification as\nsuggested.\nThe following description has been incorporated into the manuscript to present the design logic of the\nmaterial to the readers.\nPage 4-5, Line 77-91:\nMaterials design and distribution characterization\nGiven the significant impact of B-site ionic composition on perovskite activity and stability,\nwe categorized B-site cations into active elements (Fe, Co, Ni, etc.) and dopant elements (Zr, Y, Zn,\nCu, etc.) to further refine the disorder characteristics within the entropy concept. The mixed ionic-\nelectronic conductor BSCF and the H+/O2-/e- triple conductor BCFZY was selected, both exhibiting\nexcellent activity and serving as benchmarks. The B-site composition of BSCF comprises entirely\nactive elements, whereas BCFZY represents a structure regulated by 10% Zr and 10% Y co-doping.\nFurtherly, the multi-element micro-doped materials with more diverse elemental regulation was\ndesigned, where B-site Co ions are modulated by 20% co-doping of (Zr, Ti, Zn, In, Cu, Mo). This\nconfiguration includes low valence of Zn2+ and Cu2+, isovalent In3+, Zr4+, and Ti4+, and high valence\nMo6+, producing an overall valence effect consistent with 10% Zr and 10% Y co-doping. The design\nincorporates a gradient ionic radius distribution from 0.6 Å (Cu2+) to 0.72 Å (Zr4+), inducing\ncontrollable lattice strain without destabilization, while balancing covalency and oxygen vacancy\nformation through transition metal elements (Zn, Cu, Mo) and non-transition metal elements (Zr, Ti,\nIn). This yielded the multi-element micro-doped BCZTZICM, which together with BSCF and BCFZY\nestablishes a B-site elemental dispersion gradient design concept.\nComment 2:\nIn lines 242–245, the discussion on ASR variations raises some ambiguity. Since ASR is a key metric\nto assess degradation during stability or cycling tests, reporting the changes across the three electrode\ntypes is valid. However, the method of obtaining these values should be clarified. Were the reported\npercentages of resistance increase calculated per cycle, or do they represent cumulative values over\n13\nthe full test? If not calculated per cycle, providing the variation per cycle would be more appropriate\nto substantiate the superior stability of the designed electrode. Additional explanation here would\nstrengthen the credibility of the results.\nResponse to C2:\nWe thank the reviewer for this insightful comment regarding the ambiguity in reporting the ASR\nincrease. The reviewer is correct that clarifying the methodology is crucial for assessing the\ndegradation rate. In the original manuscript, the reported percentages (256% for BSCF, 85% for\nBCFZY, and 23% for the BCZTZICM) represented the cumulative increase in ASR over the entire\nduration of the stability test (40 cycles). We have revised the relevant section in the manuscript to\nclarify this point and have added new table/figures/sentence stating the average ASR increase per hour\nor per cycle. The recalculated the ASR degradation rates on a per-cycle basis for all three electrode\ntypes as follows: 6.4% for BSCF, 2.1% for BCFZY, and 0.58% for the BCZTZICM.\nWe believe these additions will significantly strengthen the credibility of our stability assessment\nand provide readers with a more transparent and comprehensive understanding of the electrode\nperformance degradation. We appreciate the reviewer's attention to this important detail, which has\nimproved the scientific rigor of our work.\nWe have added Supplementary Fig. 24, and revised the expression in manuscript as below:\nPage 13, Line 254-259\nASR increase over 100 hours and 40 complete cycles (Fig. 3e) demonstrated superior electrode\nstability: BSCF showed polarization resistance (Rp) 256% degradation (6.4% per cycle), BCFZY 85%\n(2.125% per cycle), while BCZTZICM exhibited only 28% (0.58% per cycle). The initial, 20 and 40\ncycle polarization impedances are shown in the Supplementary Fig. 21. This substantially lower\ndegradation rate confirms the exceptional long term and multi-thermal cycle stability of the\nBCZTZICM, enabling reliable operation under demanding conditions.\nSupplementary information:\nPolarization resistance of BSCF, BCFZY, and BCZTZICM electrodes initially, after 20 and 40 thermal\ncycles, respectively.\nComment 3:\nIn line 334, the authors state that Duan’s method was used to fabricate the ultra-thin electrolyte.\nHowever, the exact thickness is not clearly reported—only ~7 μm can be inferred from the SEM images\nin the supplementary information. If the Nature Energy paper was used as a reference, did the authors\nattempt to fabricate an even thinner electrolyte (e.g., ~3 μm as reported there)? Demonstrating such a\nthickness could further enhance the performance and strengthen the impact of this work.\nResponse to C3:\nThank you for this insightful comment regarding the thickness of the electrolyte and the potential\nfor further performance improvement. In our work, we successfully fabricated an ultra-thin electrolyte\nwith a thickness of approximately 7 μm using Duan’s method as a reference. We agree that pursuing a\nthinner electrolyte, such as the ~3 μm benchmark reported in the Nature Energy paper, is an attractive\n15\ndirection for enhancing performance. However, in our current experimental setup, particularly due to\ninconsistencies in the tape-casting solvent system and challenges in engineering processes related to\nmaterial details, we have not yet been able to reliably produce defect-free electrolytes significantly\nbelow 6 μm. Achieving such thin layers while maintaining structural integrity, uniformity, and\nelectrochemical stability remains technically challenging at this stage of our process optimization.\nNevertheless, we greatly appreciate the reviewer’s suggestion and acknowledge the creative\ncontributions of Prof. Duan and other researchers in advancing low-temperature protonic ceramic\nelectrochemical cells (PCECs). We are continuing to optimize our fabrication protocol, including\nsolvent formulation and processing parameters, and view the pursuit of sub-5 μm electrolytes as a key\nobjective of our ongoing research.\nComment 4:\nIn line 350, the measurement of Faradaic efficiency is indeed essential for evaluating electrolysis\nperformance. However, the methodology should be explicitly described. Was the measurement\nconducted using gas chromatography or a standard flow meter? How many repetitions were performed\nunder each condition? Given the variability and potential errors associated with Faradaic efficiency,\nthe values should be reported together with error bars to demonstrate data reliability. Furthermore,\npresenting the corresponding energy efficiency, derived from the Faradaic efficiency, would provide a\nmore comprehensive assessment of the electrolysis performance.\nResponse to C4:\nWe are grateful to the reviewer for the meticulous review and valuable insights that have allowed\nus to enhance our analysis and reduce the possibility of misinterpretation.\nThe Faradaic efficiency for electrolysis was quantified using online gas chromatography (GC,\nPANNA, A60 pro). During the measurement, the oxygen electrode side was supplied with 60 mL min-\n1 flowing air (20% H O), while the hydrogen electrode side was purged with a flowing argon stream,\n2\nwhich acted as a carrier gas directed to the GC for quantitative analysis of the produced hydrogen. The\ncell was stabilized for 20 minutes at each current density (0.8, 1.2, and 1.6 A cm-2) before data\ncollection. The GC chromatograms showed a distinct hydrogen peak, without observable nitrogen or\noxygen peaks, confirming the excellent sealing of the cell and the absence of gas cross-leakage.\nAs rightly suggested by the reviewer, the energy efficiency was calculated by referring to the paper of\nDuan et al. (Duan, C. et al., Nat Energy, 2019, 4, 230–240. Nat Energy, 2023, 8, 1145–1157.) via the\nfollowing formula:\n𝑛 𝑛 ×(𝑛×𝐹)\n𝐹𝐸(%) =\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n=\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n×100%\n𝑛 I\n𝐻2,𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙\n𝛥𝐻 ×𝑛 𝑛 ×(𝑛×𝐹)\n𝐸𝐸𝐸(%) =\n𝐻2,𝐿𝐻𝑉 𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n=\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n×100%\nI×V I×V\nWhere 𝑛 is the measured hydrogen production rate (mol s-1), 𝑛 is 2, 𝐹 is the\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\nFaradaic constant (96,485 C mol-1) and I is the applied current (A), 𝛥𝐻 is the lower heating\n𝐻2,𝐿𝐻𝑉\nvalue reaction enthalpy for steam electrolysis, with a value of 241.8 kJ mol-1, and V is the\ncorresponding voltage (V).\nAction: We have redrawn Supplementary Fig. 37 and revised the manuscript as below:\nPage 19, Line 365-369\nThe efficiencies reach 85.1% at optimal operating conditions (1.2 A cm-2, 550 °C) and remain\nabove 83% even at aggressive current densities (1.6 A cm-2). The corresponding hydrogen production\nrate of 9.264 mL cm-2 min-1 and electrical energy-to-chemical energy conversion efficiency reached\n58-68%, demonstrating capability for high-purity hydrogen production applications (Supplementary\nFig. 31).\nPage 25, Line 584-599\nThe Faradaic efficiency for electrolysis was quantified using online gas chromatography (GC,\nPANNA, A60 pro). During the measurement, the oxygen electrode side was supplied with 60 mL min-\n1 flowing air (20% H O), while the hydrogen electrode side was purged with a flowing argon stream,\n2\nwhich acted as a carrier gas directed to the GC for quantitative analysis of the produced hydrogen. The\ncell was stabilized for 20 minutes at each current density (0.8, 1.2, and 1.6 A cm-2) before data\ncollection. The GC chromatograms showed a distinct hydrogen peak, without observable nitrogen or\noxygen peaks, confirming the excellent sealing of the cell and the absence of gas cross-leakage. To\nensure statistical reliability, the measurement at each current density was repeated five times. The FE\n17\nand energy efficiency was calculated referring following formula:\n𝑛 𝑛 ×(𝑛×𝐹)\n𝐹𝐸(%) =\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n=\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n×100%\n𝑛 I\n𝐻2,𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙\n𝛥𝐻 ×𝑛 𝑛 ×(𝑛×𝐹)\n𝐸𝐸𝐸(%) =\n𝐻2,𝐿𝐻𝑉 𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n=\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\n×100%\nI×V I×V\nWhere 𝑛 is the measured hydrogen production rate (mol s-1), 𝑛 is 2, 𝐹 is the\n𝐻2,𝑚𝑒𝑎𝑠𝑢𝑟𝑒𝑑\nFaradaic constant (96,485 C mol-1) and I is the applied current (A), 𝛥𝐻 is the lower heating\n𝐻2,𝐿𝐻𝑉\nvalue reaction enthalpy for steam electrolysis, with a value of 241.8 kJ mol-1, and V is the\ncorresponding voltage (V).\nSupplementary information\nSupplementary Fig. 37.\nThe faraday efficiency, corresponding H production rate and energy efficiency of tested hydrogen\n2\nelectrode-supported ultrathin BZCYYb electrolyte single cell with BCZTZICM oxygen electrode.\nComment 5:\nIn line 363, regarding the description of Fig. 5e, the single cell demonstrates stability over 120 hours,\n18\nwhich is encouraging. However, in the reversible operation test, the voltage response under FC mode\nappears less stable compared to EC mode. Could the authors provide an explanation for this difference\nin stability?\nResponse to C5:\nWe appreciate the reviewer's insightful observation regarding the stability difference between fuel cell\n(FC) and electrolysis cell (EC) modes during reversible operation. The reduced voltage stability\nobserved in FC mode compared to EC mode can be attributed to the dynamic water vapor environment\nat the oxygen electrode. During the reversible operation test, the oxygen electrode atmosphere was\nmaintained at 10% H O. In FC mode, the oxygen electrode actively produces water as a reaction\n2\nproduct, leading to continuous fluctuations in the local water partial pressure (p(H O)). This dynamic\n2\nenvironment requires a significant equilibration period to reach steady-state conditions. The elevated\nand fluctuating p(H O) at the oxygen electrode introduces competitive adsorption between water\n2\nmolecules and oxygen species on the electrode surface. This H O-O competitive adsorption\n2 2\nphenomenon affects the oxygen reduction reaction (ORR) kinetics and leads to gradual performance\ndegradation during extended FC operation. In contrast, during EC mode, water is consumed rather than\nproduced at the oxygen electrode, resulting in a more stable and predictable p(H O) environment,\n2\nwhich contributes to the improved voltage stability. We acknowledge that this aspect requires further\ninvestigation and optimization. In future work, we will:\n(1) refine our testing protocols to better control the water vapor environment during reversible\nswitching, and\n(2) develop advanced oxygen electrode materials with improved tolerance to dynamic p(H O)\n2\nconditions and reduced sensitivity to competitive adsorption, thereby enhancing stability during\nreversible operation. We thank the reviewer for highlighting this important aspect, which will guide\nour future research directions.\nComment 6:\nFor the stability results shown in Fig. 5e and 5f, evaluating the durability of a single cell should go\nbeyond reporting the degradation rate alone. A comparison with degradation rates reported in the\nliterature is equally important to place the results in context. The authors can critically compare the\ntemperature-dependent performance, which is good; however, it is also encouraged to compare the\n19\ndegradation behavior of their electrode with previously reported works. Recent studies by Hanping\nDing and Chuancheng Duan provide relevant benchmarks, and including such comparisons would\nsignificantly strengthen the discussion and highlight the novelty of this work.\nResponse to C6:\nThank you for your thoughtful comment and suggestion. We appreciate it for pointing this out.\nWe have summarized the relevant data in Table S3, and would like to express our sincere gratitude to\nall the outstanding researchers in this field whose valuable experimental data and findings have\ncontributed to this work.\nAction: We have added Table S3 and revised the manuscript and supplementary information as below:\nPage 19, Line 380-384\nRecent studies on oxygen electrode materials and corresponding long-term stability in PCEC\nsummarized in Table S3. The remarkable stability observed at reduced temperatures, coupled with the\nsuperior activity of BCZTZICM, highlights the promising application prospects of PCEC for efficient\nand sustainable energy conversion.\nSupplementary information\nPage 44\nelectrolysis modes.\nCurrent density (mA Time Test\nOxygen electrode Electrolyte Degradation rate Ref\ncm-2) (h) condition\n-1385 600 20% H O-Air\n-1385 1200 10% H O-Air\n2\n< 0.07 mV h-1 300 100 H -Air\n< 0.15 mV h-1 -600 260 /\nPNC55 BZCYYb 1.98% 1.4 V 200 30%H O-O\n2 2\n80 mL\n10% H O-Air\n2\n19.3 μV h-1 300/500\n60 mL Dry\nH This\n80 mL Dry work\nAir\n16.9μV h-1 -300/-500\n60 mL Dry\nH\n2\n0.077 mA cm-2 h-1 20 sccm H\n2\n0.75 V 80\n(0.35%) 40 sccm O\nx\n0.134 mA cm-2 h-1 Humidified\n1.25 V 80\n(0.57%) O\n2\n1.03 % 1.3 V\nNAUP-PNC BZCYYb4411 105 20%H O-O\n2 2\n1.24 % 1.4 V\n10% H O-Air\n2\nslight activation 1018\nDry H\n2\nDry H\n2\n1 μV h-1 -1500 5013 40% H O-Air\n2\nPBNO-BCY CCS 10 μV h-1 1041\nReviewer #3 (Comments to the Author):\nThis manuscript is yet another paper on PCFC/PCEC air electrode from Zhou/Shao’s group. This\nmanuscript reports a “micro-doped” version of BaCoO perovskite which shows superior performance\n3\nas an air electrode compared with other “conventional” air electrode. The authors claim that the reason\nfor this improved performance is due to the “homogenous” distribution of oxygen vacancies in the\nlattice as well as the lowered thermochemical expansion of the new material. However, I am very\nskeptical about this explanation, since I do not see that the computational and experimental efforts\npresented in this work can directly lead to these conclusions. Further, I fail to see that the improved\nperformance is indeed from these factors as detailed by the authors (rather than any other factors or\nother properties of the new perovskite oxide). Overall, I think that this manuscript is a perfect example\nof the “deadlock” situation/conundrum of the current research on air electrode, i.e., if a new material\nis found to perform well, we do not necessarily know why it is good. Therefore, I suggest that a major\nrevision is needed before this manuscript can be published. More detailed questions below:\nComment 1:\nThe authors claim that the “micro-doping” makes oxygen vacancies distribution more random (note\nthat this is different from the case of doped SrCoO cited by the authors, un-doped SrCoO is known\n3 3\nto form ordered phase, but BSCF does not). However, no experimental evidence can show this. APT\nis not very good at showing the atomistic arrangement and short-range ordering (I believe what the\nauthors saw is probably due to sample preparation or existence of grain boundaries). The XPS O 1s\nbinding energy cannot be used to prove this, either (see for example Idriss, Surface Science 2021, 712,\n121894. https://doi.org/10.1016/j.susc.2021.121894). In fact, the conventional thinking might be the\nopposite, i.e., doping with multiple cations can lead to more trapping of protons or reduced oxide ion\nmobility, due to more defect-defect interaction (Coulombic/electrostatic interactions). The\ncomputational results are not trustworthy, either. It is well known that performing computations on\nperovskite lattices with many different dopants (each with a low concentration) is difficult. Since the\n22\nauthors use such a small superlattice, it is impossible to have an exhaustive search of all possible\nconfigurations thus statistic errors will be present.\nResponse to C1:\nThank you very much for taking the time to review our manuscript and for making such insightful\nand constructive comments.\nRegarding the cited works on doped SrCoO , in the referenced studies (Adv. Funct. Mater. 2022,\n3-δ\n2210496), SrCo Nb O , SrCo Ta O , and the co-doped SrCo Nb Ta O are disordered\n0.8 0.2 3-δ 0.8 0.2 3-δ 0.8 0.1 0.1 3-δ\nstructures with same Pm-3m space group. The remarkable enhancement in electrocatalytic activity\nobserved in SrCo Nb Ta O provided significant inspiration for our approach. We wish to\n0.8 0.1 0.1 3-δ\nreiterate the core premise of our material design strategy: we aim to compare and discuss the degree\nof ionic arrangement disorder while maintaining the same average crystal structure (i.e., all\ncompositions exhibit a disordered structure discernible by XRD).\nConcerning the experimental validation of randomly distributed oxygen vacancies, it is important\nto note that direct microscopic characterization of oxygen vacancy distributions over large scales\nremains a significant challenge with currently available advanced techniques. Our conclusion\nregarding their highly dispersity is therefore inferred from the cationic dispersion and overall chemical\nhomogeneity in BCZTZICM (Supplementary Fig. 4 and APT results). This is further supported\nindirectly by the thermal expansion coefficient (TEC) curve and the near-linear temperature\ndependence of its electronic conductivity. Since electronic conductivity is influenced by both hole\nconcentration and oxygen vacancy concentration, localized excessive formation of oxygen vacancies\nleads to the annihilation of electron holes, resulting in a conductivity decrease. The observed increase\nin conductivity for BSCF with decreasing temperature can be attributed to a reduction in oxygen\nvacancies and a consequent increase in hole concentration. In contrast, the more linear conductivity\nprofile of BCZTZICM suggests that its conductivity is primarily governed by the thermal activation\nof holes, with a relatively minor influence from oxygen vacancy generation. This behavior is consistent\nwith a more uniform distribution of oxygen vacancies on a macroscopic scale.\nRegarding the concern about potential issues with sample preparation or grain boundary\nrepresentation, The quality and integrity of our APT samples were rigorously ensured. All needle-\n23\nshaped specimens were prepared to be fully dense, following established site-specific protocols using\nfocused ion beam (FIB) lift-out techniques. The localized elemental segregation or compositional\nvariations observed in our samples are not attributable to the presence of grain boundaries.\nCharacteristic features of grain boundaries, as reported in the literature (Adv. Energy Mater. 2025, 15,\n2404410, J. Mater. Chem. A, 2022, 10, 2496–2508), demonstrate distinctly different signatures.\nTo gain deeper insights into the local chemical environment and potential nanoscale segregation\ntendencies, radial distribution function (RDF) analysis derived from APT was performed. As illustrated\nin Fig., with Co ions designated as the reference center, BCZTZICM exhibits no pronounced tendency\ntoward deep association with any single proximate ionic species. The normalized concentrations of\nhigh-abundance Co cations and oxygen anions approach unity, while trace-doped In and Zr ions\ndisplay oscillatory concentration profiles within the analyzed volume (constrained by limited\nabundance), collectively indicating a remarkably homogeneous spatial distribution of Co ions, oxygen\nions, and dopant species. In contrast, the RDF of BSCF suggests a slight tendency for Co ions to\ndisperse from Fe, implying a preferential association among like ions (Co–Co and Fe–Fe) within\nthe analyzed volume. Furthermore, oxygen anions demonstrate a stronger association with Co than with\nFe. For BCFZY, the gradually rising normalized concentration profile indicates that Co exhibits a\npreferential association with Y, along with Fe and O, while displaying a relative exclusion toward Zr.\nThese RDF patterns provide critical insights into nanoscale compositional uniformity and segregation\ntendencies, with BCZTZICM displaying superior homogeneity compared to the compositional\nclustering observed in BSCF and BCFZY.\nWe sincerely appreciate your valuable comments regarding the XPS characterization. In situ near-\nambient pressure XPS (NAP-XPS) measurements were performed on the samples under revised\nexperimental protocols, as delineated in Fig S. Following stabilization at 600 °C under 0.3 mbar air\natmosphere for one hour, then samples were cooled to ambient temperature, with spectral acquisitions\nconducted at both 600 °C and room temperature for comparative analysis. Upon heating to 600 °C,\nboth BSCF and BCFZY underwent transition from bimodal to single-peak configurations, which\npersisted upon subsequent cooling to room temperature. Conversely, BCZTZICM maintained\ncharacteristic bimodal features corresponding to adsorbed and lattice oxygen species throughout the\nentire thermal cycle. Analysis of the peak position shifts further indicates that BCZTZICM exhibits\n24\nsuperior stability of both lattice oxygen and surface-adsorbed oxygen species under the investigated\nconditions.\nWe sincerely appreciate the reviewer's insightful comment regarding the computational\nmethodology and the challenges associated with modeling multi-doped perovskite systems. We fully\nacknowledge that exhaustively sampling all possible atomic configurations in a multi-component\nsystem is computationally prohibitive, and we agree that our supercell size cannot capture every\npossible dopant arrangement. However, we would like to clarify our approach and provide additional\njustification for the computational results presented in our manuscript. While we recognize that our\nsupercell size (11.8720 Å×11.8720 Å×25.8934 Å) cannot capture all possible dopant configurations,\nour computational approach was deliberately designed to provide meaningful qualitative trends rather\nthan absolute quantitative predictions. We carefully constructed representative atomic configurations\nthat reflect the experimentally observed homogeneous dopant distribution, as confirmed by our APT\nresults shown in Fig. 1g and Supplementary Fig. 5-7. For each composition, we tested multiple initial\nconfigurations (3-5 different arrangements) and selected the most energetically favorable structures\nfor comparison, finding that the reported trends were consistent across different configurations.\nImportantly, our DFT calculations were primarily used for comparative studies between BCZTZICM,\nBSCF, and BCFZY under identical computational conditions, which minimizes systematic errors and\nmakes the relative trends more reliable than absolute values. We emphasize that our computational\npredictions are strongly validated by extensive experimental characterization across multiple\nindependent techniques. The DFT-predicted lower proton diffusion barriers for BCZTZICM (Fig. 1h)\nare directly confirmed by electrochemical impedance spectroscopy, which demonstrates that\nBCZTZICM achieves the lowest area-specific resistance of 0.163 Ω cm2 at 600°C compared to BCFZY\n(0.337 Ω cm2) and BSCF (0.311 Ω cm2) under humidified conditions (Fig. 4d). The calculated H O\n2\nadsorption and hydration energies (Fig. 4c) align remarkably well with experimental FT-IR\nspectroscopy (Fig. 4b) and XANES results (Fig. 3c-d), which independently confirm the enhanced\nhydration capability of BCZTZICM. Furthermore, the predicted uniform oxygen vacancy distribution\nand reduced vacancy formation energy variance (Fig. 2e) are corroborated by minimal thermal\nexpansion anomalies (TEC=17.67×10-6 K-1 shown in Fig. 3b), superior thermal cycling stability (Fig.\n3e-g), and lower O desorption in O -TPD measurements (Supplementary Fig. 13). Most critically, the\n2 2\nwhich formed the foundation for constructing our computational models. Most importantly, the\nultimate validation of our computational approach comes from the exceptional experimental\nperformance achieved by the BCZTZICM-based devices. Our cells demonstrate a peak power density\nof 1.56 W cm-2 at 600 °C, an electrolysis current density of 2.0 A cm-2 at 1.3V, and outstanding stability\nover 780 hours with degradation rates of only 19.3 and 16.9 μV h-1 in fuel cell and electrolysis modes,\nrespectively. These results represent state-of-the-art performance for protonic ceramic electrochemical\ncells and strongly support the validity of the design principles derived from our computational insights.\nThe fact that materials designed based on these computational predictions achieve such exceptional\nperformance provides compelling evidence that our approach, despite its inherent limitations in\nconfiguration sampling, successfully captures the essential physical and chemical trends governing\nmaterial behavior.\nAction: We have added Fig. S and revised the manuscript and supplementary information as below:\nPage 6, Line 125-137\nTo gain deeper insights into the local chemical environment and potential nanoscale segregation\ntendencies, radial distribution function (RDF) analysis derived from APT was performed. As illustrated\nin Supplementary Fig. 8, with Co ions as reference center, BCZTZICM exhibits no pronounced\ntendency toward deep association with any single proximate ionic species. The normalized\nconcentrations of high-abundance Co cations and oxygen anions approach unity, while trace-doped In\nand Zr ions display oscillatory concentration profiles within the analyzed volume (constrained by\nlimited abundance), collectively indicating a remarkably homogeneous spatial distribution of Co ions,\noxygen ions, and dopant species. In contrast, the RDF of BSCF suggests a slight tendency for Co ions\nto disperse from Fe, implying a preferential association among like ions (Co–Co and Fe–Fe) within\nthe analyzed volume. Furthermore, oxygen anions demonstrate a stronger association with Co than\nwith Fe. For BCFZY, the gradually rising normalized concentration profile indicates that Co exhibits\na preferential association with Y, along with Fe and O, while displaying a relative exclusion toward Zr.\n[Supplementary information]\nRadial distribution function (RDF) analysis of neighboring ions around Co centers within a 10 nm\nrange, derived from APT data.\nPage 9-10, Line 182-193\nIn-situ near-ambient pressure X-ray photoelectron spectroscopy (NAP-XPS) measurements\nwere conducted to probe the surface oxygen chemistry. As shown in Supplementary Fig. 14, upon\nstabilization at 600 °C, both BSCF and BCFZY underwent a transition from a double-peak to a single-\npeak configuration in the O 1s spectra, which persisted upon cooling to room temperature, indicating\nenhanced oxygen vacancy formation at elevated temperatures, accompanied by desorption of adsorbed\nspecies and rapid interconversion between lattice and adsorbed oxygen. In contrast, BCZTZICM\nretained its characteristic double-peak structure throughout the thermal cycle, demonstrating robust\noxygen adsorption capability even under high-temperature, low-P(O ) conditions. Comparative\n2\nanalysis of the 600 °C and room-temperature spectra (Supplementary Fig. 15) reveals that BCZTZICM\nexhibits strengthened metal-oxygen bonding and sustained adsorption capacity, attributable to the\nextensive distribution of disordered oxygen vacancies that provide abundant active sites for surface\n27\nexchange.\n[Supplementary information]\nSupplementary Fig. 14.\nIn-situ near-ambient pressure X-ray photoelectron spectroscopy (NAP-XPS) O 1s data of BCZTZICM,\nBCFZY, and BSCF.\nIn-situ NAP-XPS O1s spectra of BSCF, BCFZY, and BCZTZICM measured in 0.3 mbar air at (a)\n600 °C and (b) after cooling to room temperature.\nThe authors show no appreciable chemical expansion (I suggest that the authors read this paper on\nchemical expansion: Phys. Chem. Chem. Phys., 2015, 17, 10028-10039) for the BCZTZICM sample.\nThey claim that this is due to the hydration process. However, hydration process is known to be charge\nneutral (see this in any classic papers on proton conductors), i.e., no B-site cation redox involved.\nIt is very strange to see that Co XANES show a peak shift upon switching to wet atmosphere (the peak\nshift is very small though, maybe due to formation of a secondary phase in wet atmosphere).\nResponse to C2:\nThanks for reviewer’s valuable insight and comments. We appreciate it for pointing this out. The\ncontradiction between the charge neutrality principle in the classical hydration process and our\nXANES observation results you pointed out is really the core problem that we need to face and explain\nseriously. The reviewer’s opinion has greatly helped us to improve the scientific rigor of the paper.\nThe reviewer's speculation that the small XANES peak shift might originate from secondary\nphase formation under wet atmosphere is a reasonable hypothesis. We have carefully verified this\npossibility through high temperature X-ray diffraction (XRD) analysis, which shows no evidence of\nany new phase formation. Therefore, we can conclude that the observed XANES changes are intrinsic\nresponses of the BCZTZICM perovskite phase to the steam atmosphere.\nWe fully agree with the reviewer's core argument. The classical hydration reaction for proton\nincorporation:\n𝐻 𝑂+𝑂× +𝑉∙∙ = 2𝑂𝐻∙\n2 𝑂 𝑜 𝑂\nThis process is indeed charge-neutral and, in theory, should not involve a change in the valence\nstate of the B-site cations. However, our reproducible XANES data clearly show a small but consistent\nshift of the Co K-edge to higher energy upon introducing steam at 600 °C. This shift typically indicates\na decrease in local electron density around Co, suggesting a trend towards a higher apparent oxidation\nstate.\nThis observation suggests that in the complex environment of a PCEC oxygen electrode, proton\nincorporation is not solely governed by the classical, charge-neutral hydration mechanism. Beyond\nthat, proton absorption can proceed through alternative reaction pathways that do involve transition\nmetal redox chemistry. Specifically, the following two mechanisms are well-established in the\n30\nliterature for oxygen electrodes:\n2𝑂× +𝐻 𝑂+2ℎ· = 2𝑂𝐻∙ +1/2𝑂\n𝑂 2 𝑂 2\n𝑀𝑛(Ⅲ)𝑥 +𝑂· +𝐻 𝑂+𝑉∙∙ = 𝑀𝑛(Ⅳ)· +2𝑂𝐻∙\n𝑀𝑛 𝑂 2 𝑜 𝑀𝑛 𝑂\nThe third mechanism explicitly involves a change in the valence of the redox-active transition\nmetal (e.g., Co, Mn, Fe, Ni). This pathway has been repeatedly demonstrated to contribute significantly\nto proton uptake in perovskite oxygen electrodes (Chem. Mater. 2019, 31, 8383-8393; Energy Environ.\nMater. 2023, e12660; Nat. Catal. 2022, 5, 777-787). All of these mechanisms aim to form proton\nvacancies, thereby improving the proton conductivity of the electrode.\nWe acknowledge that using XANES alone to discuss this process is somewhat limited. In addition\nto the evidence mentioned above, we also characterized the material’s physical adsorption capacity\nusing H O-TPD and observed proton defect-related characteristic peaks via in-situ Fourier transform\n2\ninfrared spectroscopy. We believe that the combination of these multiple pieces of evidence supports\nthe conclusion that BCZTZICM exhibits improved proton absorption capabilities.\nWe sincerely thank the reviewer for pointing this issue, which has significantly strengthened the\nscientific rigor of our manuscript.\nAction: We have revised the manuscript and supplementary information as below:\nPage 13, Line 250-253\nThe steam tolerance of BCZTZICM was evaluated by high-temperature XRD during cyclic transitions\nbetween dry and humid air atmospheres. As shown in Supplementary Fig. 23, the material exhibits\nonly reversible lattice expansion/contraction, without emergence of secondary phases.\nPage 2, Line 26\nMoreover, the thermally driven mild oxygen release can be further offset by beneficial proton uptake,\nthereby increasing the thermomechanical durability of the oxygen electrode.\nPage 2, Line 26\nFurthermore, X-ray absorption near-edge structure (XANES) and in-situ Fourier transform infrared\n(FT-IR) spectroscopy confirms an enhanced proton uptake capability, which synergistically mitigates\nthe microstructural degradation induced by electrode cracking.\nPage 12, Line 233-236\nSteam-treated sample shows elevated valences, indicating proton uptake partially offsets thermal\nexpansion. This process (Supplementary Note 2) generates protonic defects that induce the changes in\n31\nlocal electronic structure and ion valence. OH- formation enables partial assessment of the proton\nuptake capacity via valence variations. After normalization for the elemental composition, lattice\noxygen content variations quantify proton uptake capacity (Fig. 3d). BCZTZICM exhibits an\nexceptional proton uptake capability under humid conditions, with a lattice oxygen content increase of\n0.056, compared with only 0.042 for BSCF, when switching from dry to humidified environments at\n600 °C.\n[Supplementary information]\nSupplementary Fig. 23.\nHigh temperature XRD patterns of BCZTZICM at 600 °C during switching between dry and humid\nair atmospheres, and magnified XRD patterns of the selected 2 theta range of 29.55-30.45°.\nThe proton uptake has the beneficial effect of enhancing the absorption of protons by the electrode,\nthereby strengthening the proton transport capacity. In general, proton defects can be generated by\nhydration reactions with oxygen vacancies, as illustrated by the following formula:\n𝐻 𝑂+𝑂× +𝑉∙∙ = 2𝑂𝐻∙\n2 𝑂 𝑜 𝑂\nWhere 𝑂× represents lattice oxygen, 𝑉∙∙ is oxygen vacancies, and 𝑂𝐻∙ represents proton defects.\n𝑂 𝑜 𝑂\nBeyond classical hydration pathways, PCEC oxygen electrodes can generate protonic defects through\ntransition metal redox-mediated routes.\n2𝑂× +𝐻 𝑂+2ℎ· = 2𝑂𝐻∙ +1/2𝑂\n𝑂 2 𝑂 2\n𝑀𝑛(Ⅲ)𝑥 +𝑂· +𝐻 𝑂+𝑉∙∙ = 𝑀𝑛(Ⅳ)· +2𝑂𝐻∙\n𝑀𝑛 𝑂 2 𝑜 𝑀𝑛 𝑂\nWhere ℎ· represents electronic holes, 𝑀𝑛(Ⅲ)𝑥 is Mn3+ ion. The proposed mechanisms,\n𝑀𝑛\nextendable to Fe/Co/Ni systems based on their redox activity, have been experimentally validated in\nrelevant perovskite electrodes.\nComment 3:\nIt is very difficult to understand the key contributors to the increased performance of the BCZTZICM\nsample without knowing some fundamental properties of the perovskite (i.e., oxygen non-\nstoichiometry, ionic/electronic conductivity). It is entirely possible that the better performance\ncompared with BSCF is NOT from “ionic dispersion”. The authors might want to reflect how to prove\nthis causality.\nResponse to C3:\nWe sincerely thank the reviewer for raising this critical point regarding the causality behind the\nenhanced performance of BCZTZICM. We agree that a deeper understanding of the fundamental\nproperties is essential to substantiate our claim that the \"ionic dispersion\" strategy enhances\nperformance by stabilizing the proton conduction pathway, rather than boosting oxygen exchange\n33\nkinetics. To address this, we have conducted a series of complementary experiments, the results of\nwhich are now included in the revised manuscript (Fig. 2f, Supplementary Fig. 25-26).\nBased on TG coupled with room-temperature X-ray absorption near-edge structure (XANES)\nspectroscopy to determine elemental valence states, we have determined the temperature-dependent\noxygen non-stoichiometry ratio (δ) for all samples (Fig. 2f). The results reveal distinct oxygen\ndesorption behaviors: BSCF exhibits rapid oxygen loss at elevated temperatures, generating excessive\noxygen vacancies. BCZTZICM and BCFZY demonstrate significantly milder oxygen release profiles.\nElectronic conductivity measurements were performed and are presented in Fig. S. BSCF shows\na characteristic non-monotonic temperature dependence, directly correlated with its severe oxygen loss.\nThis is due to excessive oxygen vacancies annihilating electron holes, leading to decreased carrier\nconcentration and consequently reduced electronic conductivity at high temperatures. In contrast,\nBCZTZICM and BCFZY demonstrate significantly milder oxygen release profiles. Notably,\nBCZTZICM maintains more stable oxygen content at operating temperatures, with near-linear\ntemperature-dependent conductivity behavior indicative of thermally activated electron hopping as the\ndominant charge transport mechanism, supporting the suppression of excessive oxygen vacancy\nformation. BCZTZICM maintains stable oxygen content at elevated temperatures, preserving the\nintegrity of proton diffusion pathways that depend on lattice oxygen. In contrast, BSCF undergoes\nexcessive lattice oxygen release, and the resulting vacancies disrupt proton diffusion pathways,\ncreating high-barrier diffusion sites.\nTo evaluate oxygen electrocatalytic activity, electrical conductivity relaxation (ECR)\nmeasurements were conducted to extract oxygen surface exchange coefficients (k ) and bulk\nchem\ndiffusion coefficients (D ) (Fig. S). Importantly, BCZTZICM does not show significantly superior\nchem\noxygen activity compared to BSCF under dry air conditions. BCZTZICM does not exhibit notably\nsuperior oxygen activity compared to BSCF. Symmetric cell measurements using GDC electrolytes\nfurther reveal that BCZTZICM shows slightly lower ORR activity than BSCF and BCFZY under dry\nair conditions. Therefore, the superior performance of BCZTZICM in PCECs is not due to enhanced\noxygen reduction kinetics but is primarily attributed to enhanced proton uptake and optimized proton\nconduction capability, facilitated by a stabilized proton-conducting lattice. We thank the reviewer\n34\nagain for prompting this more rigorous and nuanced discussion, which has significantly strengthened\nour manuscript.\nAction: We have revised the manuscript and supplementary information as below:\nPage 9, Line 177-180\nThermogravimetric analysis demonstrates a mass loss of 1.12% for BCZTZICM and higher losses for\nBSCF and BCFZY. Based on room-temperature XANES, the oxygen non-stoichiometric ratios were\nobtained. BCZTZICM demonstrates a higher concentration of inherent oxygen vacancies compared to\nBCFZY, which is beneficial to surface activity.\nFig. 2 Multielement micro-doping enhances structural homogenization. (a) Bond length distributions, (b)\nminima of the Electron localization function (ELF) along bond paths, and (c) O-ELF values for BCZTZICM,\nBSCF, and BCFZY, using standard deviation (σ) as the metric for internal uniformity. (d) Average O-ELF values,\n(e) oxygen vacancy formation energy, and (f) TG and oxygen non-stoichiometric ratio data of BCZTZICM BSCF\nand\nBCFZY samples.\nPage 15, Line 304-314\nElectronic conductivity measurements (Supplementary Fig. 25) reveal the characteristic of BSCF\nnon-monotonic temperature dependence, directly correlated with severe oxygen loss. Excessive\noxygen vacancies annihilate electron holes, reducing carrier concentration and conductivity at high\n35\ntemperatures. In contrast, BCZTZICM maintains stable oxygen content at operating temperatures,\nexhibiting near-linear temperature-dependent conductivity dominated by thermally activated electron\nhopping. This behavior confirms suppressed oxygen vacancy formation, preserving proton diffusion\npathways that depend on lattice oxygen. BSCF's excessive oxygen release disrupts these pathways,\ncreating high-barrier diffusion sites.\nElectrical conductivity relaxation (ECR) measurements were conducted to extract oxygen surface\nexchange coefficients (k ) and bulk diffusion coefficients (D ) (Supplementary Fig. 26). Notably,\nchem chem\nBCZTZICM shows no superior oxygen activity versus BSCF under dry air. Electrochemical\nimpedance spectroscopy (EIS) of symmetric cells determines of electrode-specific contributions.\nDry-\nair testing of Gd Ce O (GDC) based cells eliminated the confounding effects of proton conduction.\n0.2 0.8 1.9\nSupplementary information\nSupplementary Fig. 25.\nElectronic conductivities of the BCZTZICM, BSCF, and BCFZY samples at 300–800 °C in air, as\nmeasured by the 4-probe DC conductivity technique.\nThe fitted values of D and k of BCZTZICM, BSCF, and BCFZY samples at temperature range\nchem chem\nof 550–700 °C from electronic conductivity relaxation curves.\n37"}, {"round": 2, "reviewer_comments": {"reviewer_2": "All of my questions have been fully resolved in the P2P response, and the submission is suitable to be accepted.", "reviewer_3": "The authors have revised the manuscript and I have been asked by the editor to review whether my comments have been fully addressed. Unfortunately, I do not believe that the questions I raised have been fully answered. I still fail to see that the concept of “micro-doping”, which is claimed to induce a more “homogenous” distribution of oxygen vacancies, can really pan out. Still, the link between “micro-doping” to more homogenous oxygen vacancy distribution to better electrochemical performance cathode is still rather weak, in my opinion. Please note that I am only commenting on whether the authors have addressed my previous comments. I have no intention of further commenting on the quality of this work or introducing new questions.\n1. Regarding the more homogenous distribution of oxygen vacancies claimed by the authors:\nThe authors claimed that they reached this conclusion not by directly observing the distribution of oxygen vacancies in microscopic scales (due to technical difficulties). Rather, the conclusion is reached circumventively by looking at the distribution of cations. Firstly, I cannot see why the RDF analysis from the APT results show a more homogenous distribution of cations for BCZTZICM (in fact, the RDF of this sample show a higher fluctuation, indicating potential inhomogeneity). Secondly, even if the cations are randomly distributed, adding more dopant might introduce more electrostatic interactions between particular dopant and oxygen vacancies.\nThe AP-XPS data is not very meaningful, in the sense that the Oads peak has multiple origins, ranging from surface -OH group to contaminates such as adventitious hydrocarbons to sulfuric groups. Since the BCZTZICM shows a persistent Oads peak even at 600 degree C, I suspect that most likely this peak is due to surface contamination of S/Si species (see for example, Riedl et al., ACS Appl. Mater. Interfaces 2023, 15 (22), 26787–26798. https://doi.org/10.1021/acsami.3c03952). In any case, this O 1s peak CANNOT be linked to oxygen vacancies (see Wang et al., Journal of the European Ceramic Society 2024, 44 (15), 116709. https://doi.org/10.1016/j.jeurceramsoc.2024.116709.).\n2. Regarding the change of XANES energy shift upon hydration:\nThe authors explained the shift of XANES spectra by writing a “new” defect chemical reaction. However, I should note that the new defect chemical reaction is actually the hydrogenation process (i.e., splitting of H2O into H2 and O2, where H2 reacts with perovskite to form reduced cations and protons). It is by no means hydration process. Unless the authors believes that their perovskite can induce spontaneous water splitting (and magically get rid of the oxygen gas molecules, too), this process cannot happen under water steam. I suggest that the authors review fundamental defect chemistry on proton conductors by consulting classic literature contributions (e.g., Merkle et al., Annu. Rev. Mater. Res. 2021. 51:461–93).\n3. Regarding the link between micro-doping and enhanced performance:\nBased on the reasons above, as well as comments I noted in my previous review opinion, I still fail to see why the enhanced electrochemical performance HAS TO be linked with the concept of micro-doping and homogenous distribution of oxygen vacancies. In my opinion, the authors can provide clear information to (rather than mislead) the community by simply reporting that they have found a new cathode material for SOCs that can out-perform BSCF, as well as reporting the basic properties (lattices, ionic/electronic conductivities, hydration behaviors, etc.), rather than making a complicate story as shown in the current form of the manuscript.", "reviewer_4": "This manuscript reports a thoughtfully designed multielement micro-doping strategy to engineer a highly ionic-dispersed BaCoO3-based oxygen electrode for reversible proton ceramic electrochemical cells (R-PCECs). The authors convincingly demonstrate that atomic-scale compositional homogenization can simultaneously reduce proton diffusion barriers and improve thermomechanical stability under humid oxidizing conditions. Overall, the study is technically sound, well supported by experimental and computational evidence, and of clear interest to the solid-state electrochemistry and energy-materials communities. I am in favor of publication after revision, provided that the authors address the following points to further improve clarity and accessibility of the manuscript:\n1. Please clarify how the proposed multielement micro-doping strategy is fundamentally different from conventional high-entropy or random B-site doping in perovskites.\n2. Please explicitly state the core design principle (ionic dispersion vs. entropy vs. strain vs. valence diversity) responsible for the observed performance gains.\n3. Suggest provide a quantitative definition of the proton-conduction bottleneck in oxygen electrodes, with representative literature values.\n4. Suggest to clearly distinguish bulk proton transport from surface or interfacial proton transport in the discussion.\n5. Suggest to discuss whether multiple proton transport mechanisms may coexist under different temperatures or humidity conditions.\n6, Clarify how proton uptake affects oxygen vacancy concentration and lattice stability.\n7. Please clearly label all APT elemental maps and color scales and discuss dataset representativeness.\n8. Please provide key DFT modeling details (supercell size, vacancy configuration, proton concentration).\n9. Clarify whether reduced polarization resistance is limited by proton transport, oxygen exchange, or coupled kinetics.\n10, is that possible to quantify and compare crack width/density among electrodes to support mechanical stability claims.\n11. PleaSe explain whether improved thermomechanical stability arises mainly from suppressed oxygen release or proton uptake.\n12. Please discuss dominant long-term degradation mechanisms observed or anticipated beyond 780 h.\n13, Please strengthen comparison with state-of-the-art oxygen electrodes using normalized metrics where possible.\n14, I am curious whether similar performance could be achieved with fewer dopant elements."}, "author_response": "Point-to-Point Responses to Reviewers’ Remarks and Suggestions\nFirst, we express our sincere gratitude to editor for handling our manuscript and for the insightful\nfeedback provided by the reviewers. Their comments have greatly improved the quality and clarity of\nthe paper.\nBelow, the following are the point-by-point response to reviewers’ comments. All changes in the\nmanuscript have been marked in red for convenience.\nReviewer #3 (Remarks to the Author):\nThe authors have revised the manuscript and I have been asked by the editor to review whether my\ncomments have been fully addressed. Unfortunately, I do not believe that the questions I raised have\nbeen fully answered. I still fail to see that the concept of “micro-doping”, which is claimed to induce\na more “homogenous” distribution of oxygen vacancies, can really pan out. Still, the link between\n“micro-doping” to more homogenous oxygen vacancy distribution to better electrochemical\nperformance cathode is still rather weak, in my opinion. Please note that I am only commenting on\nwhether the authors have addressed my previous comments. I have no intention of further commenting\non the quality of this work or introducing new questions.\nComment 1:\nRegarding the more homogenous distribution of oxygen vacancies claimed by the authors:\nThe authors claimed that they reached this conclusion not by directly observing the distribution of\noxygen vacancies in microscopic scales (due to technical difficulties). Rather, the conclusion is\nreached circumventively by looking at the distribution of cations. Firstly, I cannot see why the RDF\nanalysis from the APT results show a more homogenous distribution of cations for BCZTZICM (in\nfact, the RDF of this sample show a higher fluctuation, indicating potential inhomogeneity). Secondly,\neven if the cations are randomly distributed, adding more dopant might introduce more electrostatic\ninteractions between particular dopant and oxygen vacancies. The AP-XPS data is not very meaningful,\nin the sense that the Oads peak has multiple origins, ranging from surface -OH group to contaminates\nsuch as adventitious hydrocarbons to sulfuric groups. Since the BCZTZICM shows a persistent Oads\npeak even at 600 degree C, I suspect that most likely this peak is due to surface contamination of S/Si\n1\nspecies (see for example, Riedl et al., ACS Appl. Mater. Interfaces 2023, 15 (22), 26787–26798.\nhttps://doi.org/10.1021/acsami.3c03952). In any case, this O 1s peak CANNOT be linked to oxygen\nvacancies (see Wang et al., Journal of the European Ceramic Society 2024, 44 (15), 116709.\nhttps://doi.org/10.1016/j.jeurceramsoc.2024.116709.).\nResponse to C1:\nWe sincerely appreciate the Reviewer's insightful comments regarding the RDF analysis and\nNAP-XPS interpretation.\nRegarding the RDF analysis from APT, the observed fluctuations in normalized concentration for\nBCZTZICM (Supplementary Fig. 8) are expected. The key metric is the proximity of the normalized\nconcentration profiles to unity across various radial distances, which indicates similar association or\nsegregation tendency between different ion species. This contrasts with BSCF and BCFZY, whose\nRDF show clearer tendencies for like-ion association (e.g., Co-Co in BSCF) or exclusion (e.g., Co-Zr\nin BCFZY). The fluctuations in RDF of BCZTZICM thus reflect a more randomized, dispersed\ncationic distribution without long-range segregation, which is the intended outcome of our multi-\nelement micro doping strategy.\nWe fully concur with the Reviewer's consideration that multi-dopants may introduce enhanced\nelectrostatic interactions. Indeed, such strengthened electrostatic interactions would suppress proton-\nelectron coupled migration. The increased average oxygen vacancy formation energy and the notably\nlower TEC for BCZTZICM are proof of the modified and homogenized chemical bonding\nenvironment, which includes these electrostatic interactions. This aligns with established literature\nwhere high-entropy or multi-doping strategies suppress lattice oxygen loss and reduce TEC.\nCritically, our proton diffusion barrier calculations reveal that BCZTZICM does not possess the\nlowest individual hopping barriers. Specifically, the minimum proton hopping barrier between\nadjacent oxygen sites in BSCF can be as low as 0.20 eV, compared to 0.24 eV for BCFZY and 0.25\neV for BCZTZICM. However, the maximum barrier sites in BSCF reaches 0.87 eV, significantly\nexceeding those of BCFZY (0.51 eV) and BCZTZICM (0.41 eV). In proton conduction processes,\n2\nhigh-barrier sites effectively obstruct potential migration pathways, creating bottlenecks that limit\noverall proton conductivity regardless of the presence of low-barrier regions.\nThe strategy of suppressing electrostatic interactions to enhance proton conductivity has been\nextensively studied in classical proton conductor materials, and we fully agree with the Reviewer's\nperspective. However, for PCEC oxygen electrodes, several unique considerations apply: electrode\nmaterials must maintain substantial ORR/OER activity, necessitating the incorporation of redox-active\ntransition metals (Co, Fe, etc.) as catalytic centers. This requirement inevitably introduces abundant\noxygen vacancies and severe local valence and coordination heterogeneity, resulting in the coexistence\nof numerous low-barrier and high-barrier regions. We believe this one of the primary bottlenecks\ncurrently limiting the development of high-performance oxygen electrode materials.\nRegarding the NAP-XPS measurements: We appreciate the Reviewer's concern about potential\nsurface contamination. We wish to clarify that all samples were prepared following identical protocols\nwithin an argon-atmosphere glovebox, and during the in-situ heating process, high-purity bottled air\n(rather than ambient atmosphere) was continuously introduced, thereby minimizing the possibility of\nadventitious contamination.\nFurthermore, evidence supporting enhanced metal-oxygen bonding in BCZTZICM is also\nderived from TGA and O -TPD.\nremoved part about the NAP-XPS O 1s spectra.\nAction: We have revised manuscript as below.\nPage 7, Line 143-153:\nTo gain deeper insights into the local chemical environment and potential nanoscale segregation\ntendencies, radial distribution function (RDF) analysis derived from APT was performed25, 26. The key\nmetric is the proximity of the normalized concentration profiles to unity across various radial distances,\nwhich indicates similar association tendency between different ion species. As illustrated in\nSupplementary Fig. 8, with Co ions as the center, the normalized concentrations of high-abundance\nCo cations and oxygen anions approach unity, while trace-doped In3+ and Zr4+ ions display vibrant\n3\nconcentration profiles within the analyzed region (constrained by limited abundance), collectively\nindicating a remarkably homogeneous spatial distribution of Co ions, O2-, and dopant ions. In contrast,\nBSCF suggests a slight tendency for Co ions to disperse from Fe, implying a preferential association\namong like ions (Co-Co and Fe-Fe) within the analyzed region, and O2- demonstrate a stronger\nassociation with Co than with Fe. For BCFZY, Co exhibits a preferential association with Y3+, Fe ions\nand O2-, while displaying a relative exclusion toward Zr4+.\nPage 8-9, Line 166-171:\nWhile BSCF contains low-barrier sites (0.20 eV), high-barrier regions (0.87 eV) obstruct\nmigration, limiting conductive pathways and the overall proton conductivity. Notably, BCZTZICM\nmaintains exceptional uniformity, with only 0.16 eV difference between maximum (0.41 eV) and\nminimum (0.25 eV) barrier sites, eliminating bottleneck effects and creating extensive networks of\nfacile transport pathways that enable rapid proton diffusion in three dimensions.\nPage 10, Line 193-195:\nThermogravimetric analysis demonstrates a mass loss of 1.12% for BCZTZICM and higher losses\nfor BSCF and BCFZY, indicating an enhanced metal-oxygen bond energy of BCZTZICM.\nComment 2:\n2. Regarding the change of XANES energy shift upon hydration:\nThe authors explained the shift of XANES spectra by writing a “new” defect chemical reaction.\nHowever, I should note that the new defect chemical reaction is actually the hydrogenation process\n(i.e., splitting of H O into H and O , where H reacts with perovskite to form reduced cations and\n2 2 2 2\nprotons). It is by no means hydration process. Unless the authors believes that their perovskite can\ninduce spontaneous water splitting (and magically get rid of the oxygen gas molecules, too), this\nprocess cannot happen under water steam. I suggest that the authors review fundamental defect\nchemistry on proton conductors by consulting classic literature contributions (e.g., Merkle et al., Annu.\nRev. Mater. Res. 2021. 51:461–93).\nResponse to C2:\nWe thank the reviewer for the insightful comment and for pointing us to the fundamental\nreference. The Reviewer correctly identifies the hydrogenation mechanism, wherein water dissociates\ninto H and O , with H subsequently reducing the perovskite to form reduced cations and protonic\n2 2 2\ndefects. This process indeed consumes electron holes, leading to decreased cation valence states and,\nin materials with weak metal-oxygen bonding, potential cation exsolution (as recently demonstrated\nin ACS Energy Lett. 2025, 10, 4948−4956 by Woochul et al.).\nHowever, the mechanism discussed in this manuscript is NOT hydrogenation. The experimental\nresults show that the valence of transition metal cations INCREASES after introducing water vapor，\nrather than decrease of valence. This is derived from lowering of antibonding O 2p states that hybridize\nwith transition metal 3d orbitals, coupled with the influence of incorporated protonic defects on the\nlocal coordination environment of transition metal ions.\nThis mechanism has been experimentally validated through operando X-ray absorption\nspectroscopy studies by Aoki and Zhou, respectively. (Chem. Mater. 2019, 31, 8383−8393, Energy\nEnviron. Mater. 2023, 0, e12660) This demonstrated that proton incorporation can induce oxidation of\nB-site cations in specific perovskite systems. Similarly, in steam electrolysis via solid oxide\nelectrolysis cells, water vapor can also induce re-oxidation of the Ni-YSZ hydrogen electrode to NiO-\nYSZ. Therefore, hydrogen is typically employed as a carrier gas to mitigate the re-oxidation of nickel.\nIt is important to note that this oxidative hydration mechanism is not universally observed across\nall perovskite compositions. For instance, Aoki et al. subsequently reported that La Ca Co Ni O\n0.8 0.2 x 1-x 3-\ndoes not exhibit this behavior (ACS Appl. Energy Mater. 2021, 4, 554−563). It depends on the\nδ\nelectronic structure characteristics of the specific perovskite system, particularly the degree of metal-\noxygen covalency, the position of the Fermi level relative to the oxygen 2p band, and the local\ncoordination flexibility around B-site cations.\nWe have revised the manuscript to clarify that the observed valence increase is attributed to this\nproton-incorporation-induced electronic modulation, and we have removed the previous simplified\ndefect reaction to avoid confusion.\nAction: We have revised manuscripts and Supporting Information as below.\nPage 11, line 227-234:\nUnder dry air conditions, BCZTZICM demonstrates minimal lattice oxygen content change (-\n0.075) compared with substantial losses for BSCF (-0.247) during heating to 600 °C (Fig. 3d),\nindicating lower thermal expansion, which were calculated using the equation in Supplementary Note\n1. Notably, steam-treated samples exhibit elevated B-site cation valences. This oxidative response\nupon proton uptake originates from a distinct mechanism: the lowering of antibonding O 2p states\nhybridized with transition metal 3d orbitals, coupled with modifications to the local coordination\nenvironment induced by protonic defect incorporation (Supplementary Note 2). This mechanism has\nbeen experimentally validated through operando XAS studies on analogous perovskite systems,\ndepending on the specific electronic structure characteristics (Supplementary Note 2). After\nnormalization for elemental composition, the lattice oxygen content variations derived from XANES\nanalysis quantify the proton uptake capacity (Fig. 3d). BCZTZICM exhibits an exceptional proton\nuptake capability under humid conditions, with a lattice oxygen content increase of 0.056, compared\nwith only 0.042 for BSCF, when switching from dry to humidified environments at 600 °C.\nSupplementary Information, Page 47:\nBeyond classical hydration pathways, PCEC oxygen electrodes can generate protonic defects through\ntransition metal oxide mediated routes32, 33, 34.\n𝑀𝑛(Ⅲ)𝑥 +𝑂· +𝐻 𝑂+𝑉∙∙ = 𝑀𝑛(Ⅳ)· +2𝑂𝐻∙\n𝑀𝑛 𝑂 2 𝑜 𝑀𝑛 𝑂\nWhere ℎ· represents electronic holes, 𝑀𝑛(Ⅲ)𝑥 is Mn3+ ion. The proposed mechanisms,\n𝑀𝑛\nextendable to Fe/Co/Ni systems based on their redox activity. This mechanism is derived from\nlowering of antibonding O 2p states that hybridize with transition metal 3d orbitals, coupled with the\ninfluence of incorporated protonic defects on the local coordination environment of transition metal\nions. It also has been experimentally validated in relevant perovskite electrodes. It should be noted that\nthis oxidative proton uptake mechanism is not universally observed across all perovskite compositions,\ndepending on the degree of metal-oxygen covalency, the position of the Fermi level relative to the\noxygen 2p band, and the local coordination flexibility around B-site cations.\nComment 3:\n3. Regarding the link between micro-doping and enhanced performance:\nBased on the reasons above, as well as comments I noted in my previous review opinion, I still fail to\nsee why the enhanced electrochemical performance HAS TO be linked with the concept of micro-\ndoping and homogenous distribution of oxygen vacancies. In my opinion, the authors can provide clear\ninformation to (rather than mislead) the community by simply reporting that they have found a new\ncathode material for SOCs that can out-perform BSCF, as well as reporting the basic properties (lattices,\nionic/electronic conductivities, hydration behaviors, etc.), rather than making a complicate story as\nshown in the current form of the manuscript.\nResponse to C3:\nWe sincerely thank the reviewer for reiterating this fundamental question. We agree that\nestablishing a clear causal link is paramount. Our revisions and the data presented aim to demonstrate\nthat the enhanced performance is not merely a coincidence but a direct consequence of the atomic-\nscale homogenization achieved specifically through our multi-element micro-doping strategy. The\nmulti-element micro-doping strategy minimizes ionic association in BCZTZICM, as evidenced by\nAPT/RDF and FT-IR. This atomic-scale homogenization, corroborated by theoretical calculations,\nhelps eliminate high energy barrier bottlenecks for proton diffusion, and evidenced by DRT analysis\nof EIS. Consequently, while strengthened electrostatic interactions slightly reduce the intrinsic ORR\nactivity, they collectively enable enhanced proton conductivity, superior proton uptake, and optimized\nthermomechanical stability.\nRegarding the homogeneity of vacancy distribution, this was a speculation based on our current\nexperimental and theoretical results. This content has been removed from the manuscript. We have\nrevised the manuscript accordingly.\nPage 10, Line 205-207:\nIntegration of APT findings with theoretical calculations reveals that materials exhibiting severe\nionic localization (e.g., BSCF) readily form excessive thermal stress ultimately inducing electrode\ncracking (Fig. 3a).\nPage 10, line 215-216:\nMultielement B-site doping in BCZTZICM effectively suppresses lattice oxygen loss through\nuniform local ion coordination environment, yielding near-linear behaviour.\nPage 12, line 237-239:\nThe structural homogeneity ensures similar local atomic coordination, which leads to consistent\nlattice oxygen desorption across long-range domains.\nPage 13, line 272-278:\nBCZTZICM exhibits an exceptional proton uptake capability under humid conditions, with a\nlattice oxygen content increase of 0.056, compared with only 0.042 for BSCF, when switching from\ndry to humidified environments at 600 °C. The multi-element micro-doping, by maximizing ionic\ndispersion, creates this enhanced structural homogeneity, which in turn leads to a greater availability\nof uniform active sites, resulting in significantly improved proton uptake capability.\nPage 16, line 320-322:\nElectronic conductivity measurements (Supplementary Fig. 25) reveal the characteristic of BSCF\nnon-monotonic temperature dependence, directly correlated with severe oxygen loss. Excessive\noxygen vacancies annihilate electron holes, reducing carrier concentration and conductivity at high\ntemperatures. In contrast, BCZTZICM maintains stable oxygen content at operating temperatures,\nexhibiting near-linear temperature-dependent conductivity dominated by thermally activated electron\nhopping. This demonstrate that lattice oxygen desorption in BCZTZICM exhibits a stochastic\ndesorption characteristic. In contrast, the non-linear conductivity responses about temperature in BSCF\nand BCFZY are associated with the inhomogeneity of local structures.\nReviewer #4 (Remarks to the Author):\nThis manuscript reports a thoughtfully designed multielement micro-doping strategy to engineer a\nhighly ionic-dispersed BaCoO -based oxygen electrode for reversible proton ceramic electrochemical\n3\ncells (R-PCECs). The authors convincingly demonstrate that atomic-scale compositional\nhomogenization can simultaneously reduce proton diffusion barriers and improve thermomechanical\nstability under humid oxidizing conditions. Overall, the study is technically sound, well supported by\nexperimental and computational evidence, and of clear interest to the solid-state electrochemistry and\nenergy-materials communities. I am in favor of publication after revision, provided that the authors\naddress the following points to further improve clarity and accessibility of the manuscript:\nComment 1:\n1. Please clarify how the proposed multielement micro-doping strategy is fundamentally different from\nconventional high-entropy or random B-site doping in perovskites.\nResponse to C1:\nWe thank the reviewer for this important question.\nHigh-entropy doping typically aims to achieve four characteristic effects: (1) the high-entropy\nstabilization effect favoring single-phase formation, (2) the lattice distortion effect arising from size\nmismatch, (3) the sluggish diffusion effect, and (4) the cocktail effect from synergistic elemental\ninteractions. This typically dilutes the concentration of catalytically active B-site ions (e.g., Co, Fe),\nwhich can compromise electrochemical activity.\nRandom B-site doping conventionally involves the incorporation of a single dopant element at\nmoderate-to-high concentrations, which predominantly modifies the local coordination structure, often\ncreating heterogeneous local strain fields.\nOur multi-element micro-doping strategy is fundamentally distinct from conventional high-\nentropy or single-element doping in its design principle: instead of pursuing configurational entropy\nmaximization or relying on the effect of an individual dopant, we employ a low total doping\nconcentration (20% of B-sites) distributed across six distinct elements (Zr, Ti, Zn, In, Cu, Mo) at\nmicro-doping levels (~3.3 % atom each). This approach prioritizes maximized ionic dispersion over\nentropy, which minimizes ion aggregation, disrupts extensive vacancy or strain clusters, and creates a\nhomogeneous local bonding environment-while preserving a high bulk concentration (80%) of active\nCo ions to maintain electrocatalytic activity. The outcome is a synergistic effect where moderate lattice\nstrain enhances structural stability and atomic-scale homogenization facilitates uniform proton\ntransport.\nWe have revised the manuscript to explicitly articulate these fundamental distinctions.\nPage 5, lines 91-99:\nThis yielded the multi-element micro-doped BCZTZICM, which together with BSCF and\nBCFZY establishes a B-site elemental dispersion gradient design concept. Differs fundamentally from\nconventional high-entropy or single-dopant approaches. Instead of maximizing configurational\nentropy, multi-element micro-doping strategy distributes a low total doping concentration (20 % of\nB-sites) across six distinct elements (Zr, Ti, Zn, In, Cu, Mo) to prioritize ionic dispersion. This design\nsuppresses ion aggregation and disrupts defect clusters while preserving 80 % active Co for high\ncatalytic activity. The resulting homogeneous local environment, combined with moderate lattice\nstrain from valence and ionic-radius diversity, enables uniform proton-transport pathways with low\nenergy barriers and yields a low, near-linear thermal expansion. Ionic dispersion is the central\nmechanism, synergistically enhanced by valence diversity and controlled strain.\nComment 2:\n2. Please explicitly state the core design principle (ionic dispersion vs. entropy vs. strain vs. valence\ndiversity) responsible for the observed performance gains.\nResponse to C2:\nThe observed performance enhancement is principally driven by maximized ionic dispersion,\nachieved through multi-element micro-doping. This strategy intentionally utilizes valence diversity\n(Zn2+ to Mo6+) and a gradient of ionic radii (0.60-0.72 Å) to create a homogeneous, moderately strained\nlattice. While configurational entropy increases naturally, it is not the main mechanism. The core aim\nis to disrupt short-range ordering of ions and their associated defect clusters-especially Co/Fe\naggregation and aligned oxygen-vacancy networks. This atomic-scale homogenization directly enables\nuniform proton-diffusion pathways with consistently low energy barriers and no severe bottlenecks,\nand suppressed anisotropic oxygen release, yielding a low, near-linear thermal expansion.\nThus, ionic dispersion is the central principle, with valence diversity and controlled strain serving\nas synergistic tools to achieve it, while preserving 80 % B-site Co to maintain high catalytic activity.\nThe manuscript has been revised to clarify this design principle.\nWe have revised the manuscript as below:\nPage 5, lines 91-99:\nThis yielded the multi-element micro-doped BCZTZICM, which together with BSCF and\nBCFZY establishes a B-site elemental dispersion gradient design concept. Differs fundamentally from\nconventional high-entropy or single-dopant approaches. Instead of maximizing configurational\nentropy, multi-element micro-doping strategy distributes a low total doping concentration (20 % of\nB-sites) across six distinct elements (Zr, Ti, Zn, In, Cu, Mo) to prioritize ionic dispersion. This design\nsuppresses ion aggregation and disrupts defect clusters while preserving 80 % active Co for high\ncatalytic activity. The resulting homogeneous local environment, combined with moderate lattice\nstrain from valence and ionic-radius diversity, enables uniform proton-transport pathways with low\nenergy barriers and yields a low, near-linear thermal expansion. Ionic dispersion is the central\nmechanism, synergistically enhanced by valence diversity and controlled strain.\nComment 3:\n3. Suggest provide a quantitative definition of the proton-conduction bottleneck in oxygen electrodes,\nwith representative literature values.\nResponse to C3:\nWe are grateful to the reviewer for valuable comment regarding the quantitative description of\nthe proton conduction bottleneck.\nThe proton conduction bottleneck in oxygen electrodes represents a critical performance\nlimitation in protonic ceramic electrochemical cells (PCECs). State-of-the-art electrolytes such as\nBaZr Y O (BZY20) and BaZr Ce Y Yb O (BZCYYb) exhibit proton conductivity of\n0.8 0.2 3-δ 0.1 0.7 0.1 0.1 3-δ\napproximately 0.01-0.02 S cm-1 at 600 °C under humidified conditions (Enrico Traversa et al. Nature\nMater. 2010, 9, 846-852., Donglin Han et al. J. Mater. Chem. A, 2024, 12, 5875., J. Mater. Chem. A,\n2018, 6, 18571). In contrast, conventional oxygen electrodes demonstrate significantly lower proton\nconductivity values. The proton conductivity of conventional oxygen electrode materials, as evidenced\nby measurements utilizing diverse techniques-such as the bulk diffusion coefficient testing for\n(Ba/Sr)(Co/Fe/W)O @PrBa Sr Co Fe O (~10-5 S cm-1 at 800 °C; Ling Zhao et al. Adv. Mater.\n3-δ 0.5 0.5 1.5 0.5 5+δ\n2024, 2405052), hydrogen permeation for La Sr Sc Fe O (1.1×10-4 S cm-1 at 600 °C; Lei Bi et\n0.8 0.2 0.5 0.5 3-δ\nal. SusMat. 2023, 3, 697-708.), and the Hebb-Wagner DC polarization method for BaGd La Co O\n0.3 0.7 2 6-δ\n(4.4×10-5 S cm-1 at 600 °C; Tadeusz Miruszewski et al. J. Mater. Chem. A, 2024, 12, 13488)-is\nconsistently and significantly lower than that of state-of-the-art proton-conducting electrolytes.\nThis dramatic conductivity mismatch (2-4 orders of magnitude) creates a severe transport bottleneck,\nconfining electrochemical reactions to the narrow triple-phase boundaries (TPBs) where electrolyte,\nelectrode, and gas phase meet. Consequently, the effective reaction zone is restricted to approximately\n1-2 μm near the electrode-electrolyte interface, severely limiting electrode utilization and overall cell\nperformance. This quantitative analysis underscores the critical need for developing triple-conducting\n(H+/O2-/e-) oxygen electrodes to extend the electrochemically active region.\nWe have revised manuscript as below:\nPage 6, lines 111-116:\nProton conduction bottlenecks in oxygen electrodes critically limit PCEC performance. State-of-\nthe-art electrolytes (BaZr Y O , BZCYYb) exhibit proton conductivity of ~0.01-0.02 S cm-1 at\n0.8 0.2 3-δ\n600 °C under humidified conditions5, 19. The proton conductivity of conventional oxygen electrode\nmaterials, as evidenced by measurements utilizing diverse techniques-such as the bulk diffusion\ncoefficient testing20 for (Ba/Sr)(Co/Fe/W)O @PrBa Sr Co Fe O (~10-5 S cm-1 at 800 °C),\n3-δ 0.5 0.5 1.5 0.5 5+δ\nhydrogen permeation21 for La Sr Sc Fe O (1.1×10-4 S cm-1 at 600 °C), and the Hebb-Wagner\n0.8 0.2 0.5 0.5 3-δ\nDC polarization method22 for BaGd La Co O (4.4×10-5 S cm-1 at 600 °C) -is consistently and\n0.3 0.7 2 6-δ\nsignificantly lower than that of state-of-the-art proton-conducting electrolytes. This severe\nconductivity mismatch confines electrochemical reactions to narrow triple-phase boundaries,\nrestricting the effective reaction zone to approximately 1-2 μm near the electrode-electrolyte interface,\nlimiting electrode utilization and overall cell performance.\nComment 4:\n4. Suggest to clearly distinguish bulk proton transport from surface or interfacial proton transport in\nthe discussion.\nResponse to C4:\nWe sincerely thank the reviewer for raising this important distinction. Bulk proton conduction in\nperovskite-based oxygen electrodes is predominantly governed by the Grotthuss mechanism in both\nfuel cell and electrolysis operational modes, involving thermally activated proton hopping between\nadjacent lattice oxygen sites via sequential O-H bond cleavage and formation. In contrast, surface and\ninterfacial proton transport originates from the proton uptake process, wherein H O dissociated to form\n2\nprotonic defects (OH-). The source of protons differs fundamentally between the two operational\nmodes: in electrolysis mode, protons are derived exclusively from H O dissociation at the oxygen\n2\nelectrode surface; whereas in fuel cell mode, the water vapor generated as a reaction product at the\noxygen electrode undergoes subsequent proton uptake, thereby augmenting the protonic conductivity\nof the electrode. We have revised the manuscript to explicitly clarify these mechanism distinctions\nbetween bulk and surface proton transport processes.\nWe have revised manuscript as below:\nPage 6, lines 122-128:\nProton conduction in perovskite oxide mechanisms exhibit temperature-dependent dominance. At\ne\nl\ne\nv\na\nt\ne\nd\nPage 14, lines 280-289：\nProton transport in perovskite-based oxygen electrodes follows the Grotthuss mechanism in both\nt\noperational modes. Surface proton transport, in contrast, originates from the proton uptake process\ne\nthrough H O dissociative chemisorption to form protonic defects. Fundamental distinctions arise in\n2\nm\nproton sources between operational modes: in fuel cell mode, protons originate from hydrogen\np\ndissociation at the hydrogen electrode, transporting through electrolyte, and forming water at the\ne\noxygen electrodes, while water vapor generated as reaction product undergoes subsequent proton\nt\nuptake, augmenting electrode protonic conductivity. Electrolysis mode generates protons exclusively\nthrough water dissociation at the oxygen electrode, with subsequent transport to the hydrogen electrode.\nThe overall efficiency critically depends on electrode ability to rapidly adsorb water vapor, facilitate\ndissociative chemisorption, and transport protons into the bulk.\nComment 5:\n5. Suggest to discuss whether multiple proton transport mechanisms may coexist under different\ntemperatures or humidity conditions.\nResponse to C5:\nWe appreciate the Reviewer for raising this insightful point regarding temperature-dependent\nproton transport mechanisms.\nIndeed, multiple proton transport mechanisms can coexist and exhibit temperature-dependent\ndominance in perovskite-based oxygen electrodes. At high temperatures (> 600 °C), the accelerated\nreaction kinetics allow proton migration via OH- transport through oxygen vacancy sites, where the\nentire OH- migrates as a mobile species, which is commonly referred to as the vehicle mechanism.\nHowever, as the operating temperature decreases, the activation energy associated with OH- vehicular\ntransport becomes prohibitively high (> 1 eV), rendering this pathway kinetically unfavorable.\nConsequently, at intermediate-to-low temperatures characteristic of PCEC operation (450-\n600 °C), proton transport predominantly occurs via the Grotthuss (hopping) mechanism, wherein\nprotons (H+) form bonds with lattice oxygen atoms and migrate through sequential bond cleavage and\nreformation between adjacent oxygen sites. This proton hopping mechanism is supported by multiple\nlines of evidence, including isotope effect studies (Yoshihiro Yamazaki et al. Nat. Mater., 2025.\nhttps://doi.org/10.1038/s41563-025-02311-w), electrochemical impedance spectroscopy (Zongping\nShao et al. Applied Catalysis B: Environmental, 2023, 331, 122682), and temperature-dependent\nconductivity measurements (K.D Kreuer, Solid State Ionics, 1999, 125, 285-302., Yashima M., Nat.\nCommun. 2023 14, 7466.). It has also been directly observed in similar perovskite systems via\ntechniques such as quasi-elastic neutron scattering (QENS) and neutron diffraction. Crucially, the\ncharacteristic activation energies we observe (typically 0.4–0.6 eV) align with the proton hopping\nmechanism, and are distinct from the higher barriers (>1 eV) typically associated with hydroxyl group\nmigration.\nWe have revised the manuscript to incorporate a discussion of these temperature-dependent\nmechanistic considerations.\nPage 6, lines 122-128:\nProton conduction in perovskite oxide mechanisms exhibit temperature-dependent dominance. At\ne\nl\ne\nv\na\nt\ne\nd\nComment 6:\n6. Clarify how proton uptake affects oxygen vacancy concentration and lattice stability.\nResponse to C6:\nWe thank the Reviewer for this question regarding the interplay between proton uptake, oxygen\nvacancy concentration, and lattice stability.\nAt high temperatures, perovskite-based oxygen electrodes undergo substantial oxygen desorption,\ngenerating oxygen vacancies that serve as catalytically active sites for surface exchange reactions.\nHowever, this extensive vacancy formation simultaneously places the lattice in a thermodynamically\nmetastable state, rendering it susceptible to structural degradation.\nProton uptake through the formation of protonic defects that occupy oxygen vacancy sites. This\nprocess effectively reduces the oxygen vacancy concentration and, more critically, suppresses the\npronounced valence changes of transition metal cations. Besides, the presence of protonic defects\npartially compensated for the lattice contraction upon cooling associated with oxygen during thermal\ncycling. This is aligned with the “hydration-induced chemical expansion offset” concept (Xie et al.\nNature Communications, 2025, 16, 3154). We have revised the manuscript to explicitly discuss this\nmechanistic framework.\nPage 13-14, lines 274-278:\nEnhanced thermomechanical stability originates from two synergistic mechanisms: the homogeneous\nionic distribution enables uniform oxygen release, yielding a reduced and near-linear thermal\nexpansion coefficient; concurrently, proton uptake capability suppresses Co reduction while the\naccompanying chemical expansion compensates for thermal contraction during cooling.\nComment 7:\n7. Please clearly label all APT elemental maps and color scales and discuss dataset representativeness.\nResponse to C7:\nThanks to the reviewers for pointing this out. The needle-shaped specimens for APT analysis\nwere prepared from dense bulk material using focused ion beam milling, thereby mitigating potential\nsurface-related artifacts. Each APT dataset captures over one million atoms, with a minimum analyzed\nvolume exceeding 20 × 20 × 200 nm3. This is far exceeding the unit-cell dimensions, ensuring\nstatistically meaningful analysis of ionic distributions and enabling the detection of possible nanoscale\nsegregation. Furthermore, the compositional homogeneity inferred from APT is corroborated by FT-\nIR (Supplementary Fig. 4). The proton diffusion energy barriers deduced from the structural analysis\nalign consistently with results from electrochemical impedance spectroscopy and DFT calculations.\nThe rest of the APT element distribution shown in Fig. 1g is fully shown in Supplementary Fig. 7.\nWe have made the following changes to the pictures of the manuscript:\nPage 8:\nFig. 1 Dispersed B-site doping eliminates proton diffusion bottlenecks in BCZTZICM perovskite electrodes.\n(a) Refined XRD profiles of the prepared BCZTZICM. (b) HR-TEM image of BCZTZICM along (011), (120), and\n(111) plane. Structure diagram of PCEC, the amplification part is a brief sketch of the ion arrangement of (c) dispersed\nand (d) aggregated at B site of perovskite-type oxygen electrode. Schematic illustrating (e) abundant proton\nconduction pathways facilitated by discrete oxygen vacancies enabled through homogeneous ionic arrangement,\nand\n(f) regions of elevated proton diffusion barrier due to contiguous vacancies. (g) 3D reconstruction of the APT data\nshowing mapping of the BCZTZICM sample with O distributions. Integrated line profiles show chemical\ncomposition alongside the z-axis of the analyzed needle. (h) Proton diffusion energies of BCZTZICM, BSCF, and\nBCFZY in the bulk.\nSupplementary Information, Page 6-8\nSupplementary Fig. 5.\n3D reconstruction of the APT data showing mapping of the BSCF sample with Ba, Sr, Co, Fe, and O\nionic distributions. Integrated line profiles show chemical composition at the (a) apex and (b)\nmidsection regions of the analyzed needle.\nSupplementary Fig. 6.\n3D reconstruction of the APT data showing mapping of the BCFZY sample with Ba, Co, Fe, Zr, Y,\nand O ionic distributions. Integrated line profiles show chemical composition at the apex regions of\nthe analyzed needle.\nSupplementary Fig. 7.\n3D reconstruction of the APT data showing mapping of the BCZTZICM sample with elements\ndistributions. Integrated line profiles show chemical composition alongside the y-axis of the analyzed\nneedle.\nComment 8:\n8. Please provide key DFT modeling details (supercell size, vacancy configuration, proton\nconcentration).\nResponse to C8:\nWe thank the reviewer for this valuable suggestion. The key parameters and their derivations are\nprovided below.\nSupercell Size\nAll density functional theory calculations were performed using a 3×3×3 supercell constructed\nfrom the cubic perovskite unit cell, containing 27 ABO formula units. The optimized supercell lattice\n3\nparameters are a=b=11.8720 Å and c=25.8934 Å (α=β=γ=90°). For the stoichiometric perovskite\nstructure, each ABO unit contains 5 atoms, comprising 1 A-site cation, 1 B-site cation, and 3 oxygen\n3\natoms. Therefore, the total number of atoms in the stoichiometric supercell is 27×5=135 atoms, which\nincludes 27 A-site cations (Ba or Sr/Ba mixture), 27 B-site cations (Co-based with various dopants),\nand 81 oxygen atoms.\nOxygen Vacancy Configuration\nTo model the oxygen-deficient perovskite structure under realistic operating conditions, one\noxygen vacancy was introduced at a specific lattice site within the supercell. The oxygen vacancy\nconcentration δ can be calculated by dividing the number of vacancies by the number of formula units.\nSince we have 1 oxygen vacancy distributed among 27 formula units, the vacancy concentration is δ\n=1/27≈0.037 per formula unit. Alternatively, considering that the supercell contains 81 oxygen sites\nin total, the percentage of vacant oxygen sites is 1/81×100%≈1.23%. This vacancy concentration was\nchosen to represent a dilute defect limit while maintaining computational efficiency.\nProton Concentration\nFor proton diffusion calculations, one hydrogen atom was incorporated into the oxygen-deficient\nsupercell to simulate the proton-conducting state of the material. The proton concentration x per\nformula unit is calculated as the ratio of hydrogen atoms to the number of formula units, yielding\nx=1/27≈0.037 per formula unit. To express this in terms of atomic percentage, we note that after\nintroducing one oxygen vacancy and one proton, the total number of atoms in the supercell becomes\n135−1+1=135 atoms. The hydrogen atomic percentage is therefore 1/135×100%≈0.74 atom %.\nThe sites of calculated vacancy formation energies are presented in Supplementary Fig. 9, while\nthe computational structures and energies related to oxygen vacancy formation energy, hydration and\nproton diffusion barriers are provided in the Supplementary Note 3. We have revised the manuscript\nas follows:\nPage 26, lines 608-611:\nCore-valence electron interactions were described by the projector-augmented wave (PAW)\npseudopotentials54, 55, and the electronic wavefunctions were expanded in a plane wave basis set with\na kinetic energy cutoff of 500 eV. The perovskite structures were constructed with a 3×3×3 supercell\ncontaining 27 ABO formula units, and lattice parameters of 11.8720 Å×11.8720 Å×25.8934 Å\n3\n(α=β=γ=90°). All model coordinates related to calculation are shown in Supplementary note 3.\nElectronic convergence was achieved when the energy difference between successive iterations was\nless than 10-5 eV, while geometric optimization was considered complete when the forces on individual\natoms were below 0.05 eV/Å.\nComment 9:\n9. Clarify whether reduced polarization resistance is limited by proton transport, oxygen exchange, or\ncoupled kinetics.\nResponse to C9:\nThank reviewer for this important question. Our systematic kinetic analyses clarify that the\nnotably reduced polarization resistance of the BCZTZICM electrode originates primarily from its\nsignificantly enhanced proton conduction. The DRT analysis of symmetric cells based on BZCYYb\n(Fig. 4e) shows a substantially weakened relaxation in the mid-frequency range, corresponding to a\nmajor improvement in bulk proton migration consistent with the low proton diffusion barriers revealed\nby DFT calculations. The weaker low-frequency response reflects enhanced surface proton uptake,\naligning with the in-situ FT-IR and H O-TPD results. Concurrently, symmetric cell tests based on\nmeasurements verify a surface oxygen exchange coefficient slightly inferior to that of BSCF, and better\nthan BCFZY. Thus, the excellent overall performance stems from the synergy between enhanced bulk\nproton conduction, improved surface proton uptake, and sufficient, stable intrinsic oxygen exchange\nactivity. We have revised manuscript as below:\nPage 17, lines 346-349:\nAs shown in Fig. 4e, BCZTZICM exhibits substantially reduced contributions across all\nfrequency ranges, with dramatic improvements in the mid-frequency regions corresponding to bulk\nproton transport. The weaker low-frequency response reflects enhanced surface proton uptake,\naligning with the in-situ FT-IR and H O-TPD results. This confirms atomic-scale compositional\n2\nhomogenization creates extensive facile proton transport pathways optimizes proton uptake, and\nretains sufficient intrinsic oxygen activity. Comparison with state-of-the-art electrode materials\nconfirm the exceptional performance of BCZTZICM (Fig. 4f). ASRs for BCZTZICM at 600 °C (0.163\nΩ cm2) substantially exceed advanced electrode metrics, establishing BCZTZICM as one of the\nhighest-performing materials currently available9, 20, 27, 29, 30, 31, 32, 33, 34, 35, 36, 37, 38.\nComment 10:\n10. Is that possible to quantify and compare crack width/density among electrodes to support\nmechanical stability claims.\nResponse to C10:\nWe thank the reviewer for this suggestion. As recommended, we have now included comparison\nto support the claims of superior mechanical stability. The thermal cycling conditions and the\ncorresponding crack widths observed in post-test microstructural analysis for reference electrodes are\nsummarized in Supplementary Table 2. This data provides direct, comparative evidence for the\nenhanced thermomechanical durability of BCZTZICM oxygen electrode.\nPage 13, lines 272-274:\nThe thermal cycling conditions and the corresponding crack widths observed in post-test\nmicrostructural analysis for reference electrodes are summarized in Supplementary Table 2.\nSupplementary Information, Page 42:\nSupplementary table 2. TEC of oxygen electrodes and crack width of electrode after thermal\ncycling.\nOxygen Thermal cycle conditions Crack\nTEC\nelectrode Heating rate Cycle width References\n( C\nmaterials ( number (μm)\n10-6 K-1) o\nSNC C min-1) o\n(based on SDC 20.5 30 7.5 40 ~1/0.6\nl\nsymmetric cell) i\nBSCF\n22.4 30 n6 40 ~0.9\n(based on SDC\ng\nsymmetric cell) r\na\nBCC\nt\n(based on\n25.89 10 6.7 35 ~0.78/0.6\nBZCYYb\nsymmetric cell)\nBCFZY\n(based on\n22.56 20 10 40 1.38/1.54 This work\nBZCYYb\nsymmetric cell)\nBSCF\n(based on\n26.08 20 10 40 1.89 This work\nBZCYYb\nsymmetric cell)\nSF\n(based on\n30 30 10 35 ~0.82/0.75\nBZCYYb\nsymmetric cell)\nPBCF 1.6 1.6 5\n(based on 10 5 15\n18.5 ~4.2\nBZCYYb\n20 10 50\nsingle cell)\nComment 11:\n11. Please explain whether improved thermomechanical stability arises mainly from suppressed\noxygen release or proton uptake.\nResponse to C11:\nThank you for this valuable insight and comments. The improved thermomechanical stability\nstems from two synergistic mechanisms. Primarily, the significantly reduced thermal expansion\ncoefficient (TEC) and its near-linear temperature dependence originate from a more homogeneous\noxygen release, enabled by the homogeneous ionic distribution. Secondly, the excellent hydration\ncapability allows for pronounced proton uptake at elevated temperatures. This process not only\nsuppresses the thermal reduction of Co ions but also utilizes the associated chemical expansion to\noffset the drastic thermal contraction during cooling cycles.\nPage 13-14, lines 274-278:\nEnhanced thermomechanical stability originates from two synergistic mechanisms: the homogeneous\nionic distribution enables uniform oxygen release, yielding a reduced and near-linear thermal\nexpansion coefficient; concurrently, proton uptake capability suppresses Co reduction while the\naccompanying chemical expansion compensates for thermal contraction during cooling.\nComment 12:\n12. Please discuss dominant long-term degradation mechanisms observed or anticipated beyond 780\nh.\nResponse to C12:\nThank you for raising this critical point. While the PCEC in this manuscript demonstrated\nexcellent stability over 780 hours at 550 °C with low degradation rates (<20 μV h-1), identifying and\nmitigating potential long-term degradation mechanisms is indeed crucial for achieving the thousands\nof hours of operation required for commercialization.\nBased on our experimental observations and analysis, beyond the intrinsic chemical and\nmicrostructural stability of the materials, a key long-term failure risk originates from the silver paste\ntypically used for sealing during button-cell testing. Under prolonged exposure to high temperature\n(>500 °C) and humid atmospheres, silver paste is susceptible to oxidation, creep, thermal stress\nmismatch, and mechanical degradation, leading to gradual seal failure. This can introduce additional\ngas leakage and crossover, not only increasing the ohmic resistance but also potentially triggering\nirreversible redox reactions at the electrodes, which could become the dominant factor in accelerated\nperformance decay. Therefore, in future research and device development, we will prioritize\noptimizing the cell sealing strategy: moving away from the susceptible silver paste toward more\npromising approaches, such as high-temperature ceramic-based sealants or compressed mica gasket\nseals. Concurrently, we will continue to monitor the phase stability of the hydrogen electrode and the\nsurface evolution of the oxygen electrode under polarization. Through these systematic engineering\nand material optimizations, we aim to achieve even lower degradation rates, meeting the stringent\nrequirement for extended operational lifetime in PCEC.\nPage 19, lines 393-395:\nThe long-term galvanostatic testing at 550 °C (Fig. 5f) demonstrates stable performance over 780\nhours, with degradation rates of 19.3 and 16.9 μV h-1 in fuel cell mode and electrolysis mode,\nrespectively. The degradation originates from seal failure induced by silver sealant oxidation, creep,\nthermal stress mismatch, and mechanical deterioration under high-temperature and humid atmospheres.\nRecent studies on oxygen electrode materials and corresponding long-term stability in PCEC\nsummarized in Supplementary Table 547, 48.\nComment 13:\n13. Please strengthen comparison with state-of-the-art oxygen electrodes using normalized metrics\nwhere possible.\nResponse to C13:\nWe thank the reviewer for raising this point. We have added a comparison of polarization resistance\nto the single-cell performance section, providing a more direct and standardized assessment of the\nelectrode performance.\nPage 19, lines 382-384:\nPerformance and impedance comparison establishes BCZTZICM among higher-performing PCEC\nsystems (Fig. 5d and Supplementary Table 3)4, 20, 27, 28, 30, 37, 38, 39, 40, 41, 42, 43, 44, 45, 46.\nSupplementary Information, Page 43:\nSupplementary Table 3. Performance and impedance comparison of oxygen electrodes in fuel\ncell and electrolysis modes at 600 °C.\nPolarization\nPeak power density Current density at\nOxygen electrode resistance at OCV References\n(mW cm-2) 1.3 V (mA cm-2)\n(Ω cm2)\nCBSLCC 1660 0.112 1760\nPBSLCC 1160 0.112 1750\nBCZTZICM 1560 0.11 2000 This work\nPLNBSCC 1210 0.17 1950\nBCT20 with PLD\n1640 0.12 1600\nlayer\nBSC-PBSCF 1640 0.04 /\nPBCHf10 1490 0.07 2780\nBCCY 743 0.113 /\nPNMCFC-PBC 1720 0.08 2170\nN-XFN 790 0.167 1700\nBCFN 1207 0.197 1511\nC/H-BSCF 1670 0.157 1230\nBSCF-RC 1100 0.13 1490\nx\nNAUP-PNC 1501 0.07 1316\nPBN-BCZYYb\n1160 0.25 1460\nCCS\nComment 14:\n14. I am curious whether similar performance could be achieved with fewer dopant elements.\nWe thank the Reviewer for this insightful question. As discussed in the manuscript, multielement\nsynergy is essential for achieving ionic dispersion, strain balance, and valence modulation, objectives\nthat cannot be readily accomplished with fewer dopant elements, as exemplified by BCFZY and BSCF.\nTo conceptualize this, we calculated the configurational entropy as a function of dopant species\nnumber at a fixed 20% B-site dopant concentration with 80% cobalt content (see figure below). The\nresults reveal that configurational entropy increases progressively but with diminishing returns as more\ndopant species are introduced, suggesting that structural stabilization is achieved beyond a certain\ncompositional complexity. However, excessive doping would compromise electrocatalytic activity\ndue to over-enhanced lattice distortion and excessively strengthened metal-oxygen bonding that\nreduces lattice flexibility.\nTo address the Reviewer's question experimentally, we synthesized single-element doped\nBaCo X O (X=Zr, Ti, Zn, In, Cu, Mo) compositions. Unfortunately, none of these exhibited the\n0.8 0.2 3-δ\ndesired cubic perovskite structure. We also synthesized a quaternary-doped cubic perovskite,\nBaCo (Zr Ti Cu Mo ) O (BCZTCM), and evaluated its electrochemical performance using\n0.8 1/4 1/4 1/4 1/4 0.2 3-δ\nsymmetric cells based on both GDC and BZCYYb electrolytes (see figure below). The results\ndemonstrate that BCZTZICM with six micro-dopants exhibits superior performance compared to the\nquaternary-doped analogue.\nIn future work, we will continue to refine this design concept by systematically varying the active\nelement content to develop further optimized oxygen electrodes."}]

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