Preserving a Kinetically‐Metastable Nanophase by Limited Calcination for High‐Performance Protonic Ceramic Cells

ABSTRACT

Sluggish oxygen reduction/evolution reactions (ORR/OER) at the air electrode critically limit the efficiency of reversible protonic ceramic cells (r‐PCCs), yet conventional high‐temperature calcination of the air electrode leads to undesired surface passivation, while low‐temperature calcination leads to insufficient crystallization, both of which severely impair electrocatalytic activity. Here we employ a thermally‐limited calcination process to kinetically retain a nanophase in the air electrode that only forms at a selected calcination temperature (termed “metastable”), thereby forming rich heterointerfaces for active ORR/OER. Specifically, controlled calcination temperature induces selective Ce incorporation into the host lattice of Ba(Co,Fe,Y)O_3‐δ while preserving the metastable BaCeO₃‐related nanophase. This nanostructure enriches oxygen‐vacancy‐related defects, accelerates surface exchange and bulk diffusion, promotes proton incorporation, and improves thermomechanical compatibility with the electrolyte. The as‐developed electrode (BaCo_0.6Fe_0.2Y_0.1Ce_0.1O_3‐δ‐BaCeO₃) exhibits a low resistance of 0.38 Ω cm² at 550°C. Single cells with this electrode deliver a high peak power density of 1.44 W cm⁻² at 650°C and an electrolysis current density of −2.47 A cm⁻² at 1.3 V. These findings establish a promising strategy of thermally‐limited calcination for designing high‐performance reversible protonic ceramic cells.

Calcination history becomes a materials‐design variable for reversible protonic ceramic cells. Thermally limited processing kinetically preserves a metastable BaCeO₃‐related nanophase within a Ce‐modified BCFY matrix, creating embedded heterointerfaces that reshape defect chemistry, promote hydration, and accelerate oxygen exchange and diffusion. This limited‐calcination‐enabled nanophase‐preservation strategy enables efficient and durable fuel‐cell/electrolysis operation.

1. Introduction

The growing deployment of intermittent renewable electricity and stable computing power demand of the data center have intensified the need for electrochemical technologies that can operate efficiently in both energy‐conversion and energy‐storage modes [1, 2, 3]. Reversible protonic ceramic cells (r‐PCCs) are particularly attractive in this context because they combine fuel‐cell and electrolysis functions within a single device while operating at intermediate temperatures, where faster start‐up, improved operational stability, and lower system cost become possible compared with conventional high‐temperature solid oxide cells [4, 5, 6, 7]. These advantages arise from the fast transport and relatively low activation energy of protonic defects in ceramic electrolytes [8, 9, 10]. Despite this promise, the performance of r‐PCCs remains strongly limited by the air electrode, which must sustain rapid oxygen reduction and evolution kinetics while maintaining coupled ionic/electronic transport and structural stability under humid oxidizing conditions [11, 12].

Considerable effort has therefore been devoted to developing mixed protonic, oxide‐ion, and electronic conducting air electrodes, particularly multi‐cation Co/Fe‐based perovskites, BaCo_0.4Fe_0.4Zr_0.2O_3‐δ (BCFZ), PrBaCo₂O_5‐δ (PBC), SrCo_0.8Fe_0.15Zr_0.05O_3‐δ (SCFZ), and Ba_0.5Sr_0.5Co_0.8Fe0_.2O_3‐δ (BSCF) [13, 14, 15, 16]. Their compositional flexibility allows broad tuning of oxygen nonstoichiometry, electronic structure, and transport behavior [17, 18]. Within this family, BaCo_xFe_1‐xO_3‐δ‐based conductors are attractive parent platforms because their cubic phases can exhibit fast coupled transport and strong oxygen electrocatalytic activity at intermediate temperatures [19]. In particular, Y‐substituted BaCo_0.7Fe_0.2Y_0.1O_3‐δ (BCFY) has been shown to deliver relatively low polarization resistance through tolerance‐factor tuning and oxygen‐vacancy generation induced by large B‐site cations [20, 21]. Even so, further improvement is still needed in terms of hydration‐related functionality and steam tolerance [22]. In this regard, Ce is an attractive dopant because its reversible Ce⁴⁺/Ce³⁺ redox chemistry and oxygen‐storage capability can regulate local defect chemistry and potentially enhance hydration‐related behavior [23, 24, 25, 26, 27]. Despite these possible advantages, most Ce‐modification strategies still remain within the framework of single‐phase optimization. In most reported systems, Ce is incorporated into a single‐phase perovskite lattice, while conventional high‐temperature calcination further drives the material toward a chemically homogeneous and thermodynamically stable state [28, 29]. Although favorable for phase purity, such a framework limits the regulation of local reaction environments and often leaves the trade‐off between catalytic activity and durability unresolved [30, 31, 32].

Dual‐phase and composite electrodes provide one possible route beyond these limitations by combining distinct functionalities within a single system [30]. However, in most reported systems, the second phase is introduced either by physical mixing or by uncontrolled phase separation during synthesis, making its distribution, length scale, and interfacial coupling difficult to control [33, 34]. This limits the understanding of the mechanism and the optimization of the material. More broadly, current air‐electrode design still relies mainly on equilibrium single‐phase selection, whereas the possible functional value of non‐equilibrium phase states has received far less attention.

Unlike post‐deposition electrode sintering, which mainly affects microstructure and interfacial contact, precursor calcination governs phase evolution and cation partitioning before electrode fabrication. This difference opens an alternative design opportunity, rather than driving the electrode fully to equilibrium during synthesis, it may be advantageous to preserve a functionally beneficial non‐equilibrium nanophase through deliberate control of calcination history. Such phases are intrinsically difficult to retain in multi‐cation oxides because continued calcination promotes cation redistribution and phase homogenization toward the thermodynamic end state. If preserved in the lower operation temperature by kinetic limitation, however, they may introduce local chemical environments inaccessible to equilibrium single‐phase materials, including defect‐rich heterointerfaces and chemically distinct reaction sites. For protonic ceramic air electrodes, this concept is especially appealing when the retained phase is itself favorable for hydration. BaCeO₃ is a representative example because of its strong water affinity and proton‐related functionality [35] Preserved as a nanoscale secondary phase, it may facilitate water uptake, enrich local vacancy chemistry, promote surface oxygen kinetics, and generate abundant heterointerfaces with the host phase [36, 37, 38]. The challenge is that such a phase is unlikely to be maintained in conventional calcination unless its evolution toward equilibrium is kinetically arrested.

Here, we show that limited calcination enables such a non‐equilibrium electrode architecture in a BaCo_0.6Fe_0.2Y_0.1Ce_0.1O_3‐δ‐BaCeO₃ (cc‐BCFY) system. Specifically, limited calcination restricts Ce cation diffusion and suppresses phase transformation kinetics, thereby preserving a kinetically “metastable”, non‐equilibrium BaCeO₃‐related nanophase that only forms at selected calcination temperatures and producing a tightly coupled dual‐phase heterostructure. Importantly, the preserved BaCeO₃‐related nanophase is not a residual impurity from incomplete synthesis, but a structurally integrated component that reshapes the local defect chemistry and interfacial reaction environment. In this work, BCFY is used as the Ce‐free benchmark, while m‐BCFYC and cc‐BCFY are introduced as stepwise controls to compare with the Ce‐containing single‐phase perovskite and the limited‐calcination‐derived dual‐phase heterostructure. Compared with single‐phase BaCo_0.7Fe_0.2Y_0.1O_3‐δ (BCFY), cc‐BCFY exhibits enhanced hydration‐related behavior, more favorable oxygen‐vacancy characteristics, faster surface and bulk oxygen kinetics, and improved thermomechanical compatibility with the electrolyte. Specifically, the resulting r‐PCC delivers a peak power density of 1.44 W cm⁻² at 650°C and an electrolysis current density of −2.47 A cm⁻² at 1.3 V, together with stable operation for over 400 h in fuel cell mode and over 150 h in electrolysis cell mode at 550°C. Rather than relying solely on the discovery of new triple‐conducting oxides, this work demonstrates that controlling thermal history to preserve kinetically‐metastable nanophase provides an additional design dimension for engineering air electrodes in reversible protonic ceramic cells.

2. Experimental Section

2.1. Powder Synthesis

All powders were synthesized via a sol‐gel route. Ba(NO₃)₂, Co(NO₃)₂·6H₂O, Fe(NO₃)₃·9H₂O, Y(NO₃)₃·6H₂O, and Ce(NO₃)₃·6H₂O (Aladdin, analytical grade) were used as starting materials, and the batch size was set to 0.05 mol. Taking BaCo_0.6Fe_0.2Y_0.1Ce_0.1O_3‐δ‐BaCeO₃ (cc‐BCFY) as an example, stoichiometric amounts of the metal nitrates were dissolved in deionized water at 80°C in a glass vessel. Ethylenediaminetetraacetic acid (EDTA) and citric acid (CA) were then added as chelating agents, with a molar ratio of EDTA:CA:total metal ions = 1:2:1. The pH was adjusted to 8 by adding aqueous ammonia. The solution was heated to evaporate water and form a gel, which was subsequently dried at 180°C for 8 h to obtain a solid precursor. The precursor was calcined in ambient air at 1000°C for 5 h and thoroughly ground to yield the final product. The electrolyte BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3‐δ (BZCYYb), as well as the air‐electrode powders BaCo_0.7Fe_0.2Y_0.1O_3‐δ (BCFY) and Ba_0.979Co_0.6Fe_0.2Y_0.1Ce_0.079O_3‐δ (m‐BCFYC), were synthesized using the same procedure. For comparison, the cc‐BCFY precursors were also calcined at 900°C and 1100°C for 5 h, denoted as BCFYC900 and BCFYC1100, respectively.

2.2. Cell Fabrication

BZCYYb powders were dry‐pressed into pellets (0.35 g each) and sintered in air at 1450°C to obtain dense electrolyte substrates. Air‐electrode inks were prepared by ball‐milling 1 g of electrode powder with 10 mL isopropanol, 2 mL ethylene glycol, and 0.8 mL glycerol for 30 min. The inks were spray‐coated onto both sides of the electrolyte substrates and fired in air at 800°C for 2 h to fabricate electrolyte‐supported symmetric cells with an effective electrode area of 0.56 cm². Ag paste was then applied onto the electrode surface at 130°C and connected with Ag wires as the current collector. Single cells with a fuel electrode (Ni+electrolyte)| electrolyte|air electrode configuration were fabricated via a dry‐pressing/spin‐coating/co‐sintering process. NiO, BZCYYb, and soluble starch were mixed at a mass ratio of 6:4:1 in ethanol and ball‐milled for 40 min. After evaporation of the ethanol, the dried mixture was ground into fine powder and dry‐pressed to form the fuel‐electrode support. The electrolyte slurry was prepared by dispersing BZCYYb powder in a vehicle containing terpineol (1.5 times the mass of BZCYYb), 4 wt% ethyl cellulose, and 1 wt% nano‐NiO, followed by ball milling for 60 min. The slurry was spin‐coated onto the support at 4000 rpm for three cycles to form the electrolyte layer, and the resulting half‐cells were co‐sintered in air at 1450°C for 5 h. Finally, the air‐electrode ink was spray‐coated onto the center of the electrolyte surface (active area: 0.28 cm²) and fired in air at 800°C for 2 h to obtain the complete single cells.

2.3. Electrochemical Measurement

Electrochemical impedance spectroscopy (EIS) was measured using a Princeton electrochemical workstation over a frequency range of 100 kHz to 0.1 Hz. Measurements were conducted from 700°C down to 500°C with a temperature step of 50°C. The EIS data were fitted with equivalent circuits using ZView software. The ohmic resistance (R_ohm) was obtained from the high‐frequency intercept in the Nyquist plots, and the polarization resistance (R_p) was determined from the difference between the low‐ and high‐frequency intercepts; the area‐specific resistance (ASR) was calculated by multiplying the corresponding resistance by the active area. Distribution of relaxation times (DRT) deconvolution was further applied to resolve individual processes in the impedance spectra. For electrical conductivity relaxation (ECR), 0.8 g of air‐electrode powder was dry‐pressed into bar‐shaped specimens and tested using a Keithley 2460. The samples were first equilibrated in air at a flow rate of 100 mL min⁻¹, followed by a rapid gas switch to 10% O_2%–90% N₂ to trigger the relaxation response. The transient voltage response was recorded until a steady state was reached, and the relaxation curves were fitted to quantify the chemical diffusion coefficient (D_chem) and surface exchange coefficient (k_chem). Single‐cell tests were performed by sealing the cell in a quartz tube using a silver‐based sealing adhesive. The tests were repeated multiple times to ensure reproducibility. Unless otherwise stated, the data shown are representative, the fuel electrode faced the inner side of the tube and was supplied with hydrogen (100 mL·min⁻¹), while the air electrode was exposed to ambient air. Polarization curves (I‐V and I‐P) were collected using a Gamry workstation.

2.4. Characterization

X‐ray absorption spectroscopy (XAS) at the Co K‐, Fe K‐, and Ce L₃‐edges was performed at the ID26 beamline of the European Synchrotron Radiation Facility (ESRF, Grenoble, France). The storage ring operated at 6.0 GeV in top‐up mode. Incident energies were selected with a Si(311) double‐crystal monochromator, and higher‐order harmonics were suppressed by detuning the monochromator. Co and Fe K‐edge spectra were collected in transmission mode over ∼7.6–8.5 keV and ∼6.9–7.8 keV, respectively. Ce L₃‐edge spectra were also measured in transmission mode over ∼5.6–6.4 keV, covering the pre‐edge, XANES, and EXAFS regions. Samples were finely ground, mixed with cellulose, and pressed into self‐supported pellets with thickness optimized for transmission measurements. All measurements were conducted at room temperature under ambient conditions. Energy calibration was performed using Co, Fe, and Ce metal foils as references. Room‐temperature powder X‐ray diffraction (XRD; Bruker D8 Advance, Germany) was used to determine the crystalline structures. Data were collected with filtered Cu Kα radiation at 40 kV and 40 mA over 2θ = 10°–90° with a scan rate of 10° min⁻¹ (step‐scan mode). In situ high‐temperature XRD (PANalytical Empyrean equipped with a high‐temperature stage) was conducted in flowing air with a heating rate of 10 K·min⁻¹; each temperature was held for 60 min to evaluate the thermal structural stability of cc‐BCFY. Microstructures were characterized using double‐Cs‐corrected scanning transmission electron microscopy (STEM‐HAADF, FEI Titan Cubed Themis G2 300), high‐resolution transmission electron microscopy (HR‐TEM, JEOL JEM‐2100F), and scanning electron microscopy (SEM, FEI QUANTA‐2000 and Hitachi S4800, 20 kV). Oxidation states were analyzed by XPS (Thermo Scientific ESCALAB 250Xi, USA) at room temperature. Thermogravimetric analysis (TGA; TA DISCOVERY 5500) was performed from 25°C to 800°C. Proton uptake of the air electrode was assessed by H₂O temperature‐programmed desorption (H₂O‐TPD; Hiden Analytica, HPR‐20) and Fourier‐transform infrared spectroscopy (FTIR, Thermo Fisher Scientific Nicolet iN10, USA), while oxygen‐release behavior was evaluated by O₂ temperature‐programmed desorption (O₂‐TPD; Hiden Analytica, HPR‐20). Thermal expansion coefficients (TECs) were measured in air from 25°C to 1000°C using a dilatometer (Netzsch DIL 402 C/3/G).

2.5. DFT Calculation

All calculations were performed within the density functional theory framework [39, 40] using VASP and the projector‐augmented wave formalism together with a plane‐wave basis [41, 42]. The exchange‐correlation energy was evaluated with the PBE functional under the generalized gradient approximation [43]. A plane‐wave cutoff energy of 450 eV was adopted. Reciprocal‐space sampling was carried out with Γ‐centered Monkhorst‐Pack grids (3 × 2 × 1 for both BCFY and cc‐BCFY) [44], constructed via VASPKIT [45]. Atomic positions were relaxed until the largest force component on any atom fell below 0.02 eV Å⁻¹, while electronic self‐consistency was deemed converged at 1 × 10⁻⁵ eV. Activation barriers were obtained from climbing‐image nudged elastic band (CI‐NEB) calculations, in which three intermediate replicas were generated by linear interpolation [46]. Structural visualization was performed using VESTA [47].

To quantify adsorption, slab models exposing the (01–1) facets of BCFY and cc‐BCFY were employed to compute the adsorption energies of H₂O and O₂. A 20 Å vacuum spacing was introduced along the surface‐normal c direction to minimize interactions between periodic slabs. The same Γ‐centered 3 × 2 × 1 k‐point sampling was used for the surface calculations.

Given the compositional complexity of the Co/Fe/Ce‐containing oxide, the present calculations were used mainly for comparative trend analysis rather than for determining absolute thermodynamic stability. No exhaustive enumeration of all possible cation arrangements, oxygen‐vacancy configurations, or interface geometries was attempted. The interface model was constructed based on the experimentally identified BCFY‐derived matrix and BCO‐related phase and was fully relaxed under the same convergence criteria as the BCFY model. Therefore, the calculated proton‐migration barriers, hydration energies, oxygen‐vacancy formation energies, and adsorption energies should be interpreted as relative descriptors for comparing BCFY and cc‐BCFY, not as a complete description of all possible local configurations.

3. Results and Discussion

3.1. Thermally‐Limited Calcination Strategy and Material Synthesis

As schematically illustrated in Figure 1a, we propose a thermally‐limited calcination concept for constructing r‐PCC air electrodes. Rather than relying on conventional one‐step synthesis to define a fixed, single‐phase cation distribution, a sol‐gel route is used to prepare Ce‐containing BCFY precursors (BCFYC), after which limited calcination directs the phase evolution. Specifically, during sintering, temperature‐dependent Ce solubility and Ba‐Ce phase formation drive controlled Ce partitioning, such that Ce is incorporated into the perovskite lattice while a nanoscale BaCeO₃ (BCO) phase is kinetically forming in situ and remains embedded within the host matrix. This process yields a tightly coupled heterostructure, denoted as cc‐BCFY, in which the phase fraction, spatial distribution, and interfacial coupling are regulated by the calcination condition, thereby enabling deliberate tuning of defect chemistry and interfacial functionality beyond single‐phase electrodes.

FIGURE 1.

Electrode material synthesis and structural characterization. (a) Schematic illustration of thermally‐limited calcination process; (b) XRD Rietveld refinement of cc‐BCFY and the extracted phase constitution; (c) In situ XRD patterns of cc‐BCFY during heating/cooling between 25°C and 800°C; (d) High‐resolution HAADF‐STEM images of the lattice fringes and heterointerface between the main phase and the kinetically‐ metastable BCO phase; (e) Elemental maps acquired by EDX in the AC‐TEM; (f) EDX line‐scan across the heterointerface.

To identify the temperature window for limited‐calcination‐enabled preservation of a kinetically‐metastable BCO‐related nanophase and controllable Ce partitioning, Ce‐containing precursors were sintered at 900°C, 1000°C, and 1100°C. The resulting air electrodes were denoted as BCFYC900, cc‐BCFY, and BCFYC1100, respectively, with BCFY sintered at 1000°C used as the benchmark. Room‐temperature X‐ray diffraction (RT‐XRD) confirms that pristine BCFY exhibits a single perovskite phase (Figure S1). In contrast, for BCFYC900 and cc‐BCFY, the XRD patterns display a weak diffraction feature at 2θ ≈ 28.8°, indicative of kinetically assisted retention of a BCO‐related phase within a specific calcination‐temperature window. Compared with the sample sintered at 1000°C, the 900°C counterpart shows broader and weaker reflections, suggesting that insufficient crystallization and incomplete phase formation may accompany low‐temperature processing. When the sintering temperature is increased to 1100°C, this weak peak disappears and the diffraction pattern reverts to a single‐phase perovskite signature, consistent with enhanced Ce solubility at elevated temperature that suppresses the retention of the BCO phase [48, 49]. These results directly support that the dual‐phase heterostructure is not an incidental by‐product, but a limited‐calcination‐controlled outcome governed by the precursor sintering temperature. SEM images (Figure S2) show well‐sintered grains and a dense microstructure; nevertheless, the low fraction and fine dispersion of the preserved phase make direct phase contrast challenging in SEM, consistent with the nanoscale nature of the secondary phase.

To quantify the phase constitution and structural parameters of the limited‐calcination‐enabled heterostructure, XRD Rietveld refinement was performed for cc‐BCFY (Figure 1b). The material is dominated by a cubic perovskite phase (Pm‐3m), Ba_0.979Co_0.6 Fe_0.2Y_0.1Ce_0.079O_2.958 (98.02 wt%, a = b = c = 4.158409 Å), accompanied by ∼1.98 wt% of a Pm‐3m BCO phase (a = b = c = 4.39696 Å) (Supporting Information Table S1). We additionally synthesized a single‐phase Ce‐doped perovskite, Ba_0.979Co_0.6Fe_0.2Y_0.1Ce_0.079O_3‐δ (denoted as m‐BCFYC, corresponding to the main phase of cc‐BCFY), and single‐phase BaCeO₃ (corresponding to the secondary phase of cc‐BCFY) by sintering at 1000°C (Figure S3). In this way, m‐BCFYC serves as a single‐phase Ce‐containing matrix control corresponding to the main phase of cc‐BCFY, while cc‐BCFY further represents the limited‐calcination‐derived dual‐phase sample. This stepwise comparison allows the parent BCFY chemistry, the Ce‐containing single‐phase matrix, and the preserved BCO‐containing heterointerface to be compared more clearly. Rietveld refinements for m‐BCFYC and BCFY (Figure S4) show that the lattice constant increases markedly after Ce incorporation. Considering the larger ionic radii of Ce³⁺ (1.01 Å) and Ce⁴⁺ (0.87 Å) relative to Co³⁺/Co⁴⁺ (0.61 Å and 0.53 Å) [50, 51], the observed lattice expansion provides direct crystallographic evidence that Ce is incorporated into the perovskite B site rather than merely forming an external Ba‐Ce phase.

To evaluate thermo‐structural stability, in situ XRD was collected during heating from 25°C to 800°C and cooling back to 25°C. As shown in Figure 1c, cc‐BCFY remains phase‐stable throughout the cycle, with no detectable impurity peaks or peak splitting. The diffraction peaks shift to lower 2θ values upon heating and largely recover upon cooling, indicating reversible lattice expansion/contraction and robust phase integrity under thermal cycling, which is essential for durable operation in r‐PCCs.

We then validate the dual‐phase heterointerface at the nanoscale by high‐resolution HAADF‐STEM. As shown in Figure 1d, the (110) lattice‐fringe spacings of the main phase and the kinetically‐metastable BCO phase in cc‐BCFY are 0.2942 and 0.3156 nm, respectively, whereas BCFY exhibits a smaller (110) spacing of 0.2720 nm (Figure S5), consistent with the refined lattice parameters. A well‐defined atomic‐level contact is observed across the heterointerface, and a semi‐coherent feature leads to enlarged lattice spacing of the main phase near the interface. Geometrical phase analysis (Supporting Information Figure S6) further reveals that the strain is predominantly localized within a few nanometers around the interface: ε_xx and ε_yy exhibit clear tensile/compressive contrast, while ε_xy is markedly enhanced in the same region, indicating strong heterogeneous deformation. Such a localized interfacial strain field provides a structural basis for defect formation and diffusion near the heterointerface, thereby offering a potential handle for performance regulation. The energy‐dispersive X‐ray (EDX) elemental mapping of cc‐BCFY acquired under AC‐TEM is shown in Figure 1e. Distinct Ba‐ and Ce‐enriched regions embedded in the main matrix are clearly observed, and their elemental constitution is consistent with a BCO‐related phase. Re‐prepared cc‐BCFY shows similar nanoscale Ba/Ce‐rich features (Figure S7), supporting the reproducibility of the thermally limited calcination process rather than a sample‐specific observation. Line‐scan profiles across the heterointerface (Figure 1f) further show that within a ∼100 nm region, Ba, Ce, and O signals increase markedly while Co, Fe, and Y remain essentially unchanged, providing direct compositional evidence for a spatially coupled dual‐phase configuration. By contrast, BCFY exhibits a largely homogeneous elemental distribution with no discernible nanoscale secondary phase (Figure S8). To further visualize the temperature‐dependent phase‐evolution pathway, TEM‐EDX mapping was additionally performed for the BCFYC900 and BCFYC1100 (Figures S9 and S10). For the BCFYC900, Ba/Ce‐enriched segregated regions are already observable, but these regions still contain detectable Fe and Y signals, suggesting that the low‐temperature secondary phase remains incompletely partitioned rather than forming a well‐defined BCO nanophase. By contrast, the BCFYC1100 no longer exhibits distinct Ba/Ce‐rich secondary domains, and the elemental distribution becomes substantially more homogeneous, consistent with re‐homogenization into a single‐phase perovskite at elevated temperature. This temperature‐dependent evolution argues against the interpretation that the secondary phase in cc‐BCFY is merely a residual impurity from incomplete reaction. Instead, it supports the view that the BCO‐related nanophase is kinetically retained within a limited calcination window.

Chemical compatibility at the electrode/electrolyte interface is essential for durable intermediate‐temperature operation. Since the air electrode is sintered at 800°C for 2 h, we assessed the reactivity between cc‐BCFY and BZCYYb by co‐sintering a 1:1 powder mixture at 800°C for 2 h followed by XRD analysis. As shown in Figure S11, no additional diffraction peaks or peak shifts are observed, indicating negligible interfacial reaction under electrode‐processing conditions. Therefore, cc‐BCFY exhibits good chemical compatibility with BZCYYb, providing a prerequisite for a low‐impedance and durable electrode‐electrolyte interface.

3.2. Interface Defects and Chemical Properties

To reveal how thermally‐limited calcination regulates charge distribution at the cc‐BCFY heterointerfaces, we first probed the electronic structures of Co and Fe in BCFY and cc‐BCFY using X‐ray absorption spectroscopy. As shown in Figure 2a,b, relative to BCFY, the Co and Fe K‐edge XANES spectra of cc‐BCFY exhibit an overall shift toward lower energies, accompanied by pronounced changes in white‐line intensity. Given the established relationship between K‐edge positions and transition‐metal oxidation states(with higher oxidation states typically shifting the absorption edge to higher energies) [52], the concurrent edge downshift and white‐line evolution indicate a reduction in the average oxidation states of Co and Fe in cc‐BCFY. This observation implies a heterointerface‐associated interfacial charge‐compensation process during heterostructure formation. In addition, the pre‐edge features in Figure 2a,b provide insights into local symmetry. For 3d transition metals, K‐edge pre‐edge signals generally arise from 1s to 3d excitations (quadrupole‐allowed) and dipole‐intensity gain enabled by 3d–4p hybridization [53]. When the local structure deviates from centrosymmetry, enhanced p‐d mixing increases the pre‐edge intensity. In contrast, the Co and Fe K‐edge pre‐edge intensities of cc‐BCFY are lower than those of BCFY, suggesting a more centrosymmetric local environment around Co and Fe sites in cc‐BCFY.

FIGURE 2.

Valence evolution and interfacial charge compensation in cc‐BCFY. (a) Co K‐edge and (b) Fe K‐edge XANES spectra of BCFY and cc‐BCFY; WT EXAFS of the (c) Co K‐edges and (d) Fe K‐edges for cc‐BCFY; FT EXAFS spectra at the (e) Co K‐edges and (f) Fe K‐edges for BCFY and cc‐BCFY; (g) Schematic of the STEM‐EELS line scan across the heterointerface, from the kinetically‐metastable BCO phase to the main phase; (h) Comparison of Ba M‐edge EELS collected in different regions; (i) Comparison of Co L‐edge EELS collected in different regions; (j) Comparison of Ce M‐edge EELS collected in different regions.

To correlate the electronic‐structure variations with the evolution of oxygen vacancies, we further employed EXAFS to probe changes in local coordination, and performed multidimensional signal decoupling by combining the WT contour maps of cc‐BCFY (Figure 2c,d) with those of BCFY provided in Figure S12, together with Fourier transform (FT) analyses (Figure 2e,f). Compared with BCFY, the EXAFS features of Co and Fe in cc‐BCFY exhibit reduced amplitudes along with discernible peak‐position changes, indicating a reconstructed local coordination shell [54]. Quantitative fitting of the Co and Fe K‐edge EXAFS (Figure S13) reveals that, for Co, cc‐BCFY exhibits an approximately 14% decrease in first‐shell coordination number together with an average Co‐O bond‐length elongation of ∼0.02 Å, with Fe showing a similar trend (Supporting Information Table S2). This coupled signature, reduced coordination number, elongated metal‐oxygen bonds, and increased structural disorder (Debye‐Waller factor, σ²), provide quantitative evidence for oxygen‐vacancy enrichment [55]. To clarify the role of Ce in this charge redistribution, Ce L‐edge XANES was collected for cc‐BCFY; meanwhile, the corresponding FT/WT analyses and EXAFS fitting (Figure S14 and Table S2) indicate a mixed Ce⁴⁺/Ce³⁺ valence state, confirming that Ce participates in redox‐active charge compensation within the heterostructure.

We next examined the spatial extent of this charge compensation by performing STEM‐EELS line‐scan analysis across the cc‐BCFY heterointerface (Figure 2g). As shown in Figure 2h, the Ba M‐edge spectra acquired within the BCO nanodomains and near the heterointerface exhibit nearly identical peak positions and M₅/M₄ intensity ratios, indicating a similar Ba chemical environment across kinetically‐metastable BCO phase and the interfacial region. Notably, the Co L‐edge signal is only clearly detected in the main phase region (Figure 2i), consistent with the Co‐free nature of the BCO phase. This spectroscopic contrast corroborates the compositional distinction between the preserved kinetically‐metastable BCO phase and the matrix and supports the formation of the dual‐phase heterostructure. This observation supports that the preserved BCO phase does not exist as isolated Ce‐containing particles, but instead remains in intimate contact with the main phase, forming a genuine heterostructure rather than a physical mixture. Consistently, high‐resolution TEM images (Figure S15) reveal that the nanoscale BCO domains remain encapsulated by the main phase, forming intimate and partially semi‐coherent interfaces instead of isolated particles. Such an embedded heterostructure enables both structural and electronic coupling across the interface, providing a physical basis for interfacial charge redistribution and defect modulation.

In contrast, the Ce M‐edge EELS spectra exhibit pronounced spatial variation. As shown in Figure 2j, the preserved metastable BCO domains are dominated by Ce⁴⁺, whereas the interfacial region adjacent to the perovskite matrix displays a clear mixed‐valence Ce⁴⁺/Ce³⁺ signature [56, 57]. Given that reduced transition‐metal oxidation states often correlate with elevated oxygen‐vacancy concentrations [58], Ce may participate in local charge compensation via the reversible Ce³⁺/Ce⁴⁺ redox couple and promote oxygen‐vacancy enrichment at the heterointerface, thereby lowering the average oxidation states of Co and Fe. This process may, in turn, enhance bulk oxygen transport and surface oxygen‐exchange kinetics.

3.3. Electrochemical Performance of the Electrodes

Using BZCYYb as the electrolyte, we evaluated the electrocatalytic performance of the air electrodes in symmetric‐cell configurations by electrochemical impedance spectroscopy (EIS) under humidified air (3 vol% H₂O) over 500°C–700°C. By using the samples sintered at varying temperatures, this symmetric‐cell configuration allows direct assessment of how thermally‐limited calcination influences interfacial kinetics. As shown in Figure 3a, symmetric cells fabricated from BCFYC precursors sintered at 1000°C exhibit the lowest area‐specific resistance (ASR) of 0.38 Ω cm² at 550°C, whereas those sintered at 900°C and 1100°C deliver higher ASR values of 0.64 and 0.51 Ω cm², respectively. Although sintering at 900°C enables the formation of the BCO phase, insufficient crystallinity and incomplete development of the heterointerface likely limit effective charge and mass transport. In contrast, sintering at 1100°C promotes excessive solid‐solution formation, suppressing the metastable BCO phase and driving the system back toward a single‐phase configuration. These results indicate that the optimal electrochemical performance emerges only within a narrow thermal window, where Ce partitioning and BCO phase retention are simultaneously enabled and balanced, highlighting the critical role of thermally‐limited calcination in constructing a functional heterostructure.

FIGURE 3.

Comparison of electrochemical performance in BZCYYb‐supported symmetric cells. (a) ASR comparison of symmetric cells with BCFYC air electrodes prepared from precursors sintered at 900°C, 1000°C, 1100°C in 3% H₂O‐air at 550°C; (b) ASR comparison of symmetric cells with BCFY, m‐BCFYC, and cc‐BCFY air electrodes in 3% H₂O‐air at 550°C; (c) Arrhenius relationship of ASR and comparison of apparent activation energies for symmetric cells with different air electrodes in 3% H₂O‐air over 500°C–650°C; (d) DRT analysis of impedance spectra for symmetric cells with BCFY and cc‐BCFY air electrodes in 3% H₂O‐air at 550°C; (e) Stability test (>200 h) of symmetric cells with BCFY, m‐BCFYC and cc‐BCFY air electrodes in 3% H₂O‐air at 550°C; (f) Nyquist plots before and after the stability test; (g) Benchmark ASR comparison at 550°C against representative air electrodes reported in the literature.

Figure 3b compares the ASR values of BCFY, m‐BCFYC, and cc‐BCFY at 550°C. The ASR decreases from 0.82 Ω cm² for BCFY to 0.52 Ω cm² for m‐BCFYC and further to 0.38 Ω cm² for cc‐BCFY. It should be noted that this comparison does not completely decouple all structural variables, because the heterointerface is intrinsically linked to phase formation during calcination. Nevertheless, all electrodes were processed under the same final electrode‐fabrication condition, and the stepwise decrease in ASR from BCFY to m‐BCFYC and then to cc‐BCFY supports the view that Ce incorporation first improves the perovskite‐derived matrix, while the retained BCO‐containing heterointerface provides an additional kinetic advantage. ASR values over the full temperature range are summarized in Figure S16. Notably, although single‐phase BCFYC1100 exhibits competitive performance at elevated temperatures, its ASR increases markedly in the intermediate‐temperature range, whereas cc‐BCFY maintains consistently low ASR, underscoring the importance of the dual‐phase structure under practical r‐PCC operating conditions. The Arrhenius plots of ASR (Figure 3c) further elucidate the kinetic origin of this behavior. Ce‐containing electrodes exhibit systematically lower apparent activation energies than BCFY (∼125.4 kJ·mol⁻¹), indicating facilitated electrode kinetics. Among them, m‐BCFYC shows the lowest activation energy (∼90.4 kJ·mol⁻¹), reflecting the intrinsic benefit of Ce substitution in reducing the activation barrier for electrocatalysis. Importantly, the limited‐calcination‐enabled cc‐BCFY heterostructure combines a comparably low activation energy with uniformly lower ASR across the entire temperature range, indicating that interfacial polarization losses are further suppressed by the dual‐phase configuration. These results suggest that Ce incorporation in the perovskite‐derived matrix is associated with a lower kinetic barrier, while the retained BCO‐containing heterointerface further suppresses interfacial polarization resistance. Together, this limited‐calcination‐enabled dual‐phase synergy enables concurrent enhancement of ORR activity and oxygen activation capability, consistent with the interfacial charge‐compensation mechanism revealed in Figure 2.

To further elucidate, from a kinetic perspective, how the biphasic heterostructure promotes ORR, we applied the distribution of relaxation times (DRT) analysis to deconvolute the electrochemical impedance spectra of the electrodes [59, 60]. As shown in Figures 3d and S17, all electrodes exhibit three characteristic relaxation processes located in the high‐frequency (HF), intermediate‐frequency (IF), and low‐frequency (LF) regions. These processes are commonly associated with bulk oxygen‐ion transport, electrode–electrolyte interfacial charge‐transfer and surface‐exchange reactions, and gas diffusion, respectively [61, 62].

cc‐BCFY exhibits a pronounced reduction in both the high‐frequency (HF) and intermediate‐frequency (IF) relaxation contributions compared with BCFY, whereas the low‐frequency (LF) component remains relatively unchanged. The concurrent suppression of the HF and IF processes indicates that limited‐calcination‐enabled dual‐phase heterostructures simultaneously facilitate oxygen‐ion transport and interfacial reaction kinetics. Notably, the reduction in the IF contribution is substantially more pronounced than that in the HF region, identifying interfacial charge‐transfer and surface oxygen‐exchange processes as the dominant source of kinetic enhancement. The suppression of the IF relaxation process is fully consistent with the interfacial charge‐compensation and oxygen‐vacancy enrichment revealed by XAS and EELS analyses (Figure 2), which are expected to primarily accelerate surface exchange and interfacial reaction steps. Together, these results indicate that the limited‐calcination‐enabled dual‐phase heterointerface contributes to improved overall oxygen transport and is closely associated with suppression of the rate‐limiting interfacial process governing ORR kinetics under intermediate‐temperature conditions.

To examine whether the limited‐calcination‐enabled dual‐phase heterostructure is associated with improved operational stability, the air electrodes were subjected to continuous EIS monitoring in symmetric‐cell configurations at 550°C under humidified air (3 vol% H₂O). As shown in Figure 3e, the ASR of BCFY increases steadily during operation, indicating progressive kinetic degradation. In contrast, m‐BCFYC exhibits a smaller ASR increase than BCFY, suggesting that Ce incorporation into the perovskite‐derived matrix already contributes to improved stability. Compared with m‐BCFYC, cc‐BCFY shows a further reduction in ASR increase during the testing period, indicating an additional stabilizing effect associated with the preserved BCO‐containing heterointerface. The corresponding impedance evolution before and after testing is compared in Figure 3f. After 200 h, the ASR of BCFY nearly doubles, while the ASR of cc‐BCFY remains essentially unchanged, supporting the improved stability of the limited‐calcination‐enabled heterointerface under humid oxidizing conditions. To benchmark the electrochemical performance, the ASR of cc‐BCFY at 550°C is compared with that of representative air electrodes reported in the literature (Figure 3g; Table S3). cc‐BCFY delivers one of the lowest ASR values among state‐of‐the‐art air electrodes under comparable conditions, highlighting its low‐polarization characteristics at intermediate temperatures. Together, the DRT and stability results support that the limited‐calcination‐enabled dual‐phase heterostructure provides additional kinetic and durability advantages for r‐PCC operation beyond those achieved by Ce incorporation alone.

3.4. Characterizations of Physicochemical Properties

To probe the evolution of oxygen nonstoichiometry and the propensity for oxygen loss in air at elevated temperatures, thermogravimetric (TG) measurements were performed for BCFY and cc‐BCFY under flowing air with a heating rate of 10.00°C·min⁻¹. As shown in Figure 4a, both samples exhibit a slight mass loss below 200°C, which is mainly attributed to the desorption of physically adsorbed water. From 200 to 800°C, the mass of both materials decreases continuously with increasing temperature, reflecting progressively increased oxygen‐vacancy concentration, which is commonly associated with the breaking of Co‐O bonds and concomitant lattice‐oxygen release [63]. Notably, cc‐BCFY exhibits a slightly larger mass loss than BCFY throughout the high‐temperature region, reaching total mass losses of 2.95% and 2.71% at 800°C, respectively. This enhanced oxygen‐release tendency is consistent with the higher oxygen‐vacancy concentration enabled by limited‐calcination‐enabled heterointerface‐associated charge compensation.

FIGURE 4.

Chemical stability and defect‐chemistry characterization. (a) The TG curves of BCFY and cc‐BCFY from RT to 800°C in dry air; XPS spectra of BCFY and cc‐BCFY powders: (b) Co 2p and (c) O 1 s; (d) H₂O‐TPD profiles of BCFY and cc‐BCFY; (e) ECR curves of BCFY and cc‐BCFY measured at 550°C; (f) Temperature‐dependent D_chem and k_chem extracted from ECR fittings; (g) O₂‐TPD profiles of BCFY and cc‐BCFY; (h) EPR spectra of BCFY and cc‐BCFY; (i) Average TEC from 25°C to 1000°C.

Given that ORR kinetics and interfacial charge transfer in perovskite oxides are strongly governed by the oxidation states of B‐site transition metals and the concentration of oxygen vacancies [64, 65], X‐ray photoelectron spectroscopy (XPS) was employed to probe the surface chemical states. As shown in Figure 4b, the Co 2p_3/2 spectra reveal that the fraction of Co⁴⁺ decreases from 23.72% in BCFY to 16.17% in cc‐BCFY, accompanied by an increased Co³⁺ population. This shift toward a more reduced Co chemical environment is commonly associated with enhanced oxygen‐vacancy formation. Consistently, the O 1s spectra (Figure 4c) show a reduced O_lattice/O_ads ratio for cc‐BCFY relative to BCFY, together with an increased contribution from reactive oxygen species (O₂ ²⁻/O⁻), indicating a higher density of oxygen‐vacancy‐related surface sites that are favorable for ORR activity. These trends are in good agreement with the XAS and EELS analyses discussed above.

Hydration is a prerequisite for introducing protons into the air electrode and plays a pivotal role in proton conduction and the kinetics of proton‐coupled electrode reactions. In perovskite oxides, water incorporation proceeds via vacancy‐associated hydration reactions, forming hydroxyl‐type defects (OH^· _O) that migrate through the lattice via the Grotthuss mechanism [66, 67] To probe hydration‐related behavior, H₂O temperature‐programmed desorption (H₂O‐TPD) measurements were conducted (Figure 4d). Both samples exhibit pronounced desorption features below 200°C, attributed to molecularly adsorbed water, whereas desorption above 200°C is mainly associated with hydroxyl‐related species. In the intermediate‐temperature range (200°C–450°C), cc‐BCFY shows a larger integrated desorption area than BCFY, accompanied by a clear shift of the desorption peak toward lower temperature, indicating a reduced apparent desorption barrier for H₂O and hydroxyl‐related species. These features suggest facilitated hydration kinetics enabled by the limited‐calcination‐enabled defect structure. To further corroborate the enhanced hydration behavior, Fourier‐transform infrared (FTIR) spectroscopy was performed after exposure to 3% H₂O‐air at 550°C. As shown in Figure S18, cc‐BCFY exhibits a stronger and broader O─H stretching band (3700–3000 cm⁻¹) than BCFY, indicating enhanced formation of hydroxyl‐type defects.

To probe the influence of the limited‐calcination‐enabled dual‐phase heterostructure on near‐surface oxygen exchange and bulk oxygen transport, electrical conductivity relaxation (ECR) measurements were carried out for BCFY and cc‐BCFY over 500–600°C (Figure 4e; Figure S19). At all tested temperatures, cc‐BCFY exhibits a markedly faster normalized conductivity recovery than BCFY, indicating simultaneous enhancement of surface‐exchange kinetics and bulk chemical diffusion. The temperature dependences of the chemical diffusion coefficient (D_chem) and surface exchange coefficient (k_chem) extracted from the ECR analysis are summarized in Figure 4f. For both electrodes, D_chem and k_chem increase monotonically with temperature, reflecting thermally activated oxygen transport and surface exchange. Across the entire temperature range, however, cc‐BCFY consistently delivers higher D_chem and k_chem values than BCFY, supporting more efficient oxygen incorporation and migration in the limited‐calcination‐enabled dual‐phase heterostructure. At 600°C, D_chem reaches 9.5983 × 10⁻⁴ cm²·s⁻¹ for cc‐BCFY, compared with 2.5236 × 10⁻⁴ cm²·s⁻¹ for BCFY. Meanwhile, the weaker temperature dependence of k_chem for cc‐BCFY indicates a reduced apparent activation barrier for surface exchange, consistent with the interfacial charge compensation and oxygen‐vacancy enrichment revealed above. Arrhenius analysis in Figure 3c further yields a smaller slope for cc‐BCFY, suggesting its lower activation energy for oxygen transport. Although ECR does not directly observe the heterointerface under applied bias, together with the DRT analysis it provides transport‐ and reaction‐relevant evidence that cc‐BCFY exhibits faster oxygen surface exchange, enhanced bulk diffusion, and lower interfacial kinetic resistance than the reference electrodes.

To assess the availability of near‐surface oxygen species in the air electrodes, O₂ temperature‐programmed desorption (O₂‐TPD) measurements were conducted. As shown in Figure 4g, both electrodes display medium‐temperature desorption peaks associated with chemically adsorbed oxygen species (e.g., O²⁻/O⁻) that are weakly bound to the surface [68]. The main desorption peak of cc‐BCFY appears at a slightly lower temperature (409°C) than that of BCFY (412°C), accompanied by a larger peak area, indicating a higher reservoir of accessible oxygen species. These observations further support that the limited‐calcination‐enabled dual‐phase configuration facilitates oxygen migration and enhances surface‐exchange activity. To further assess charge‐transport properties, the electronic conductivity was measured as a function of temperature (Figure S20). Across the measured temperature range, cc‐BCFY exhibits lower electronic conductivity than BCFY. Therefore, the superior electrochemical performance of cc‐BCFY cannot be simply attributed to enhanced electronic transport, but is more consistent with the improved hydration‐related behavior, oxygen transport, and interfacial kinetics revealed above.

To further probe defect structures, particularly the evolution of oxygen‐vacancy‐related defects, electron paramagnetic resonance (EPR) measurements were performed for BCFY and cc‐BCFY. As shown in Figure 4h, both samples exhibit a pronounced resonance signal at g ≈ 2.003, which is commonly attributed to oxygen‐vacancy‐associated trapped electrons (F centers) and/or under‐coordinated oxygen defects, indicating the presence of oxygen‐vacancy‐related defect species in both materials [69]. Compared with BCFY, cc‐BCFY shows a markedly stronger EPR signal, suggesting a higher density of paramagnetic defect centers, typically consistent with an increased concentration of oxygen‐vacancy‐related defects. Such a defect‐rich lattice is expected to facilitate oxygen‐ion transport and surface oxygen exchange. This trend is in line with the larger TG mass loss observed for cc‐BCFY, corresponding to a more pronounced change in oxygen nonstoichiometry, and collectively supports a more active defect evolution and redox responsiveness in cc‐BCFY.

Beyond transport kinetics, long‐term durability of the air electrode also requires good thermo‐mechanical compatibility with the electrolyte [70]. As shown in Figure 4i, the average thermal expansion coefficient (TEC) of cc‐BCFY measured over 25–1000°C is 18.42 ×10⁻⁶ K⁻¹, which is lower than that of BCFY (19.44 ×10⁻⁶ K⁻¹) and closer to that of the BZCYYb electrolyte. This improved TEC matching is expected to alleviate interfacial thermal stress during thermal cycling. Post‐test cross‐sectional observations (Figure S21) reveal pronounced delamination at the BCFY and electrolyte interface after 200 h operation, whereas cc‐BCFY maintains a dense, continuous electrode–electrolyte contact without visible cracking, highlighting the thermo‐mechanical advantage of the dual‐phase design.

3.5. Single‐Cell Performance and DFT Calculation

To assess the benefits of the dual‐phase heterostructure in single‐cell application, NiO‐BZCYYb fuel‐electrode‐supported r‐PCCs were fabricated using BCFY and cc‐BCFY as air electrodes. A dense BZCYYb electrolyte (∼10 µm thick) was employed as the proton‐conducting layer. It should be noted that single‐cell performance can be affected by various factors beyond the air‐electrode heterointerface, including sealing quality, electrolyte thickness, and other contributions from the fuel electrode and electrolyte. Therefore, BCFY‐ and cc‐BCFY‐based cells were fabricated and tested under the same conditions as closely as possible to enable a fair comparison, so as to verify the benefits of rich heterointerface on cell performance. As shown in Figure 5a, the reduced single cell with the cc‐BCFY air electrode exhibits a well‐adhered and crack‐free electrode‐electrolyte interface, indicating good structural integrity after cell fabrication and reduction. The current voltage and power‐density (I‐V‐P) curves under fuel‐cell (FC) operation are compared in Figures 5b and S22. Single cells employing cc‐BCFY deliver peak power densities of approximately 1.44, 0.88, 0.59, and 0.33 W cm⁻² at 650°C, 600°C, 550°C, and 500°C, respectively, all substantially higher than those of BCFY‐based cells. This improved FC performance is consistent with the heterointerface‐related kinetic benefits inferred from the symmetric‐cell comparison, including accelerated ORR kinetics and reduced interfacial polarization. The single‐cell comparison therefore indicates that the electrode‐level advantages associated with the kinetically preserved BCO‐related heterointerface are reflected in practical r‐PCC operation, although it does not independently isolate this effect. Under electrolysis‐cell (EC) operation (Figure 5c; Figure S23), the cc‐BCFY air electrode likewise exhibits enhanced activity. At an applied voltage of 1.3 V, the cc‐BCFY‐based cell achieves an electrolysis current density of approximately −2.47 A cm⁻², significantly exceeding that of the BCFY‐based cell (−1.21 A cm⁻²). This enhanced EC performance is again consistent with the improved kinetics and interfacial charge‐transfer behavior associated with the heterointerfaces, as suggested by the EIS results of symmetric‐cells. Notably, under FC and EC modes, the performance of the cc‐BCFY electrode is comparable to that of representative air electrodes reported recently (Figure 5d,e; Tables S4 and S5).

FIGURE 5.

r‐PCCs performance and durability. (a) Cross‐sectional SEM image of the single cell employing the cc‐BCFY air electrode; (b) I‐V and power‐density curves of single cells with BCFY and cc‐BCFY air electrodes in FC mode at 650°C; (c) I‐V curves and electrolysis current densities of single cells with BCFY and cc‐BCFY air electrodes in EC mode at 650°C; Performance comparisons of single cells in (d) FC mode and (e) EC mode; Long‐term durability tests at 550°C under constant current in (f) EC and (g) FC modes; (h) Voltage‐switching cycling stability of r‐PCC operation conducted at 600°C.

The durability of the cc‐BCFY air electrode was further evaluated under galvanostatic operation at 550°C. As shown in Figure 5f,g, the cell operates stably for over 150 h in EC mode and maintains continuous output for more than 400 h in FC mode at 550°C. Linear fitting of the galvanostatic stability curves gives a voltage change rate of −0.09 mV h⁻¹ in EC mode at −250 mA cm⁻² and a voltage degradation rate of 0.09 mV h⁻¹ in FC mode at 240 mA cm⁻², indicating good operational stability under the tested conditions. To further examine whether the retained secondary phase persists after operation, post‐test TEM‐EDX mapping was performed on the air electrodes after long‐term EC and FC testing (Figure S24). In both cases, Ba/Ce‐enriched nanodomains remain observable after operation, indicating that the BCO‐related secondary phase is still present after the tested FC and EC operation. Together with the in situ XRD stability over 25°C–800°C, these results support the structural persistence of the retained BCO‐related nanodomains within the single cell operation temperature range. For reversible proton ceramic cell operation, current‐switching durability tests were conducted at 600°C and 550°C by alternately cycling the cell voltage between 0.9–1.1 V and 0.8–1.2 V under both FC and EC modes. As shown in Figures 5h and S25, stable responses are maintained over 25 consecutive cycles without noticeable degradation. These results demonstrate that the limited‐calcination‐enabled dual‐phase air electrode enables not only high single‐cell performance but also durable and reversible operation, underscoring its promise for practical r‐PCC applications.

To gain atomistic insight into possible changes in key reaction‐related processes, density functional theory (DFT) calculations were carried out for BCFY and cc‐BCFY. Structural models were constructed based on the experimentally identified phases (Figure S26), and proton migration was first examined using the climbing‐image nudged elastic band (CI‐NEB) method (Figure 6a). Given that protons are accommodated as O─H species within the oxygen sublattice, the migration pathway was decomposed into two elementary steps [71]. O─H reorientation at the same oxygen site, followed by proton transfer between neighboring oxygen sites. All energies are referenced to the initial state.

FIGURE 6.

DFT calculation results. (a) Energy profile and migration barrier for proton transport; TDOS and PDOS of (b) BCFY and (c) cc‐BCFY; (d) Descriptor comparison: hydration energy, the lowest oxygen vacancy formation energy, d‐band center, and H₂O/O₂ adsorption energies.

For BCFY, a maximum energy barrier is observed during the O─H reorientation step, with a barrier of 0.546 eV, indicating a substantial kinetic penalty for local proton reorientation. In contrast, the corresponding barrier in cc‐BCFY is reduced to 0.366 eV, representing a decrease of approximately 0.18 eV, suggesting that the limited‐calcination‐enabled heterointerface environment significantly lowers the energetic cost of proton reorientation. For the subsequent proton transfer step, BCFY exhibits almost isoenergetic (∼0.27 eV), whereas cc‐BCFY displays a downhill energy trend from −0.077 to −0.259 eV, indicating enhanced proton transfer ability in the heterostructured environment.

To elucidate the electronic‐structure origin of this kinetic advantage, the total and projected densities of states (TDOS/PDOS) were analyzed, and the d‐band centers of the transition‐metal sites were evaluated (Figure 6b,c). While both systems exhibit substantial transition‐metal 3d contributions near the Fermi level, cc‐BCFY shows a reorganized d‐state distribution accompanied by an upshift of the average d‐band center from −1.62 to −1.37 eV. Site‐resolved analysis (Supporting Information Table S6) further reveals a broader distribution of Co/Fe d‐band centers in cc‐BCFY, with several sites positioned closer to the Fermi level. Such electronic‐structure features generally indicate strengthened coupling between metal d states and adsorbate orbitals, which is favorable for the activation of key intermediates (e.g., O*, OH*, and H₂O*) involved in oxygen and proton‐related reactions [72].

Based on the adsorption sites for water and oxygen (Figure S27), we further computed the hydration energy, the lowest oxygen‐vacancy formation energy, and the adsorption energies of H₂O and O₂. As shown in Figure 6d, cc‐BCFY exhibits a more negative hydration energy (−1.548 eV) than BCFY (−1.225 eV), together with a lower minimum oxygen‐vacancy formation energy (−0.599 eV vs. −0.317 eV). These trends suggest that the local environment in the cc‐BCFY model is more favorable for defect formation and stabilization of protonic hydroxyl species. In addition, stronger adsorption is predicted for both H₂O (−0.974 eV) and O₂ (−0.325 eV) on cc‐BCFY compared with BCFY (−0.723 and −0.263 eV), suggesting more favorable capture and activation of reactants at the surface. Given the complexity of the multicomponent heterostructure and the simplified nature of the computational model, these calculated values should be interpreted as qualitative comparative descriptors. Thus, the DFT results provide atomistic support for the experimentally observed trends in proton migration, hydration‐related behavior, oxygen‐vacancy formation, surface adsorption, and electronic‐structure evolution in cc‐BCFY, rather than serving as a stand‐alone or quantitative proof of the full mechanism.

4. Conclusion

In summary, this work demonstrates that preserving a kinetically‐metastable nanophase through thermally‐limited calcination is effective for engineering high‐performance and durable air electrodes for reversible protonic ceramic cells. Through limited calcination of a Ce‐containing precursor, a metastable BCO nanophase is kinetically retained within a Ce‐containing perovskite matrix, producing a tightly coupled dual‐phase heterostructure with strong interfacial interaction. This heterostructure reshapes local defect chemistry and interfacial reaction environments, leading to reduced Co/Fe oxidation states, enriched oxygen vacancies, enhanced hydration behavior, and accelerated surface oxygen exchange and bulk oxygen transport. As a result, cc‐BCFY delivers a low ASR of 0.38 Ω cm² at 550°C and, in single cells, achieves a peak power density of 1.44 W cm⁻² at 650°C and an electrolysis current density of −2.47 A cm⁻² at 1.3 V, together with stable operation for over 400 h in FC mode and over 150 h in EC mode at 550°C. More broadly, rather than relying solely on the discovery of new triple‐conducting oxides, this work shows that controlling thermal history to kinetically retain metastable, non‐equilibrium nanophases offers an additional design dimension for air electrodes in reversible protonic ceramic cells.

Author Contributions

Y.Z. and B.C. conceived and designed the project. Y.P., H.L., K.L., and L.F. carried out materials synthesis, electrochemical measurements, and structural characterizations. Z.L. performed advanced electrochemical analysis and contributed to data interpretation. Y.P., J.L., R.D., S.T., and H.Z. conducted microscopic characterizations and interfacial analysis. B.C. and H.X. supervised the project. Y.P., Z.L., and Y.Z. drafted the manuscript, with critical revisions from B.C. and Z.S. All authors discussed the results and reviewed the manuscript.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: adma73819‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that support the findings of this study are available in the supplementary material of this article.