Hydration-induced stiffness enabling robust thermal cycling of high temperature fuel cells cathode

Domain: earth_science
Published: 2025-04-02
Editorial outcome: major_revision

Abstract
Thermo-mechanics of cathode is closely related to the durability of high-temperature solid oxide fuel cells (SOFCs), with two main mechanical failures during thermal cycling: interface delamination and bulk cracking of cathode. Bulk cracking, caused by insufficient fracture strength/stiffness is a big concern but often overlooked. Here, we introduce chemical hydration to offset the thermal expansion, enhancing the cathodic mechanical stiffness and fracture strength, thus promoting the thermo-mechanical durability of cathode in proton ceramic fuel cells (PCFCs). Such chemical-induced expansion offset is achieved by strengthening intergranular bonding inside the bulk cathode after the hydration, preventing granule detachment during thermal shrinkage. As a demonstration, the stiffness-enhanced air electrode (BaCo0.7Ce0.15Y0.15O3, noted as s-BCC-Y) exhibits 86% enhancement of fracture strength, thus thermal cycling stability with almost no degradation after 35 harsh thermal cycles between 600 and 300 °C, surpassing pristine BaCo0.7Ce0.3O3 and many cobalt-free PCFC cathodes. Benefitted from the improved stiffness of cathode, full cell with the s-BCC-Y electrode demonstrates enhanced power output. This work highlights the importance of bulk cathode thermo-mechanics in developing robust SOFCs for high temperature energy applications.

Introduction
Solid oxide fuel cells (SOFCs) can directly convert the chemical energy of fuels into electrical energy. Their high energy conversion efficiency and environmental friendliness align well with current trends in energy markets for higher sustainability. At high operating temperatures (i.e., above 600 °C), their overall efficiency could exceed 80%. Unfortunately, the high-temperature operation brings challenges to the industrialization of SOFCs, including multiple degradation issues at levels and scales from particle to system, especially during thermal cycling1,2. In particular, the performance degradation of air electrodes (also called cathodes) is largely attributed to mechanical failures including interface delamination and bulk cracking3. Early studies have been aimed at solving the interface delamination issue, through the reduction of the thermal expansion coefficient (TEC) of cathodes to match with the rigid electrolyte4. For example, developing low Co or Co-free cathode5, compositing with the electrolyte powder6–10, or negative thermal expansive addictive3,11 were all proven to be effective (description in details in SI-1). However, merely reducing the TEC may not fulfill all the thermal-mechanical requirements of a robust cathode since the bulk crack issue of the cathode remains.

Principally, the porous cathode is the most vulnerable layer to mechanical cracks in the membrane-electrode assembly (MEA) of SOFC due to its low fracture strength, originating from its high grain boundary embrittlement. In an anode-supported SOFC, the cathode layer is normally prepared by co-firing with anode substrate at ~900–1200 °C, lower than the sintering temperature of the dense electrolyte layer alone (>1300 °C). The contact bonding of cathode grains is therefore substantially lower in strength than that of the electrolyte layer or the anode later. For example, the classic fracture strength of 8YSZ12 and BaZr0.1Ce0.7Y0.1Yb0.1O3-δ (BZCYYb) is respectively 145 and 84.43 MPa (measured in this work), much higher than that of bench-marking cathodes, i.e., 34 MPa for BSCF13 and 40 MPa for LSCF14. Therefore, during both cathode layer fabrication (i.e., cooling process after the firing) and operation (i.e., thermal cycling between room temperature and working temperature), the bulk cracks are more likely to occur in cathode layer than other layers, when the cell is subjected to internal thermal stress and external mechanical loading. Particularly in practical operation of large-size SOFC in stacks, large compression load is needed to ensure good current collecting, which could cause cracks in cathode due to the press-in of interconnectors15. To solve the bulk crack issue, the stiffness of cathode (measured by the modulus and hardness) and mechanical strength that are jointly determined by the grain contacting strength and mechanical expansion behavior, need to be improved16,17. But unfortunately, this has been long-time overlooked.

Here, we took the state-of-art triple conducting BaCo0.7Ce0.3O3 (BCC)18,19 as a parent perovskite material for the development of cathode, which was selected because of it high oxygen reduction reaction (ORR) activity, but it also unfavorably high TEC and low bulk stiffness and fracture strength (~28 MPa). We further proposed a chemical-induced expansion offset strategy to facilely improving the stiffness and anti-crack capability of the cathode by doping Y at the B site of pristine ABO3 structure. Characterizations indicate that the chemical-induced expansion offset originates from the Y-doping enhanced hydration of the oxide lattice, thus strengthening the intergranular bonding inside the bulk cathode by counteracting the detaching of cathode granules during the thermal shrinkage (from calcination temperature to room temperature). As a result, the stiffness-improved BaCo0.7Ce0.15Y0.15O3 (s-BCC-Y) demonstrated better material mechanical and electrochemical performance.

Results
Hydration-induced chemical expansion concept
As known, the cathode experiences significant thermal stress due to the mismatch of the TEC with the electrolyte materials, especially during temperature cycling. This thermal stress leads to interface delamination and cathode fracture. Numerical studies (SI-1) have focused on optimizing interface delamination from the perspective of TEC, while the issue of cathode fracture has been largely neglected for a long time. Figure 1a illustrates that at the calcination temperature, the cathode could be free of thermal stress with all granules well connected with neighbors. During cooling, the thermal shrinkage would cause intergranular faults, deteriorating the stiffness and thus causing cathode fracture. Supposing additional chemical expansion during this cooling process, the faults would be alleviated and the stiffness can be restored—denoted as the concept of “hydration-induced chemical expansion offset”.

To practice this concept, we used Y doping to create the desired chemical induced expansion and study its effect on the stiffness of traditional cobalt-containing perovskite since it was reported that the substitution of trivalent Y3+ in the B site of perovskite generally introduced significant chemical expansion (i.e., by about 73.7% from 0.019 to 0.033 expansion in BaZr0.95Y0.05O2.975 and BaZr0.8Y0.2O2.9)20. Moreover, possibly increasing oxygen vacancy, which means that oxygen ions have more pathways for transport, which facilitates the migration of oxygen ions and the reduction reaction, thereby enhancing the catalytic activity of the SOFC cathode21–23. The BaCo0.7Ce0.3O3 18,19 was selected as the parent perovskite as it has promising ORR/(oxygen evolution reaction) OER activity due to its high triple-conducing properties and stability in water-containing atmosphere, while its TEC is still too high to maintain good thermo-mechanical robustness. As schemed in Fig. 1b at the granule/lattice level, Y3+ doped BCC (s-BCC-Y) perovskite is expected to generate more oxygen vacancies for enhanced hydration reactions, which would expand the lattice during cooling since the hydration is thermodynamically exothermic20,24. This lattice expansion is induced chemically by Y doping and could counteract the generation of intergranular faults.

To verify this, we first prepared s-BCC-Y (BaCo0.7Ce0.15Y0.15O3) and BCC by a facile sol-gel method by calcination at 1000 °C in air. X-ray diffraction (XRD) patterns of the parent BCC and as-obtained s-BCC-Y were observed at room temperature, as shown in Fig. 1c, d. As expected, the s-BCC-Y material was self-assembled into two crystalline phases, a major cubic perovskite phase (89.7%, Pm-3m, BaCo0.74Ce0.094Y0.166O3) and a minor rhombohedral phase (10.3%, R-3c, BaCo0.11Ce0.86Y0.03O3) based on the Rietveld refinement (details shown in SI-2). In addition, lattice spacings were measured from this selected region, BaCoO3 (102) and BaCeO3 (110) in BCC, Pm-3m (100) and R-3c (110) in s-BCC-Y were clearly observed, in good accordance with the XRD refinement results, as shown in Fig. 1e, f. These refinement results correspond well with the interplanar spacings calculated from Selected area electron diffraction (SAED) pattern, and the Energy-dispersive X-ray (EDX) analysis also confirmed the presence of Co-rich and Ce-rich phases (shown in SI-3). To better present how Y affected these two phases, we conducted DFT calculations and performed iterative self-consistent cycles for structural optimization (mainly focusing on bond strength and angles). All structures were fully relaxed until the maximum force on each atom was less than 0.01 eV/Å, and the energy convergence criterion was set to 10–8 eV (as shown in SI-4). More importantly, it can be observed in the high resolution TEM images that the introduction of Y elements assuredly reduced the intergranular fault regions between grains, resulting in a thinner intergranular amorphous phase at grain boundaries, which was reported to be not only beneficial for both proton conduction and phase stability against decomposition at the same time25, but also for enhancing the material’s stiffness26,27.

Based on the results above, we can confirm the successful substitution of Y in the positions of Co and Ce in the B-site of BCC. Due to the larger ion radius of Y3+ than Ce4+ (0.9 Å vs. 0.87 Å)28, doping Y into the lattice of BCC at the synthesis temperature would create lattice expansion and more lattice oxygen vacancies, if then cooled down to room temperature, the created lattice oxygen vacancy could bring in more hydration, forming proton defects and therefore realizing the mitigation of thermal shrinking due to refill of oxygen vacancies by hydration as already widely verified in proton conducting perovskite24,29,30. To further validate the enhanced hydration capacity introduced by Y doping in BCC, we compared the weight changes of BCC and s-BCC-Y from 30 to 1000 °C in dry air, as shown in the TG and DTG analyses in Fig. 1g. Beyond the typical water loss in the 100–200 °C range, commonly observed in porous materials, both BCC and s-BCC-Y exhibit notable mass losses from 400 to 1000 °C, particularly around 450 °C and 750 °C. This behavior aligns with that of many cathode materials31,32. Extensive studies33–35 have attributed this phenomenon primarily to the loss of lattice oxygen in perovskite cathodes, which is an indication of hydration. Therefore, the observed enhanced weight loss at 450 °C could be caused by the enhanced hydration in Y doped s-BCC-Y. Temperature-programmed desorption (TPD) tests were also conducted in water and oxygen environments to confirm the TG results. The results shown in Fig. 1h (H2O-TPD) and SI-5 (O2-TPD), indicate that the mass loss in BCC in TG is mainly due to loss of lattice oxygen, while s-BCC-Y’ mass loss not only contributes to lattice oxygen but a significant amount of water. This observation again highlights the superior hydration capability of s-BCC-Y compared to BCC.

Reduction of TEC and interface optimization
The hydration-induced expansion is expected to mitigate the dilatometric response, as evidenced by the experimental curves of bar-shaped samples in ambient air with a slight amount of moisture. As shown in Fig. 2a. Y doping reduced the TEC of BCC by 34.1% (BCC: 25.89×10-6 K-1, s-BCC-Y: 17.06×10-6 K-1, in the range of 300 to 800 °C). The calculation formula for TEC and the specific experimental test data are depicted in SI-6, and the atomic coordinates of the optimized computational models were provided in our Supplementary Data 1. The ab initio Car-Pareinello molecular dynamics (AIMD) in the framework of the density functional theory (DFT)36 was further used to validate the TEC mitigation, as marked by dotted values in Fig. 2a, where the main Co-rich phase (BaCo0.74Ce0.094Y0.166O3) of s-BCC-Y composite was simplified to BaCo0.7Y0.15O3 for a reasonable DFT calculation. The simulation expansion behavior of lattice after Y doping shows smaller changes in the lattice constant upon temperature in case (see Table S6). This reduced variation can be attributed to the influence of Y on the outer electrons of Co atoms and the strengthening of the Co-O bond energy37,38. Interestingly, the simulated TECs are slightly smaller than experimental values, especially for s-BCC-Y below 600 °C. Considering the simulation process of TECs by DFT calculation does not account for the hydration-induced chemical expansion, we attribute this discrepancy to the chemical expansion that offsets the thermal shrinkage when lowering the temperature. To further explain the reduction of TEC, the partial density of states (PDOS) of cubic perovskite BaCoO3 and BaCo0.7Y0.15O3 as model systems were calculated, which were commonly used to analyze changes in cathode electrochemical performance39, as shown in Fig. 2b. Co ions show a considerable density of states around the Fermi level, indicating Co4+ can easily transform into Co3+ 40. Specifically, a distinct resonance peak of Co and O elements was observed above the Fermi level, which means electron clouds of Co and O further overlap and suggests that the Co-O equilibrium interatomic distance (the bond length)41 turns to be shorter. As is commonly recognized, a decrease in bond length is accompanied by an increase in bond energy42–44. The observed increase of Co-O bond energy after the Y doping indicates an improved resistance to thermal lattice expansion, thereby reducing the overall TEC. This trend also conforms with the calculation data obtained from DFT (SI-6, Table S6). In addition, we also validated these PDOS results by XPS characterizations. The fitted XPS plots of Co 2p3/2 in the BCC and s-BCC-Y are shown in Fig. 2c. In BCC, the Co3+ and Co4+ are 46.66% and 53.34%, while they are 67.13% and 32.87% for s-BCC-Y, respectively. This not only confirms the analysis results from DFT, which indicates Co4+ can more easily transform into Co3+ but also indicates that the introduction of Y lowers the overall valence state of BCC. It is conducive to the formation of oxygen vacancies to enhance the ORR activity45–48. The Ce-rich phase was further characterized using PDOS from DFT calculations and XPS, as shown in SI-7. The PDOS results indicate that Ce’s electronic states cross the Fermi level, suggesting enhanced conductivity, consistent with previous reports on the high conductivity of yttrium-doped barium zirconate49. XPS analysis further reveals a reduction in Ce’s chemical valence, which correlates with the observed decrease in Co’s chemical valence.

Although the TEC of this cathode has been significantly reduced, it still like many other excellent cathode materials, such as BSCF50 (19.95×10-6 K-1, 50–1000 °C) and BCFZY51 (18.38×10-6 K-1, 400–800 °C), when compared to the electrolyte material (9–12×10-6 K-1). In practical engineering, TEC is only one of the important factors affecting interfacial delamination, and it is influenced by many other factors such as the thickness, porosity, and hydration level of the cathode. To quantify the beneficial effect of hydration-induced TEC mitigation, as well as the specific interfacial optimization effect brought by this degree of TEC reduction. Numerical simulations were performed to probe the interface thermal stress at a typical operating temperature (600 °C) using the Finite Element Method for a 2D-axisymmetrical model of the porous cathode coated on a dense proton conducting electrolyte (BZCYYb) based on reconstructed geometries from cross sectional SEM images. The details of simulation parameters were shown in SI-8. The simulated von Mises stress at the cathode-electrolyte interface revealed a clear decrease due to the mitigation of TEC. This reduction is particularly pronounced at the cathode edges, where thermal stress is concentrated, as shown in Fig. 2d. The 34.1% reduction in TEC leads to a 49.8% decrease in von Mises stress at the interface edges (from 4.94 to 2.48 GPa). These simulated results suggest that hydration-induced TEC mitigation can significantly reduce the risk of interfacial delamination between the cathode and electrolyte.

Hydration-enhanced bulk stiffness at a micro-scale
As already stated in the main, delamination is not the only cause for mechanical failure of the cathode, and bulk cracking also plays a significant role52,53. Particularly at high temperature, the cathode materials are more susceptible to damage and cracking due to the thermal-reduced stiffness54,55. Here we compared the bulk stiffness and fracture strength of the hydration-enhanced cathode after the Y doping. First, according to the definition of stiffness, Young’s modulus is viewed as a key parameter to characterize it16,17. Therefore, high-temperature nanoindentation tests were conducted to measure the Young’s modulus and hardness of the cathode. A Berkovich indenter was used for the nanoindentation test at a displacement control mode at varied temperatures (Fig. 3a), and the load-displacement curves at maximum depth to 1500 nm are shown in Fig. 3b. s-BCC-Y exhibited better resistance to the indenter intrusion. Specifically, at 25 °C, s-BCC-Y showed a maximum load of 147 µN, while BCC only displayed 98 µN. As the temperature rises to 600 °C, s-BCC-Y maintained a maximum load of 80 µN compared to 35 µN for BCC. The most commonly used analysis method for obtaining hardness and modulus was developed by Oliver and Pharr56,57. With their method, Young’s modulus and hardness were calculated, and the results were shown in Fig. 3c and the test curves were presented in SI-9. As expected, similar to the most ceramic perovskites, both Young’s modulus and hardness decreased with increasing temperature58,59. As expected, s-BCC-Y always exhibited superior values, which indicates enhanced stiffness at all temperatures than BCC.

For a more detailed mechanical assessment, ultra-precision pressure sensors and Cube corner probe were used in 100 nm indentation, shown in Fig. 3d, e. Different from the Berkovich probe, which is often conducted to calculate Young’s modulus and hardness of bulk samples60,61, this Cube corner probe with a smaller half angle and a 40 nm tip curvature radius can be used to assess the stiffness of individual grains without the interference of grain boundaries. In this 100 nm indentation test, s-BCC-Y achieved a maximum load of 360.6 µN, higher than BCC (280.1 µN), further indicating that the Y-doping enhanced the stiffness of individual grain in a lattice-scale condition.

The modulus of the Co-rich (BaCoO3) and Ce-rich (BaCeO3) structures, as well as their Y-doped variants were also simulated by DFT calculations. The modulus calculations were conducted using VASP, with plane-wave basis sets and projector augmented-wave (PAW) pseudopotentials. Structural optimizations were performed until the maximum atomic force was reduced to less than 0.01 eV/Å, and the energy convergence criterion was set to 10-8 eV. Seven different strains were applied, ranging from −0.015 to 0.015 with an increment of 0.005. VASPKIT was utilized for data extraction and post-processing of the calculated results. All Voigt62, Reuss63, and Hill64 methods were used, and the computational results are presented in (SI-10). The Young’s modulus, bulk modulus, and shear modulus of the Co-rich phase in BCC and s-BCC-Y are shown in Fig. 3f, which indicates the modulus increases after the Y doping, consistent with the results of nanoindentation tests. The Pugh’s ratio (k=G/B) is often used to evaluate the quasi-ductile/brittle nature of ceramics65,66. When G/B is less than 0.571, it indicates a ductile property of the ceramic67–69. Both BCC and s-BCC-Y met Pugh’s criterion, and s-BCC-Y (k=0.511) had a lower k than BCC (k=0.565). This suggests that s-BCC-Y has exhibited certain quasi-ductility characteristics, making it resistant to fracture. It is also consistent with the cathode fracture experiments we conducted next in Fig. 4.

Hydration-improved fracture stiffness at a macro-scale
To verify the positive effect of bulk stiffness enhancement on fracture strength, we examined the fracture strength of dense BCC and s-BCC-Y pellets using a ball-on-ring fixture, a more reliable method for testing the strength of flat brittle specimens than traditional 3-point bending test70, as depicted in Fig. 4a. The specific calculations for strength were shown in the method section (details in SI-11). Fracture strength from room temperature to 700 °C of both BCC and s-BCC-Y are illustrated in Fig. 4b that the highest value of 52.1 MPa was achieved for s-BCC-Y at room temperature, over 80% higher than that of pristine BCC (28 MPa). Interestingly, the enhancement in fracture strength was more pronounced at lower temperatures, while it became negligible at higher temperatures (>600 °C). We attribute this phenomenon to chemical hydration, which occurs predominantly below the hydration temperature turning point. As is well known, perovskite-based cathode materials undergo hydration during the cooling process24, which can optimize grain contact interfaces and, in turn, increase material stiffness.

This trend coincides well with the discrepancy of TEC curves below 600 °C in Fig. 2a, implying that the enhanced hydration indeed altered the thermal mechanical behaviors, in term of both TEC and fracture strength. Accordingly in Fig. 4c, the load vs. displacement curves of BCC and s-BCC-Y specimen till fracture were both further enhanced if treated with a higher 10 vol.%-H2O wetted air at 800 °C for 5h with subsequent cooling. All these experimental findings indicate that enhanced hydration resulted in the increase of fracture strength in s-BCC-Y.

To further elucidate the universality of the strategy of hydration induced for mechanical enhancement, we also compared the fracture strength of the classic single cubic phase cathode Ba0.5Sr0.5Co0.8Fe0.2O3 (BSCF) and B-site 10% Y doped Ba0.5Sr0.5(Co0.8Fe0.2)0.9Y0.1O3 (noted as BSCFY with XRD profile shown in SI-12), before and after the full hydration (800 °C, 10 vol.%-H2O in air, 5h), which were shown in Fig. 4d. As it presented, Y-doping effectively enhanced the fracture strength of BSCF, both before and after full hydration, which means hydration is not limited to the composite BCC system but also applies to single-phase cubic perovskites. However, this improvement is less significant compared to that observed in BCC. There might be two reasons for that. First, the high fraction of Co and Fe in BSCF already results in high fracture of oxygen vacancy that the further doping of Y does not show significant improvement of hydration capacity, Second, the high reducibility of Co3/4+ and Fe3/4+ ions (absent in valence-fixed Y3+) at high temperature is the origin of its large TEC (>20×10-6 K-1), and thus large thermal stress between grains. That implies the high fraction of Co/Fe in BSCF than BCC could dominate the fracture strength, other than the hydration capacity. Besides, we also doped the Yb element into BCC (nominally BaCo0.7Ce0.15Yb0.15O3, noted as BCCYb with XRD profile shown in SI-12), which is also a trivalent dopant, creating oxygen vacancies for hydration, similar to the Y element. The BCCYb also showed improved strength than BCC (see Fig. 4d), further confirming that the enhanced hydration due to trivalent doping can enhance fracture strength.

To get further direct evidence of the improved stiffness for anti-crack formation, FIB-SEM was used to scrutinize the internal microstructures of BCC and s-BCC-Y by ion milling as depicted in Fig. 4e. After sintering at 1150 °C for 5h, the samples were naturally cooled to room temperature and then carefully transferred to the electron microscope for observation. There was an obvious distinction between BCC and s-BCC-Y. BCC displayed clear cracks along its phase boundaries and showed a tendency to propagate from their boundaries towards the interior of the grains. However, the s-BCC-Y showcased intimate intergranular contact, and almost no cracks were observed between the grain boundaries. Besides, a rapid quenching method was conducted to examine their crack resistance. Dense BCC and s-BCC-Y samples were quenched from high temperature (1150 °C) to RT, by direct taking the samples from furnace to the water (RT). which subjects the samples to rapid shrinkage and induces possible thermal cracks. As shown in the following surface (SI-13, Fig. S17), the cracks formed in the BCC sample were much more severe than the Y-doped samples, demonstrating the superior thermal shrinkage resistance of s-BCC-Y. This also correlates well with the interfacial observations at the lattice scale by TEM (in Fig. 1e, f). This difference of microstructures is solid evidence of the mechanism of stiffness enhancement as depicted in Fig. 1a due to the enlargement of the perovskite structure after Y-doping, leading to grain expansion countering the thermal shrinkage when cooled, which results in a closer contact between different grains.

As a short summary of the above mechanical assessment, a noticeable corelationship between enhancement in stiffness/fracture strength and hydration successfully demonstrated the effectiveness of facile Y-doping for developing mechanical robust cathode of SOFCs.

Electrochemical evaluation of hydrated s-BCC-Y during thermal cycling
The electrochemical activity and thermal-mechanical durability of s-BCC-Y as an oxygen-reduction electrode were evaluated. First, electrochemical impendence spectroscopy (EIS) measurements of BCC and s-BCC-Y symmetrical cells with BZCYYb electrolyte were performed under open-circuit conditions from 500 to 700 °C in ambient air atmosphere. The Arrhenius plots and Nyquist plots of the area-specific resistance (ASR, noted as Rp) of BCC and s-BCC-Y were compared in SI-14, Fig. S18. s-BCC-Y demonstrated consistently lower polarization resistance at all temperatures than BCC. For example, the Rp of s-BCC-Y was only 0.23 Ω cm2, 28.1% lower than that of BCC (0.32 Ω cm2). The DRT analysis (SI-14, Fig. S18) shows the differences between BCC and s-BCC-Y are mainly observed at mid- and high-frequencies, suggesting that s-BCC-Y has a better oxygen ion conductivity and redox catalytic activity. The stability of the Rp of symmetric cells at OCV was also tested. After 300 h, the Rp of BCC increased from 0.325 Ω cm² to 0.71 Ω cm², showing a degradation of 1.28 mΩ cm²/h, while the decay of s-BCC-Y’s is only 0.58 mΩ cm²/h (SI-14, Fig. S19). More importantly, s-BCC-Y demonstrated a much improved thermal cyclability of Rp, as shown in Fig. 5a. After 35 cycles (for ~120 h) of rapid temperature cycling between 300-600 °C, the Rp value of the s-BCC-Y electrode only slightly increased from 0.23 to 0.33 Ω cm2, while BCC increased from 0.34 to 0.59 Ω cm2 by almost 2 times (corresponding Nyquist plots and DRT shown in SI-14, Fig. S20). DRT analysis reveals that the main difference between the two occurs at low-frequency, indicating that slow processes, such as gas diffusion within the BCC cathode, begin to change, likely due to interfacial delamination or electrode fracture. In contrast, s-BCC-Y shows smaller changes at low frequencies. Besides, a long-term stable cathode also need to have good CO2 resistance71, and s-BCC-Y exhibited that better (see SI-15, Fig. S21). DRT results show significant enlargement of peaks in the low-frequency region of s-BCC-Y and BCC, which indicates that gas diffusion is suppressed, likely due to cathode surface carbonization and the coverage of carbonates impeding oxygen diffusion.

We examined the SEM morphology of BCC and s-BCC-Y before and post-cycling (Fig. 5b). As anticipated, not only delamination but also bulk crack (leads to the loss of partial cathode in the bulk) were observed in the BCC sample post-cycling. As illustrated in the schematic diagram (Fig. 5b), we attribute the extensive cathode loss observed after long-term cycling to fractures caused by thermal stress at the interface of a low-stiffness cathode. By contrast, the improved stiffness and fracture strength enabled s-BCC-Y to retain its intact porous framework without any delamination and bulk cracks, which is key to superior durability. In addition, it should be mentioned that no reactions were observed for both air electrodes and BZCYYb (SI-16, Fig. S24), Further confirming that the cracking and delamination observed at the interface were mainly caused by poor stiffness and thermo-mechanical properties in this work, rather than chemical side reactions.

To fully demonstrate the concept of chemically induced expansion offset as a general method for thermal-mechanical enhancement, we further tested s-BCC-Y as the cathode in both oxygen ion and proton-conducting single cells (SOFC and PCFC), using H2 as the fuel. In Fig. 5c, d, the s-BCC-Y cathode reaches 1220 mW cm–2 at 800 °C, using NiO-YSZ as anode and YSZ as electrolyte, greatly outperforming the BCC cathode (only 603 mW cm–2 at 800 °C). On the NiO-BZCYYb|BZCYYb half-cell, s-BCC-Y also showed a better power density than BCC (504 mW cm–2 vs. 258 mW cm–2 at 500 °C, This more significant improvement of performance in PCFC than SOFC can be explained by the more harsh thermal mechanical environment in PCFC than SOFC, that was resulted by the lower TEC of BZCYYb than YSZ72,73. Meanwhile, the long-term stability of single cells was also assessed. Both cells were able to sustain a consistent power output over an extended duration. In Fig. 5e, the cell output shows a voltage drop from 0.797 V to 0.758 V after 380 hours, corresponding to a 0.103 mV/h degradation. In Fig. 5f, the voltage decreases from 0.753 V to 0.71 V, representing a 0.153 mV/h degradation. This strongly suggests that the s-BCC-Y composite cathode, improved by stiffness enhancement, exhibits remarkable performance in single fuel cells. This also indicates that the cathode, with enhanced hydration demonstrates good durability in both SOFC and PCFC configurations.

Finally, we further assessed the performance of s-BCC-Y on large-scale PSOFC (NiO-BZCYYb|BZCYYb|s-BCC-Y) with an active cathode area of 9.62 cm2, that is above 35 times higher than button cell testing (~0.28 cm2) as compared in Fig. 5g. Considering cathode with a larger area is more vulnerable to cracking and delamination failures caused by the higher thermal stress in the multi-layered film due to geometry effect and strong local joule heating effect, the high large single cell performance of s-BCC-Y (a total power of 3.5 W at 650 °C, corresponding to a practical power density of ~0.36 W/cm-2) again confirms its good thermo-mechanical robustness.

Methods
Powder preparation
All perovskite materials in this paper, including BaCo0.7Ce0.3O3-delta, BaCo0.7Ce0.15Y0.15O3-delta, BZCYYb, YSZ were prepared via a complexing sol-gel method. For instance, when preparing BaCo0.7Ce0.15Y0.15O3-delta, stoichiometric amounts of Ba(NO3)2-99.5%, Co(NO3)2·6H2O-99.9%, Ce(NO3)3·6H2O-99.95% and Y(NO3)3·6H2O-99.99% according to the nominal composition, were dissolved into water and heated at 180 °C under stirring. Ethylenediaminetetraacetic acid (EDTA-99.99%) and citric acid (CA-99.5%) were added to act as complexing agents at a molar ratio of 1:1:2 for total metal ions, EDTA and CA. NH3·H2O was used to adjust the PH value of water in the solution to around 7. A sol-gel precursor was obtained after the removal of water in the solution by evaporation, which was pre-decomposed at 180 °C for 8 h to obtain the precursors. The as-synthesized precursors were then calcined at 1000 °C for 5 h in air to obtain the final product. The SEM is shown in SI-17. All materials mentioned above were procured from Aladdin Scientific Corporation.

Fabrication of symmetric cells and single cells
Symmetrical cells with the configuration of cathode|BZCYYb|cathode were fabricated for the electrochemical impedance spectroscopy (EIS) test. First, the as-synthesized BZCYYb powders were drily pressed into pellets and calcined at 1450 °C to obtain dense pellets, placed in a normal environment for standby use (~25 °C, 60% relative humidity). Next, the cathode suspension, consisting of the 1g as-synthesized cathode materials, 0.8 ml glycerol, 2 ml ethylene glycol, and 10 ml isopropyl alcohol, was prepared by ball milling for 40 min. The suspension was sprayed onto both sides of BZCYYb disks and calcined at 800 °C for 2 h in air. On both sides of the cathode, silver paste was applied to create silver grids for the current collection, and silver wires were used to connect the cell and the electrochemical workstation for the needed test.

NiO+BZCYYb anode-supported single cells with thin-film electrolytes were prepared by dry-pressing and high-temperature co-sintering. Anode powders consisting of 6g NiO, 4g BZCYYb, 1g starch, and appropriate ethanol were prepared by ball milling. Then, manually mixed and left to dry at room temperature. First, 0.35g anode powder was pressed into a pellet with a pressure of 4 MPa, and the 0.02g BZCYYb electrolyte was evenly spread onto the anode surface and co-pressed under a pressure of 8 MPa. Then the bi-layer BZCYYb electrolyte disk was calcined at 1450 °C for 5 h in air. Finally, the cathode suspension was sprayed onto the electrolyte surface with an effective area of 0.27 cm-2 and calcined at 800 °C for 2 h. The final thickness of the electrolyte layer is about 10 μm (See SI-17 for details). Regarding the large area single cell with a cathode area of 9.62 cm2, the anode is prepared by casting processing, followed by spin-coating the electrolyte onto the hot-pressed anode, co-sintering at 1450 °C to obtain the test half-cell and the cathode is sprayed onto the electrolyte surface too. The oxygen ion conducting SOFC with NiO-YSZ|YSZ half cells are purchased from Ningbo SOFCMAN company, with a 400 μm NiO-YSZ anode, 15 μm YSZ electrolyte, 3 μm GDC barrier layer.

Material characterizations
The crystal structures of BCC and s-BCC-Y were carried out by Rigaku Miniflex 600 X-ray diffractometer (XRD) with Cu source. The diffraction patterns were collected in the 2-theta range between 10 and 80° with intervals of 0.02° at a scanning rate of 2° min-1, Rietveld refinement was performed using jade and X’Pert HighScore plus software. The morphology of the cathode particles and fuel cell cross-section were investigated by Thermo Scientific Scios 2 scanning electron microscopy (SEM). A further micro-structure of s-BCC-Y was investigated using transmission electron microscopy (TEM, FEI Talos F200x). energy-dispersive X-ray spectroscopy (EDX) measurements were performed in the same instrument. XPS measurement were conducted with a spectrometer (Thermo Scientific K-Alpha). About 100 mg of cathode materials were pressed into a pellet on a sample holder. The sample was then introduced into the analysis chamber when the pressure was below 2.0×10-7 mbar. The spot size was set to 400 µm, operating at 12 kV and a filament current of 6 mA. Full-spectrum scans employed pass energy of 150 eV with a 1 eV step size, while narrow-scan spectra used a pass energy of 50 eV with a 0.1 eV step size. The experimental data peak fitting was performed using the software Avantage.

Electrochemical measurements
The I-V and I-P curves of cells were obtained using a Keithley 2460 source meter with a four-probe configuration. The anode of a single cell was fed by H2 fuel with a flow rate of 80 ml min-1, while the cathode was exposed to ambient air. The impedance of single and symmetrical cells was investigated by the EIS measurement by the Princeton electrochemical workstation (autocalibration) at open circuit voltage (OCV) conditions. The frequency of the EIS measurement ranged from 0.01 to 100,000 Hz and the single amplitude was 100 mV. During the testing process, the temperature was first raised to the maximum operating condition temperature, and then gradually decreased to conduct the I–V curve tests at subsequent temperatures. Taking the cell performance test in Fig. 5g as an example, after securing the cell in the fixture, H2 fuel with a flow rate of 80 ml min-1 was introduced to the anode side. The temperature was then increased from room temperature to 650 °C at a rate of 10 °C/min, held for 50 °C, with the testing conducted at each operating temperature, until the testing concluded at 450 °C.

Fracture strength testing
The fracture strength tests were conducted using a high-temperature mechanic testing machine (OTF-1500X-S-CR, Hefei Kejing, China). The fixture consists of a spherical punch and a circular sample holder with an internal ring. The sample holder, is made of pressureless sintered SiC, which is suitable for high-temperature testing, and is placed on a high-precision pressure sensor for accurate measurements. Several studies74,75 have used a verified formula method to calculate the fracture strength (σ), the test profile and calculation parameter are shown in Fig. 4a. The formula (1) is used to calculate the fracture strength of the material,

σ = A × F / t2 (1)

where t is the specimen thickness, F is the applied force, and A is a dimensionless factor that depends on the geometry of the specimen and loading ball, the ring diameter, and Poisson’s ratio of the loading ball and test material. The factor A is calculated using Eq. 2 for this test,

A = 3 / 4π × [2(1+υs) × ln(a/b) + (1-υs) × (a2-b2) / 2R2 + (1+υs)] (2)

where υs is the Poisson’s ratio of the test material, a is the radius of the support ring, R is the radius of the test specimen, and b is the contact radius of the loading ball. The contact radius of the loading ball b can be simplified to b=t/3 in this test.

The punch approach speed was 3 mm/min, and the approach threshold (the punch contact with the sample) was set to 0.05 kg. After contact, the approach speed was reduced to 0.1 mm/min until the cathode sample fractured. During high-temperature testing, the chamber’s temperature was increased at a rate of 10 °C /min, in a vacuum environment and cooled by 25 °C circulating water. To eliminate experimental variability, each kind of sample underwent 20 trials at room temperature and 10 trials at high temperature. Data within the (μ-σ, μ+σ) confidence interval, about 65% confidence, were considered valid, and the average of these valid data was accepted in this paper.

Nanoindentation testing
This test was performed using the Bruker Hysitron PI 89 Picoindenter. High-temperature Berkovich Probe (diamond, Half Angle 65.27°, Included Angle 142.30°, C0 24.5, Young’s modulus 1140 GPa, Poisson’s ratio 0.07) and Cube Corner probe (diamond, Half Angle 35.26°, Included Angle 90°, C0 2.598, Young’s modulus 1140 GPa, Poisson’s ratio 0.07) were used to test.

The linear fitting models proposed by Doerner and Nix76 and the exponential fitting models proposed by Oliver and Pharr56,57 are used to calculate equivalent elastic modulus Er,

Er = (π / 2) × S / sqrt(A) (3)

Where S is the initial unloading slope, and A represents the contact area between the probe and the material. The contact area A is determined using the Oliver-Pharr method by relating it to the contact depth hc, where A=F(hc), and hc is given by:

hc = hmax - ε × Pmax / S (4)

Where hmax is the maximum displacement during testing, ε is a geometric parameter depending on the shape of the indenter. Then, Young’s modulus of materials was obtained by:

1/Er = (1-υ1^2) / E1 + (1-υ2^2) / E2 (5)

Where E1 and υ1 are the Young’s modulus and Poisson’s ratio of the tested material, E2 and υ2 are the Young’s modulus and Poisson’s ratio of the testing probe material. The hardness H was calculated by:

H = Pmax / A (6)

where Pmax is the maximum load applied, and A is the contact area between the indenter and the tested material.

The probe approach speed was 150 nm/s, and the approach threshold was 50 µN. Once the probe reached a depth of 1500 nm, it remained at this depth for 10 s, and then the probe withdrew at a speed of 150 nm/s. During high-temperature testing, the chamber’s temperature was increased at a rate of 30 °C/min in a vacuum environment and cooled by 25 °C circulating water. To eliminate experimental variability, each kind of sample underwent 150 trials at room temperature (25 °C) and 100 trials at high temperature (150–700 °C). Data within the (μ-σ, μ+σ) confidence interval, about 65% confidence, were considered valid, and the average of these valid data was accepted in this paper.

Numerical simulations
The interface thermal stress numerical simulation was conducted by COMSOL Multiphysics 5.5, and the model was built by actual SEM porous cathode-dense electrolyte interface characteristics. Mesh independence analysis was conducted for all simulations, and unstructured grids were employed. The detailed process is presented in the SI-5.

DFT computations
The TEC was calculated by Ab initio Car-Pareinello molecular dynamics (AIMD) within the framework of DFT. The exchange-correlation term was determined using the generalized gradient approximation (GGA) within the Beke-Lee-Yang-Parr (BLYP) implementation77,78. Norm-conserving pseudopotentials were used, with s and p as nonlocal orbitals and d as local orbitals. The AIMD calculations were performed using the SIESTA code, which is based on a flexible linear combination of atomic orbitals (LCAO) basis sets79 and boasts efficient Order-N scaling80. A cutoff radius of 250 Ry was employed to ensure that the self-consistent procedure converged to less than 1 meV/atom. During atomic relaxation, a maximum of 25 electronic steps were used for each ionic step. For molecular dynamics, the Parinello-Rahman method81 with variable shape and volume of the unit cell was applied (NPT ensemble). A 3×3×3 supercell was built based on the optimized unit cell in different temperatures. Monkhorst-Pack (2×2×2) was set for Brillouim zone sampling. The temperature range investigated was T=600 K–1000 K, with increments of ΔT=100 K for the SIESTA calculations. Finally, the TEC of the structure was calculated using the Phonopy software through the quasi-harmonic approximation (QHA) method.

The modulus calculations were conducted using VASP, employing plane-wave basis sets and the projector augmented-wave method. We utilized the generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) parametrization82 for the exchange correlation potential. The energy cutoff was set to 500 eV, and the Brillouin-zone integration was performed using a Γ-centered Monkhorst-Pack mesh83 of 3×3×3 for BaCoO3 and 5×5×2 for BaCeO3, as implemented by VASPKIT84. The structural optimizations were carried out until the maximum force on each atom reached less than 0.01 eV/Å, and the energy convergence criterion was set to 10−8 eV. We set seven different strains with a range of −0.015 to 0.015 and a step size of 0.005. The elastic tensor matrices for different structures were calculated, and the Young’s modulus (Y), bulk modulus (B), and shear modulus (G) of the structures were obtained using the Voigt, Reuss, and Hill methods.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript reported effects of Y in BaCo(Ce)O3 (BCC) cathode on mechanical stiffness and stability to thermal cycles of the cell. The content of this manuscript is interesting however following points are not clear.\n1)The mechanism for increase stiffness of BCC by doping Y is not clear. The authors estimate TEC or Toughness by using DFT and so detail discussion on mechanism is required.\n2)The increase in TEC is also observed and supported by DFT calculation. However, the estimated TEC was around 1,7 x10-5 K-1 which is not so much difference from that of non-doped one (BCC) however still much different from YSZ (1.1x10- 5K-1). Therefore, increased performance and stability may not be explained by matching of TEC.\n3) Increased cell performance was analyzed with electrochemical impedance in Fig.S16. The assignments of impedance arcs are required. It seems that diffusion overpotential may be changed by Y doping which is not related with mechanical property of cathode.\n4)Degradation is also observed in Fig.5 e and f and assignment of degradation of performance is also required.\n5) Reaction of BCC is also expected and doped Y may be prevent reaction between cathode and BZCYYb electrolyte. So, suppression of chemical reaction is another reason for increased performance of the cell.\n6) In summary, the reason why only Y can increase the mechanical property of BCC is required in spite of small amount of doped and still large difference in TEC", "reviewer_2": "This paper investigates the mechanical and electrochemical performance of protonic ceramic fuel cells (PCFCs), focusing on enhancing the durability of the cathode through chemical hydration. The authors introduce Y-doping (yttrium doping) into a perovskite-based material, BaCo₀.₇Ce₀.₁₅Y₀.₁₅O₃-δ (s-BCC-Y), to mitigate thermal expansion during high-temperature cycling, thereby enhancing mechanical strength and reducing the risk of cracks and delamination in the cathode. Moreover, the Y-doped cathode demonstrates superior electrochemical performance and long-term stability compared to conventional cathodes, especially under thermal cycling conditions. These results are interesting; however, there are several issues that need to be addressed before the paper can be published, requiring major revisions. Please refer to the comments provided below.\n1. In Figures 1c-f, the author notes that Y-doping causes changes in the crystal structure. However, the precise atomic-level modifications, such as how Y ions are substituted within the lattice and how they affect bond lengths and angles, are not thoroughly investigated. Thus, authors should further examine these crystal structure changes at the atomic scale.\n2. In Figure 1g, a detailed explanation of how Y-doping increases oxygen vacancies is recommended. Additionally, since the authors only present the weight change curve for s-BCC-Y, it may give the impression that s-BCC-Y and BCC behave identically under dry-air conditions. Revising this figure would help prevent potential misunderstandings.\n3. The authors demonstrate that Y-doping influences oxygen vacancy formation. However, the work lacks a detailed analysis of the atomic interactions and the specific role these vacancies play within the lattice (Figures 2b and 2c). Furthermore, since the Y dopant can alter the oxidation states of both Ce and Co, excluding the Ce ion from the DFT calculations is not justified. Authors should conduct a more in-depth analysis and provide a detailed explanation\n4. When comparing the Y-doped and undoped BSCFs in Figure 4b, the difference in the increase in fracture strength from the untreated to the hydrated state is negligible. Regarding this, authors need to explain the detail.\n5. In Figure 4e, the comparison of the two samples using FIB-SEM does not specify whether they underwent any particular treatment prior to analysis. Cracks may have formed during the sampling process, so it is essential to confirm if the samples were exposed to treatments such as rapid heating-cooling cycles. Including data from treated samples would more effectively demonstrate their superior thermal shrinkage resistance and provide a more thorough comparison.\n6. To strengthen the authors' claims, SEM images of BCC and s-BCC-Y before operation should be included alongside the images presented in Figure 5b.", "reviewer_3": "Title: Hydration-induced stiffness enabling robust thermal cycling of high temperature fuel cells cathode\nManuscript No.# NCOMMS-24-47288-T\nThe manuscript details a study on improving the mechanical durability of high-temperature fuel cell cathode materials to prevent degradation. It proposes doping BaCo0.7Ce0.3O3 perovskite with yttrium (Y) to enhance thermal expansion tolerance and mechanical properties. This approach has been validated through structural, mechanical, and electrochemical evaluations, showing increased cathodic stiffness and fracture strength and resulting in better fuel cell performance and stability. The innovation is significant for advancing SOFCs and PCFCs.\nHowever, there are still a few points that need attention before the manuscript can be recommended for acceptance:\n1. The manuscript is dense with technical details and data, which is commendable, but it may benefit from a more streamlined presentation to enhance clarity and readability. Simplifying the more complex sections, or providing clearer explanations, might make the document more accessible to a broader audience.\n2. While the mechanical improvements are clear, the paper could discuss any potential drawbacks of the Y-doping strategy more extensively. For instance, does the doping affect the long-term stability of the material or introduce any other trade-offs in terms of cost or material availability?\n3. It would be interesting to include a comparison to other strategies aimed at solving the same problem. How does this approach compare to others in terms of effectiveness and practicality?\n4. While the findings are promising, the scalability of the solution presented is not extensively discussed. Insights into manufacturing challenges and the economic viability of scaling up the doping process could strengthen the manuscript’s applicability to the industry.\n5. The longevity of the cathode material is crucial for commercial application. The paper could explore the long-term performance of the cathodes under continuous operation, beyond cyclic testing.\n6. In the method section, the terminology \"atropurpureus gel\" appears undefined or nonstandard. Should this refer to a distinct variety of gel, it would be prudent to elucidate its composition or an appropriate amendment to its designation.\n7. Given the complications associated with the performance of 8YSZ electrolytes at intermediate temperatures, how do the introduction of chemical hydration and the improvement in mechanical stiffness and fracture strength, as demonstrated by the stiffness-enhanced air electrode s-BCC-Y, impact the thermo-mechanical durability of cathodes in solid oxide fuel cells (SOFCs), particularly with proton ceramic fuel cells (PCFCs)? Can these strategies be applied to optimize other intermediate-temperature SOFC materials, such as doped ceria?"}, "author_response": "Response to Reviewers\nReviewer #1\nThis manuscript reported effects of Y in BaCo(Ce)O (BCC) cathode on mechanical\n3\nstiffness and stability to thermal cycles of the cell. The content of this manuscript is\ninteresting however following points are not clear.\n[Comment 1]: The mechanism for increase stiffness of BCC by doping Y is not clear.\nThe authors estimate TEC or Toughness by using DFT and so detail discussion on\nmechanism is required.\n[Response 1]: Thank you for your careful review of our work. We acknowledge that the\nexplanation of the mechanism behind the stiffness enhancement in the previous version was\ninsufficiently clear. In the revised manuscript, we have made substantial improvements to clarify\nthis point and facilitate readers’ understanding. Based on our experimental findings and\ncomputational analysis (DFT), we were able to attribute the mechanism of increase in stiffness\nmainly to two key factors: TEC mitigation and the enhanced hydration induced by Y-doping.\nFor the first factor-TEC mitigation, DFT calculations were used to determine the theoretical\nTEC and stiffness. We have expanded the discussion on DFT and the mechanism in this reviewed\nmanuscript, specifically in lines 249-254 (also see below highlighted in RED). Briefly, the presence\nof Y atoms suppresses the change of the lattice constant with temperature due to the influence of Y\non the strengthening of the Co-O bond energy. In detail, Co ions show a considerable density of\nstates around the Fermi level, indicating Co4+ can easily transform into Co3+ [Kuklja et al.,\nM.Physical Chemistry Chemical Physics 15, 5443-5471 (2013)]. Specifically, a distinct resonance\npeak of Co and O elements was observed above the Fermi level, which means electron clouds of Co\nand O further overlap and suggests that the Co-O equilibrium interatomic distance (the bond length)\n[ Pauling, L et al., The Journal of Physical Chemistry 58, 662-666 (1954)] turns to be shorter. As is\ncommonly recognized, a decrease in bond length is accompanied by an increase in bond energy\n[Gordy, W, L et al., The Journal of Chemical Physics 15, 305-310 (1947), Phillips, L et al., Reviews\nof Modern Physics 42, 317 (1970)]. The observed increase of Co-O bond energy (equivalent to bond\nstrength) after the Y doping indicates a suppressed change of the lattice constant (∆L) [Wang, X. et\nal. Physical Chemistry Chemical Physics 20, 17781-17789 (2018)] and thus improving the lattice\nresistance to thermal expansion, thereby reducing the overall TEC.\nFor the second factor-hydration, in addition to the intrinsic TEC mitigation due to the lattice\nstructural modification (i.e. Co-O bond energy as explained above) after Y doping, the Y doping\nmay also cast large influence on the hydration process of perovskite. As known, the hydration\nusually induces chemical expansion [Chen, T. et al. The Journal of Physical Chemistry C 125,\n26216-26228 (2021). Guo S, et al. Advanced Functional Materials, 2304729 (2024)], therefore, it\nalso plays a key role to the TEC and stiffness enhancement. Here, we utilized Y-doping (an acceptor\ndoping strategy) to promote hydration, which significantly contributes to offset the thermal lattice\nshrinkage, and thus the increased stiffness. In the previous version, we did not make this point clear.\n2\nwith explanations added on lines 142-154. Furthermore, in Fig. 4d, we also proved the positive\neffect of hydration on strength/stiffness by cross-comparing the same perovskite before vs. after\nhydration and Y doped vs. un-doped. Detailed descriptions of these findings have been\nsupplemented at lines 292–308.\nWe believe these revisions as follows offer a more comprehensive and clear explanation of the\nmechanisms involved for the stiffness increase by doping Y, enhancing the clarity of the manuscript\nand facilitating a better understanding for the readers.\nFig.1 g, h TG and DTG analysis of hydrated BCC and s-BCC-Y from RT to 1000℃. H O-TPD of\nFig 4. d Comparison of fracture strength before and after hydration for BCC, s-BCC-Y, BSCF,\nBSCFY, and BCCYb cathode materials.\nAt Lines 249-254:\n“The modulus calculations were conducted using VASP, with plane-wave basis sets and projector\naugmented-wave (PAW) pseudopotentials. Structural optimizations were performed until the\nmaximum atomic force was reduced to less than 0.01 eV/Å, and the energy convergence criterion\nwas set to 10-8 eV. Seven different strains were applied, ranging from -0.015 to 0.015 with an\nincrement of 0.005. VASPKIT was utilized for data extraction and post-processing of the calculated\nresults.”\nAt Lines 142-154:\n“To further validate the enhanced hydration capacity introduced by Y doping in BCC, we\n2\ncompared the weight changes of BCC and s-BCC-Y from 30 to 1000 °C in dry air, as shown in\nthe TG and DTG analyses in Fig. 1g. Beyond the typical water loss in the 100–200 °C range,\ncommonly observed in porous materials, both BCC and s-BCC-Y exhibit notable mass losses from\n400 to 1000 °C, particularly around 450 °C and 750 °C. This behavior aligns with that of many\ncathode materials 31,32. Extensive studies 33-35 have attributed this phenomenon primarily to the loss\nof lattice oxygen in perovskite cathodes, which is an indication of hydration. Therefore, the observed\nenhanced weight loss at 450°C could be caused by the enhanced hydration in Y doped s-BCC-Y.\nTemperature-programmed desorption (TPD) tests were also conducted in water and oxygen\nenvironments to confirm the TG results. The results shown in Fig. 1h (H O-TPD) and SI-5 (O -\n2 2\nTPD), indicate that the mass loss in BCC in TG is mainly due to loss of lattice oxygen, while s-\nBCC-Y’ mass loss not only contributes to lattice oxygen but a significant amount of water. This\nobservation again highlights the superior hydration capability of s-BCC-Y compared to BCC.”\nAt Lines 292–308:\n“To further elucidate the universality of the strategy of hydration induced for mechanical\nenhancement, we also compared the fracture strength of the classic single cubic phase cathode\nBa Sr Co Fe O (BSCF) and B-site 10% Y doped Ba Sr (Co Fe ) Y O (noted as\n0.5 0.5 0.8 0.2 3 0.5 0.5 0.8 0.2 0.9 0.1 3\nBSCFY with XRD profile shown in SI-12), before and after the full hydration (800°C, 10 vol.%-\nH O in air, 5 h), which were shown in Fig. 4d. As it presented, Y-doping effectively enhanced the\n2\nfracture strength of BSCF, both before and after full hydration, which means hydration is not limited\nto the composite BCC system but also applies to single-phase cubic perovskites. However, this\nimprovement is less significant compared to that observed in BCC. There might be two reasons for\nthat. First, the high fraction of Co and Fe in BSCF already results in high fracture of oxygen vacancy\nthat the further doping of Y does not show significant improvement of hydration capacity, Second,\nthe high reducibility of Co3/4+ and Fe3/4+ ions (absent in valence-fixed Y3+) at high temperature is\nthe origin of its large TEC (>20 ×10-6 K-1), and thus large thermal stress between grains. That implies\nthe high fraction of Co/Fe in BSCF than BCC could dominate the fracture strength, other than the\nhydration capacity. Besides, we also doped the Yb element into BCC (nominally\nBaCo Ce Yb O , noted as BCCYb with XRD profile shown in SI-12), which is also a trivalent\n0.7 0.15 0.15 3\ndopant, creating oxygen vacancies for hydration, similar to the Y element. The BCCYb also showed\nimproved strength than BCC (see Fig. 4d), further confirming that the enhanced hydration due to\ntrivalent doping can enhance fracture strength.”\n3\n[Comment 2]: The increase in TEC is also observed and supported by DFT calculation.\nHowever, the estimated TEC was around 1.7 x10-5 K-1 which is not so much difference\nfrom that of non-doped one (BCC) however still much different from YSZ (1.1x10-5 K-\n1). Therefore, increased performance and stability may not be explained by matching of\nTEC.\n[Response 2]: Thank you very much for pointing out this issue. As shown in our manuscript, we\nindeed did not achieve perfect matching of TEC between s-BCC-Y and YSZ, the actual TEC of s-\nBCC-Y was only decreased to some extent, from 25.89×10-6 K-1 to 17.06×10-6 K-1 by doping, and\nthere is still a noticeable difference compared to the electrolyte material (YSZ, 11×10-6 K-1).\nHowever, this “imperfect” reduction of TEC may still result in improvement to the mechanical\ndegradation of air electrode (especially delamination) since the porous nature of the air electrode\nwill also aids in reduction of the apparent TEC of the whole porous electrode. In our work, the\nporosity of air electrode is approximately 36.7%, which could help explain the increased stability\nand performance along with the “imperfect” reduction of TEC. On this basis, we also can see many\nother high TEC electrodes have the potential for good durability in the literature such as BSCF\n(19.95×10-6 K-1, 50-1000°C) [ Wei, B. et al. Journal of the European Ceramic Society 26, 2827-\n2832 (2006)] and BCFZY (18.38×10-6 K-1, 400-800°C) [Yu, Z. et al. Advanced Functional Materials\n34, 2309698 (2024)].\nOn the other hand, we would like to clarify that we do not attribute the improvement in\nperformance and stability solely to the reduction in TEC. The optimization of the actual interface\ninvolves many factors, such as modulus, hydration, etc. The cathode thickness and porosity also\nplay a role in influencing the true mismatch and stress at the interface.\nConsidering this complexity of correlation between multiple mechanical factors and durability,\nwe have performed numerical simulations to specifically evaluate the effects of the TEC reduction\nfrom 25.89×10-6 K-1 to 17.06×10-6 K-1 on the interface stress. The results show that the maximum\ninterface stress is reduced by 49.8%, which is a significant improvement, that draw us to the\nconclusion that the improved mechanical properties (not only TEC, but also modulus, hardness and\nPugh ratio) results in the improved performance and durability.\nTherefore, we have made a clarification in our manuscript to offer readers a clearer understanding\nat Line 197-203:\n“Although the TEC of this cathode has been significantly reduced, it still large like many other\nexcellent cathode materials, such as BSCF (19.95×10-6 K-1, 50-1000°C) and BCFZY (18.38×10-6\nK-1, 400-800°C), when compared to the electrolyte material (9-12×10-6 K-1). In practical, TEC is\nonly one of the important factors affecting interfacial delamination, and it is influenced by many\nother factors such as the thickness, porosity, and hydration level of the cathode. To quantify the\nbeneficial effect of hydration-induced TEC mitigation, as well as the specific interfacial\noptimization effect brought by this degree of TEC reduction.”\n4\n[Comment 3]: Increased cell performance was analyzed with electrochemical\nimpedance in Fig.S16. The assignments of impedance arcs are required. It seems that\ndiffusion overpotential may be changed by Y doping which is not related with\nmechanical property of cathode.\n[Response 3]: Thank you for pointing out this issue. We appreciate your suggestion, and we\nrealize that assignments of the impedance arcs and more detailed descriptions would be beneficial\nfor the readers. As a response, we have performed a DRT analysis (which is a more “precise” tool\nto analysis impedance data than arcs assignment by fitting) on all the impedance spectra presented\nin the manuscript. We hope this will offer a deeper investigation into the potential mechanisms\nbehind the observed impedance changes.\nWe believe that the changes observed in the low-frequency region are related to diffusion.\nSpecifically, s-BCC-Y maintains stable diffusion properties under thermal cycling, whereas BCC\ndoes not. However, the primary impedance contributions of both electrodes do not originate from\nthe low-frequency diffusion region, as diffusion is not the rate-limiting process in this context.\nThe observed differences in low-frequency behavior (which are very minor) may not be directly\nrelated to the mechanical properties of the electrodes. Mechanical strength determines the\noccurrence of mechanical failure, which in turn affects the three-phase conductive network, thus\ninfluencing the middle-to-high frequency impedance, as observed in our experiments (see Fig. S18).\nWe hypothesize that the difference in low-frequency behavior between BCC and s-BCC-Y,\nparticularly after thermal cycling (see Fig. S20), may be indirectly related to the overall\nmicrostructural changes in the electrodes. Thermal cycling could cause delamination or other\nmicrostructural alterations, disrupting the ionic and electronic conductivity networks, which reduces\nthe effective reaction area. This could lead to localized high current density and, consequently, an\nincrease in low-frequency impedance. As for the primary factors influencing low-frequency\nimpedance, we have conducted preliminary studies on CO poisoning, which significantly impacts\n2\nlow-frequency impedance by forming carbonate layers that hinder gas diffusion (see Fig. S22).\nHowever, this is only a preliminary analysis and an explanation for the minor changes in diffusion\noverpotential. Since we lack direct evidence and a more detailed study, and because low-frequency\nimpedance is not the dominant factor in our investigation, we have chosen not to discuss it in detail\nin the main text.\nAs for the specific revision:\nFirst, we have added Nyquist plots (SI-14, Fig. S18c) and DRT analyses at three different\ntemperatures (550 °C, 600 °C, and 650 °C) in Fig. S18d-f to complement the limited information\nprovided by the Arrhenius plots. Additionally, we have included DRT data analyses of EIS Nyquist\nplots under cyclic testing, as shown in SI-14, Fig. S20.\nAs known, DRT is widely used for fuel cells to analyze reaction processes based on the cell’s\nresponse at different frequencies. In SOFCs, the response at low frequencies(<10Hz) represents\nslow processes such as gas diffusion. The response at mid-frequencies (10-1000Hz) reflects the\nelectrochemical kinetics and the rate of redox reactions, while high-frequency (>1000Hz) responses\nare associated with fast processes like ion diffusion and electron transport. In this work, we define\nlow frequencies as <10 Hz, mid-frequencies as 10–1000 Hz, and high frequencies as >1000 Hz.\n[Sumi H. et al. Journal of power sources 226, 354-358 (2013). Sumi H. et al. Scientific Reports 11,\n10622 (2021). Lim D-K. et al. Ceramics International 42, 3776-3785 (2016)].\n20%, and 30%), and (b) Arrhenius plots for BCC and s-BCC-Y (15%), and (c) EIS Nyquist plots\nfor s-BCC-Y and BCC, including the DFT analysis for them at 650 (d), 600 (e), 550°C (f).\nAs shown in the DRT of Fig. S18, the differences between BCC and s-BCC-Y are mainly\nobserved at mid- and high-frequencies, suggesting that s-BCC-Y may have a better oxygen ion\nconductivity and redox catalytic activity.\nIn the cyclic testing DRT analysis shown in Fig. S20, Thermal cycling could cause delamination\nor other microstructural alterations, disrupting the ionic and electronic conductivity networks, which\nreduces the effective reaction area. This could lead to localized high current density and,\nconsequently, an increase in low-frequency impedance.\nFig. S20 Representative EIS Nyquist plots for the BCC (a) and s-BCC-Y (b) symmetric cell with\nBZCYYb electrolyte after 35th cycles, and the DRT for BCC (c) and s-BCC-Y (d).\nthe data presented in Fig. S16 (now change to Fig. S22 and Fig. S23). The differences are also\nreflected in the low-frequency region. The CO poisoning significantly hinders gas diffusion by\n2\nforming carbonate layers.\nFig. S22 ASR for BCC and s-BCC-Y symmetric cell at different time during CO poison and air\n2\nrecovery.\nFig. S23 The DRT for BCC (a, c) and s-BCC-Y (b, d) symmetric cell at different time during CO\n2\npoison (CO on) and air recovery (CO off).\n2 2\nTherefore, we revised the manuscript by referring to DRT analysis in the main text and providing\ndetailed discussion in the SI. The revisions in main text on Line 347-363 include:\n“…For example, the Rp of s-BCC-Y was only 0.23 Ω m2, 28.1% lower than that of BCC (0.32\nΩ cm2). The DRT analysis (SI-14, Fig. S18) shown the differences between BCC and s-BCC-Y are\nmainly observed at mid- and high-frequencies, suggesting that s-BCC-Y has a better oxygen ion\nconductivity and redox catalytic activity. The stability of the Rp of symmetric cells at OCV was\nalso tested. After 300 hours, the Rp of BCC increased from 0.325 Ω cm² to 0.71 Ω cm², showing\na degradation of 1.28 mΩ cm²/h, while the decay of s-BCC-Y’s is only 0.58 mΩ cm²/h (SI-14,\nFig. S19)..…After 35 cycles (~120 h) of rapid temperature cycling between 300-600°C, the Rp\nvalue of the s-BCC-Y electrode only slightly increased from 0.23 to 0.33 Ω cm2, while BCC\nincreased from 0.34 to 0.59 Ω cm2 by almost 2 times (Nyquist plots and DRT shown in SI-14, Fig.\nindicating that slow processes, such as gas diffusion within the BCC cathode, begin to change, likely\ndue to interfacial delamination or electrode fracture. In contrast, s-BCC-Y shows smaller changes\nat low frequencies. Besides, s-BCC-Y also exhibited better CO resistance (see SI-15, Fig. S21).\nBCC, which indicates that gas diffusion is suppressed, likely due to cathode surface carbonization\nand the coverage of carbonates impeding oxygen diffusion.”\n8\n[Comment 4]: Degradation is also observed in Fig.5 e and f and assignment of\ndegradation of performance is also required.\n[Response 4]: Thanks for your kind reminder. We have marked the performance degradation and\ndegradation rates to better serve the readers. In Fig. 5e, the cell output shows a voltage drop from\n0.797 V to 0.758 V after 400 hours, corresponding to a 0.103 mV/h degradation. In Fig. 5f, the\nvoltage decreases from 0.753 V to 0.71 V, representing a 0.153 mV/h degradation. It should be\nnoted that this degradation can be caused by many factors, including the degradation of seal\n[Chou, Y. et al. Journal of Power Sources 191, 384-389 (2009)], anode microstructure\n[Holzer, L. et al. Journal of Power Sources 196, 1279-1294 (2011)], current collector etc. [Guo, Y. et al.\nJournal of Power Sources 196, 5511-5519 (2011)].\nIn addition, we would like to emphasize that the performance of the cell largely depends on the\nfabrication experience, which can vary between laboratories. The stability data we present for s-\nBCC-Y may still be observable compared to the highly stable cathodes reported in the literature\nwith a degradation rate of 0.104 mV/h under YSZ and 0.153 mV/h under BZCYYb electrolytes of\ns-BCC-Y, which, however, is still a good performance based on our specific testing environment\nand fabrication process.\nFor reviewer’s reference, we provided stability test data of other classic materials and state-of-art\nmaterials under development in our lab, as shown in following table, which suggests that the testing\nconditions may not have fully showcased the potential stability of this cathode. We are confident\nthat with further optimization of the fabrication and testing processes, the cathode with our hydration\nstrategy (s-BCC-Y) has significant potential for further improvement of the durability. Considering\nthat these are experiments that have not been published yet, we did not include them in the main\ntext or SI.\nTable. Some of the SOFC stability data currently being tested in our laboratory\nCathode Electrolyte Anode Condition Duration Decay Rate\n550 °C\nBCFZY BZCYYb Ni- BZCYYb 350h 0.4469 mV/h\nH -air\n2\n550 °C\nBCFZY BZCYYb Ni-BZCYYb# 140h 0.2823 mV/h\nH -air\n2\n550 °C\nBCFZY BZCYYb Ni-BZCYYb# 260h 0.1883 mV/h\nH -air\n2\n700 °C\nBCFZY BZCYYb Ni- BZCYYb 150h 0.345 mV/h\nH -3%H O-air\n2 2\n700 °C\nBSCF SDC Ni-SDC 120h 0.778 mV/h\nH -air\n2\n750 °C\nPBCC# YSZ Ni-YSZ 200h 0.281 mV/h\nH -air\n2\n550 °C\nBCFN-Ni# BZCYYb Ni- BZCYYb 200h 0.35 mV/h\nH -air\n2\n550 °C\nBCFYN# BZCYYb Ni- BZCYYb 200h 0.175 mV/h\nH -air\n2\n550 °C\nBFCCC# BZCYYb Ni- BZCYYb 300h 0.32 mV/h\nH -5%H O-air\n2 2\n550 °C\nBCFN# BZCYYb Ni- BZCYYb 230h 0.14 mV/h\nH -3%H O-air\n2 2\n600 °C\nBCFZ# BZCYYb Ni- BZCYYb 200h 0.157 mV/h\nH -air\n2\n# indicates materials being developed that has not been published.\nelectrolyte(f)\nBased on these results, we added revisions as follows:\nLines 382-387: “Both cells were able to sustain a consistent power output over an extended duration.\nIn Fig. 5e, the cell output shows a voltage drop from 0.797 V to 0.758 V after 380 hours,\ncorresponding to a 0.103mV/h degradation. In Fig. 5f, the voltage decreases from 0.753 V to 0.71\nV, representing a 0.153 mV/h degradation. This strongly suggests that the s-BCC-Y composite\ncathode, improved by stiffness enhancement, exhibits remarkable performance in single fuel cells.”\n10\n[Comment 5]: Reaction of BCC is also expected and doped Y may be prevent reaction\nbetween cathode and BZCYYb electrolyte. So, suppression of chemical reaction is\nanother reason for increased performance of the cell.\n[Response 5]: Thank you very much for this comment. we indeed overlooked this, but it is\nundoubtedly a very important point. To verify whether our cathodes (BCC and s-BCC-Y) react with\nthe electrolyte (BZCYYb), we conducted additional XRD tests on their mixtures after calcination.\nWe mixed the cathode and electrolyte powders and calcinate the mixture at 800 °C for 5 hours.\nThe XRD results showed no impurities or new peaks observed, indicating that neither BCC nor s-\nBCC-Y reacts with the electrolyte. We have included this description in the main text and added the\nXRD figures to the supplementary information.\nFig. S24 XRD patterns of (a) s-BCC-Y, s-BCC-Y+ BZCYYb (1:1, weight ratio) and (b) BCC,\nBZCYYb, BCC+ BZCYYb (1:1, weight ratio) after the calcination at 800 °C for 2 h in air.\nBased on these results, we added revisions as follows on Lines 370-373:\n“In addition, it should be mentioned that no reactions were observed for both air electrodes and\nBZCYYb (SI-16, Fig. S24), Further confirming that the cracking and delamination observed at the\ninterface were mainly caused by poor stiffness and thermo-mechanical properties in this work,\nrather than chemical side reactions.”\n11\n[Comment 6]: In summary, the reason why only Y can increase the mechanical\nproperty of BCC is required in spite of small amount of doped and still large difference\nin TEC.\n[Response 6]: Thank you for your question. We would like to clarify that we are not suggesting\nthat only yttrium (Y) doping improves the mechanical properties of (BCC). The central idea of our\nwork is that material engineering techniques, such as acceptor-doping (with Y being one example,\nbut not the only option), can enhance the material’s hydration ability. This improved hydration,\nalong with the associated “chemical expansion”, significantly impacts the material’s behavior,\nparticularly in terms of thermal expansion and performance under temperature fluctuations, which\nare typical in high-temperature fuel cells.\nAlthough the difference in TEC between s-BCC-Y and other materials like YSZ remains\nsubstantial, even small amounts of Y doping significantly enhance the hydration ability. This change\nleads to chemical expansion, which, while not fully compensating for the TEC mismatch, greatly\nimproves the mechanical properties of the cathode material, especially its resistance to cracking.\nThis improvement in anti-cracking behavior is the main factor contributing to the observed increase\nin mechanical robustness.\nBase on reviewer’s comments and suggestions, we have further emphasized this point in the\ndiscussion section of the manuscript. We hope this explanation clarifies the rationale behind the Y-\ndoping strategy and its effects on the mechanical properties of BCC. The enhanced hydration leads\nto significant changes in the material’s stiffness and its overall robustness.\nBased on these results, we added revisions as follows on Lines 403-411:\n“This study indicates that material engineering techniques, such as acceptor-doping (with Y being\none example, but not the only option), can enhance the material’s hydration ability. This improved\nhydration, along with the associated “chemical expansion”, significantly impacts the material’s\nthermal expansion and stiffness under temperature fluctuations, which are typical in high-\ntemperature fuel cells. Although the difference in TEC between s-BCC-Y and other materials like\nBZCYYb or YSZ remains substantial, even small enhancement of hydration ability can greatly\nimprove the mechanical properties of the cathode material, especially its resistance to cracking. This\nimprovement in anti-cracking behavior is the main factor contributing to the observed increase in\nmechanical robustness, taking slight degradation and enhanced long-term stability during extended\nthermal cycling tests.”\nThis paper investigates the mechanical and electrochemical performance of protonic\nceramic fuel cells (PCFCs), focusing on enhancing the durability of the cathode through\nchemical hydration. The authors introduce Y-doping (yttrium doping) into a perovskite-\nbased material, BaCo Ce Y O (s-BCC-Y), to mitigate thermal expansion\n0.7 0.15 0.15 3-δ\nduring high-temperature cycling, thereby enhancing mechanical strength and reducing\nthe risk of cracks and delamination in the cathode. Moreover, the Y-doped cathode\ndemonstrates superior electrochemical performance and long-term stability compared\nto conventional cathodes, especially under thermal cycling conditions. These results are\ninteresting; however, there are several issues that need to be addressed before the paper\ncan be published, requiring major revisions. Please refer to the comments provided\nbelow.\n[Comment 1]: In Figures 1c-f, the author notes that Y-doping causes changes in the\ncrystal structure. However, the precise atomic-level modifications, such as how Y ions\nare substituted within the lattice and how they affect bond lengths and angles, are not\nthoroughly investigated. Thus, authors should further examine this crystal structure\nchanges at the atomic scale.\n[Response 1]: Thank you very much for pointing out this issue. To address your question\nregarding the lattice change after Y doping, we conducted both TEM and XRD analyses, which\nrevealed the lattice structure, as shown in Fig. 1c-f. In the original paper, we have brief described:\n“As expected, the s-BCC-Y material was self-assembled into two crystalline phases, a major cubic\nperovskite phase and a minor rhombohedral phase based on the Rietveld refinement”. To make this\nclearer, we added/change more comments to enhance the conclusion: “To better present how Y\naffected these two phases, we conducted DFT calculations and performed iterative self-consistent\ncycles for structural optimization (mainly focusing on bond strength and angles). All structures\nwere fully relaxed until the maximum force on each atom was less than 0.01 eV/Å, and the energy\nconvergence criterion was set to 10⁻⁸ eV (as shown in SI-4)”.\nAlthough we did not observe a very distinct Y column in HR-TEM due to the limitation of TEM\nresolution, the clear lattice structure observed in XRD and TEM are sufficient to conclude the\nsuccessful homogeneous doping of Y. We then employed DFT calculations to study the lattice\nstructure (in terms of bond lengths and angles), as seen in Fig. S3, S4.\nIn the DFT calculations, Y doping of two different phases (BaCoO and BaCeO ) are simulated\n3 3\naccording to the XRD refinements of s-BCC-Y. all structures were firstly fully-relaxed until the\nmaximum force on each atom was less than 0.01 eV/Å, and the energy convergence criterion was\nset to 10-8 eV. Ultimately, we obtained predicted changes in the crystal structure before and after Y-\ndoping, including variations in bond lengths, bond angles, and other parameters. These results have\nbeen included in the SI.\nFor the Co-rich phase (BaCoO , Pm-3m), the introduction of Y alters the original octahedral\n3\nconfiguration, resulting in changes in the Co-O bond lengths and angles, as shown SI-4, Fig. S3.\nSpecifically, the Co-O bond length along the Co-Y direction decreases from 1.988 Å to 1.825 Å,\nindicating an enhancement of the Co-O bond strength. This strengthening contributes to the\nimproved structural stiffness according to [Wang, X. et al. Physical Chemistry Chemical Physics 20,\n13\n17781-17789 (2018)]. In terms of bond angles, the deformation of the octahedron results from the\noverall influence of the surrounding Y atoms, whereas the Pm-3m structure does not exhibit\nsignificant changes in its overall structure. Additionally, the newly introduced Y-centered\noctahedron has Y-O bond lengths of 2.237 Å.\nFig. S3 Changes in the crystal structure, bond lengths, and bond angles of the Co-rich phase (Pm-\n3m) before and after the introduction of Y.\nFor the Ce-rich phase (BaCeO , R-3c), we applied the same processing and analysis methods as\n3\nthose used for the Co-rich phase (Pm-3m). The results are shown in Fig. S4. Unlike the Co-rich\nphase (Pm-3m), the Ce-O bond length along the Ce-Y direction increases slightly from 2.263 Å to\n2.279 Å due to the larger atomic radius of Ce compared to Co and its similarity to Y. The changes\nin bond angles of B-O-B are also minor, shifting from 158.05° to 158.22°. Although the increase is\nmodest, the larger B-O-B bond angle tends to exhibit better electrochemical performance. [Nan, H.-\ns. et al. Materials Science in Semiconductor Processing 94, 35-50 (2019), Lei, N. et al. Chemical\nEngineering Journal 421, 127790 (2021)]. Based on the results above. we have added the\ndescription after the XRD and TEM sections (lines 119-125, 127-131 and all SI-4) in the main text\nto provide readers with a clearer depiction of the lattice changes at atomic level and demonstrate\nhow Y atoms influence the unit cell, as well as the bond lengths and bond angles.\n3c) before and after the introduction of Y.\nBased on these results, we added revisions as follows on\nLines 119-125:\n“As expected, the s-BCC-Y material was self-assembled into two crystalline phases, a major\ncubic perovskite phase (89.7%, Pm-3m, BaCo Ce Y O ) and a minor rhombohedral phase\n0.74 0.094 0.166 3\n(10.3%, R-3c, BaCo Ce Y O ) based on the Rietveld refinement (details shown in SI-2). In\n0.11 0.86 0.03 3\naddition, lattice spacings were measured from this selected region, BaCoO (102) and BaCeO (110)\n3 3\nin BCC, Pm-3m (100) and R-3c (110) in s-BCC-Y were clearly observed, in good accordance with\nthe XRD refinement results, as shown in Fig. 1e, f.”\nLines 127-131:\n“To better present how Y doping affected these two phases, we conducted DFT calculations and\nperformed iterative self-consistent cycles for structural optimization (mainly focusing on bond\nstrength and angles). All structures were fully relaxed until the maximum force on each atom was\nless than 0.01 eV/Å, and the energy convergence criterion was set to 10⁻⁸ eV (as shown in SI-4).”\nSI-4:\nIn DFT calculations and iterative self-consistent cycles for structural optimization1, 2, all\nstructures were fully relaxed until the maximum force on each atom was less than 0.01 eV/Å, and\nthe energy convergence criterion was set to 10-8 eV. After completing the calculations in VASP, the\nresults were visualized using VESTA 3.\nFor the Co-rich phase (Pm-3m), Y-doping primarily affects the Co-centered regular octahedral\nstructure. The introduction of Y alters the original octahedral configuration, resulting in changes in\nthe Co-O bond lengths and angles, as shown in the Fig. S3. Specifically, the Co-O bond length\nalong the Co-Y direction decreases from 1.988 Å to 1.825 Å, indicating an enhancement of the Co-\nO bond strength. In terms of bond angles, the deformation of the octahedron results from the overall\ninfluence of the surrounding Y atoms, whereas the Pm-3m structure does not exhibit significant\nchanges in its overall structure. Additionally, the newly introduced Y-centered octahedron has Y-O\nbond lengths of 2.237 Å.\n3m) before and after the introduction of Y.\nFor the Ce-rich phase (R-3c), we applied the same processing and analysis methods as those\nused for the Co-rich phase (Pm-3m). The results are shown in Fig. S4. Unlike the Co-rich phase\n(Pm-3m), the Ce-O bond length along the Ce-Y direction increases slightly from 2.263 Å to 2.279\nÅ due to the larger atomic radius of Ce compared to Co and its similarity to Y. The changes in bond\nangles of B-O-B are also minor, shifting from 158.05° to 158.22°. Although the increase is modest,\nthe larger B-O-B bond angle tends to exhibit better electrochemical performance.\nFig. S4 Changes in the crystal structure, bond lengths, and bond angles of the Ce-rich phase (R-\n3c) before and after the introduction of Y.\n16\n[Comment 2]: Figure 1g, a detailed explanation of how Y-doping increases oxygen\nvacancies is recommended. Additionally, since the authors only present the weight\nchange curve for s-BCC-Y, it may give the impression that s-BCC-Y and BCC behave\nidentically under dry-air conditions. Revising this figure would help prevent potential\nmisunderstandings.\n[Response 2]: Thank you very much for pointing out this issue. In Fig. 1g of the original paper,\nthe presentation was too complex, which may cause misunderstandings to the readers.\nRegarding how oxygen vacancies are influenced by Y doping, in brief, Y3+ doping is a commonly\nused acceptor doping strategy for B site cations, that are required to be 4+. Therefore, due to the\ncharge balance, the acceptor doping at B site (Y3+ doping) will lead to the loss of lattice oxygen,\nforming vacancy. This is a very accepted theory [Wexler. et al., Journal of the American Chemical\nSociety 143, 13212-13227 (2021)]. To make this clear to readers, we enhance this at Line 107-109.\nThe reviewer is also suggested to refer to our response to the next comment for more details\n(Comment 3).\nRegarding the TG curves, in this response, we attempted to amend the TG experimental as\nrequired by the reviewer. The TG curve of BCC in dry air is added as seen in the update Fig. 1g,\nHowever, regarding the water-containing TG measurement, the TG equipment we previously used\n(TA TGA 550), actually, like most TG devices, does not sustain the inflow of high-water vapor for\nlong time due to potential damage to sensors. After we continuously introducing water vapor for\ntesting, this equipment malfunctioned and has been under repair since then. And more pitifully, all\nother third parties we sought also refused to conduct water-contained TG test. We sincerely apology\nfor we may not be able to provide the experimental data of water-contained TG for BCC.\nNonetheless, we are able to make up using water contained Temperature Programmed Desorption\n(TPD) test to demonstrate the material’s water absorption and hydration behavior along with the dry\nTG results. The TPD technique, as a classic and commonly used method, can detect the desorption\nof substances from the material at high temperatures, particularly suitable for detecting desorbed\nwater and oxygen.\nFig. 1 g, h TG and DTG analysis of hydrated BCC and s-BCC-Y from RT to 1000℃ (g). H O-\nFirst, we performed dry TG tests of BCC and s-BCC-Y (from RT to 1000℃) that are fully\nhydrated in wet air before test. In the TG analysis in dry air, the mass loss of hydrated BCC and s-\nBCC-Y occurs in three stages: 100–200 °C, 400–500 °C, and 700–800 °C. Water loss within the\n100–200 °C range corresponds to the release of adsorbed and crystallized water from the porous\nsurface and pore structure of the powders [Gao, Y. et al. Advanced Functional Materials, 2416625\n(2024). Aguadero, A. et al. Chemistry of Materials 19, 6437-6444 (2007). Ding, X. et al. Applied\nCatalysis B: Environmental 243, 546-555 (2019)]. Key differences in mass loss between the two\nmaterials are observed at approximately 475 °C and 770 °C, indicating distinct behaviors in\nhydration and lattice stability.\nTo comprehend this, we also conducted O -TPD (Fig. S5) and H O-TPD (Fig. 1h) tests. The O -\n2 2\nTPD revealed that the mass loss in both BCC and s-BCC-Y is indeed due to the loss of lattice oxygen\nat as many other papers reported [Song, Y. et al. Joule 3, 2842-2853 (2019)]. However, in H O-TPD,\n2\nthe case of s-BCC-Y, at temperatures of 475°C and 774°C, in addition to detecting oxygen signals,\nthe H O signals were also observed. This indicates that the oxygen vacancies in s-BCC-Y contain a\n2\nsignificant amount of water molecules introduced by hydration. This also suggests that s-BCC-Y\nhas a stronger hydration ability compared to the BCC. [ Liu, Z. et al. Nature Communications 15,\n472 (2024)].\nBased on these results, we added revisions as follows:\nAt lines 107-109:\n“Moreover, possibly increasing oxygen vacancy, which means that oxygen ions have more\npathways for transport, which facilitates the migration of oxygen ions and the reduction reaction,\nthereby enhancing the catalytic activity of the SOFC cathode 21-23.”\nAt lines 142-154:\n“To further validate the enhanced hydration capacity introduced by Y doping in BCC, we\ncompared the weight changes of BCC and s-BCC-Y from 30 to 1000 °C in dry air, as shown in the\nTG and DTG analyses in Fig. 1g. Beyond the typical water loss in the 100-200 °C range, commonly\nobserved in porous materials, both BCC and s-BCC-Y exhibit notable mass losses from 400 to\n1000 °C, particularly around 450 °C and 750 °C. This behavior aligns with that of many\ncathode materials 31,32. Extensive studies 33-35 have attributed this phenomenon primarily to the loss\nof lattice oxygen in perovskite cathodes, which is an indication of hydration. Therefore, the observed\n18\nenhanced weight loss at 450°C could be caused by the enhanced hydration in Y doped s-BCC-Y.\nTemperature-programmed desorption (TPD) tests were also conducted in water and oxygen\nenvironments to confirm the TG results. The results shown in Fig. 1h (H O-TPD) and SI-5 (O -\n2 2\nTPD), indicate that the mass loss in BCC in TG is mainly due to loss of lattice oxygen, while s-\nBCC-Y’ mass loss not only contributes to lattice oxygen but a significant amount of water. This\nobservation again highlights the superior hydration capability of s-BCC-Y compared to BCC.”\n19\n[Comment 3]: The authors demonstrate that Y-doping influences oxygen vacancy\nformation. However, the work lacks a detailed analysis of the atomic interactions and\nthe specific role these vacancies play within the lattice (Figures 2b and 2c). Furthermore,\nsince the Y dopant can alter the oxidation states of both Ce and Co, excluding the Ce\nion from the DFT calculations is not justified. Authors should conduct a more in-depth\nanalysis and provide a detailed explanation.\n[Response 3]: Thank you very much for pointing out this issue. Our explanation of oxygen\nvacancies (OVs) in the previous manuscript was too simplistic, causing some confusion for readers.\nIn response, we have added an experimental measurement of the OVs and further provided a detailed\nexplanation on atomic interactions and the specific role of OVs in the revised manuscript.\nIn general, the doping of trivalent Y will lead to the charge imbalance in ABO system, of which\n3\nthe B site cation is advised to be 4+. Many studies have proved that the Y doping can lead to the\nformation of additional oxygen vacancy in ABO perovskite [Wang, Y. et al. Applied Catalysis B:\ninteractions with the other B site by change of valence. In the previous manuscript, we have already\nstudied the valence changes of Co elements which is a good indication of increase of oxygen\nvacancies. It has been confirmed by numerous studies on oxygen vacancies and B site valence in\nperovskite, especially in SOFC cathode. For example, Ji Yang deduced the increase of oxygen\nvacancies in LSCO from the Co²⁺/Co³⁺ ratio (from 0.87 to 1.47) in XPS [Yang, J. et al. ACS Catalysis\n9, 9751-9763 (2019)]. Manjeet Kumar concluded that the increase in oxygen vacancies results from\nthe conversion of Ce⁴⁺ to Ce³⁺ [Kumar, M. et al. Electrochimica Acta 284, 709-720 (2018)]. There\nare also similar studies on oxygen vacancies in Mn and Fe valence decreasing [Yang, W. et al.\nApplied Catalysis B: Environmental 260, 118150 (2020), Sedykh, V. et al. Ceramics International\n49, 25640-25648 (2023)].\nRegarding the reviewer’s comment on the specific role these vacancies play within the lattice,\nnormally, during heating, metal cations at the B site undergo valence changes, and oxygen ions are\nreduced and escape from the lattice, forming oxygen vacancies. As the transport channels for oxygen\nions in perovskites, oxygen vacancies, enhance the diffusion of oxygen ions, thus improving ionic\nconductivity. They can also serve as active sites for reactions, thereby increasing catalytic activity.\nThe metal cations that coordinate with the missing oxygen ions will shift to maintain local charge\nneutrality. In this process, the interaction between metal ions and the remaining lattice oxygen\nchanges to reduce lattice strain, and the atomic bonding (chemical bonds) between the ions also\nchanges. Based on this knowledge, we have also made brief comments in the main text (lines 110-\n112) that: Moreover, possibly increasing oxygen vacancy, which means that oxygen ions have more\npathways for transport, which facilitates the migration of oxygen ions and the reduction reaction,\nthereby enhancing the catalytic activity of the SOFC cathode. We hope this will help readers quickly\nunderstand the role of the oxygen vacancies mentioned in the manuscript.\nLastly, regarding the reviewer’s concern about excluding Ce elements, we agree that this was\nnot justified. Therefore, we have further added DFT-PDOS calculations and XPS valence state\nanalysis for Ce in both BCC and s-BCC-Y as shown in the following Fig. S6. As we considered in\nthe initial draft, the Ce-rich phase in our materials is relatively small (~10%), so its influence is\nlimited, and this is why we have placed these results in the supporting information.\nAccording to DFT calculations, Ce’s PDOS crosses the Fermi level, which suggests an\n20\nimprovement in the material’s conductivity, consistent with the high conductivity of Ce-rich phases\nreported previously. [Song, Y. et al. Joule 3, 2842-2853 (2019)].\nCe’s XPS analysis [Wang H, et al. Nature Communications 15, 3765 (2024). Pereira-Hernández\nXI, et al. Nature Communications 10, 1358 (2019)] also shows a reduction in Ce’s chemical valence,\nCe3+ increased from 37.11% to 44.91%, which is consistent with the reduction in Co’s chemical\nvalence mentioned in the main text, promoting the increase of oxygen vacancies and enhancing the\nmaterial’s oxygen ion conductivity.\nFig. S6 The PDOS (a) and XPS (b) of Ce-rich phase in BCC and s-BCC-Y.\nBased on these results, we added details were shown in SI-7 and made revisions in the main text\nas follows on Lines 107-109 and 200-205:\nAt SI-7:\n“Through the DFT analysis of Co, we observe the appearance of resonance peaks, indicating the\nstrengthening of Co-O chemical bonds. However, the effect is not as pronounced as that of Ce. This\nis attributed to the similar ionic radii of Ce and Y, resulting in a relatively minor impact from Y\ndoping. This observation is also reflected in the lattice calculations performed using DFT. More\nnotably, Ce’s PDOS crosses the Fermi level, which suggests an improvement in the material’s\nconductivity, consistent with the high conductivity of Ce-rich phases reported previously 1. Ce’s\nXPS analysis also shows a reduction in Ce’s chemical valence, which is consistent with the\nreduction in Co’s chemical valence mentioned in the main text, promoting the increase of oxygen\nvacancies and enhancing the material's oxygen ion conductivity. The results shown in Fig. S6.”\nAt lines 107-109\n“Moreover, the increased oxygen vacancy, which means that oxygen ions have more pathways\nfor transport, could facilitate the migration of oxygen ions, thereby enhancing the catalytic activity\nof the SOFC cathode 21-23.”\nAt lines 191-196\n“The Ce-rich phase was further characterized using PDOS from DFT calculations and XPS, as\nshown in Fig. S6. The PDOS results indicate that Ce’s electronic states cross the Fermi level,\nsuggesting enhanced conductivity, consistent with previous reports on the high conductivity of\nYttrium-doped barium zirconate48. XPS analysis further reveals a reduction in Ce’s chemical\nvalence, which correlates with the observed decrease in Co’s chemical valence.”\n21\n[Comment 4]: When comparing the Y-doped and undoped BSCFs in Figure 4b, the\ndifference in the increase in fracture strength from the untreated to the hydrated state is\nnegligible. Regarding this, authors need to explain the detail.\n[Response 4]: Regarding the reviewer’s comment that the difference between BSCF and BSCFY\nbefore and after hydration is not significant. However, actually, the difference is not negligible (or\nbetter to say, the increase is “mild”) in the views of authors. For BSCF, the fracture strength\nincreases from 38.5 to 57.3 MPa by 48.8%. For BSCFY, the fracture strength increases from 44.0\nto 71.3 MPa by 62%. If we compare the BSCF and BSCFY, the improvement is 14.3% and 24.4%\nfor un-hydrated and hydrated samples, respectively.\nThe authors acknowledge that this increase is not as much as BCC and s-BCCY. This is possible\nbecause other factors that also govern the low fracture strength of BSCF. To explain the “mild”\nimprovement of strength of BSCF after Y doping, we would like to conjecture that:\nFirst, the high fraction of Co and Fe in BSCF already results in high fracture of oxygen vacancy\nthat the further doping of Y does not show significant improvement of hydration capacity;\nSecond, the high reducibility of Co3/4+ and Fe3/4+ ions (absent in valence-fixed Y3+) at high\ntemperature is the origin of its large TEC (>20 ×10-6 K-1), and thus large thermal stress between\ngrains. That implies the high fraction of Co/Fe in BSCF than BCC could dominate the fracture\nstrength, other than the hydration capacity.\nIt should also be noted to the reviewer and readers that explanation of why BSCF shows low\nfracture strength than BCCY system is very challenging as too many factors need be considered,\nincluding bulk/surface/boundary physical-chemical properties, such as porosity, phase/lattice\nstructure, heat treatment history, grain size, reactivity/adsorption activity to CO /H O, bond energy,\n2 2\ngrain boundary structure (coherency), etc. However. BSCF and BCCY are “too different” in the\nabove-mentioned terms. More investigation into this is therefore needed, but beyond the main scope\nof this work. Inspired by the review, the authors deem this as a very interesting topic to study (it is\nalso in the research scope of our group) and need to be systematically revisited the correlation\nbetween strength and those factors. As the reviewer may beware of, currently there is not too many\ngroups are looking into this area.\nTo make this clear to readers, we redraw Fig. 4d (see below) to better present the difference of\nstrength in BSCF and BSCFY by adding the actual values on the bar chart and provide the comments\nin the main text on Line 292-308 as follows:\n“To further elucidate the universality of the strategy of hydration induced for mechanical\nenhancement, we also compared the fracture strength of the classic single cubic phase cathode\nBa Sr Co Fe O (BSCF) and B-site 10% Y doped Ba Sr (Co Fe ) Y O (noted as\n0.5 0.5 0.8 0.2 3 0.5 0.5 0.8 0.2 0.9 0.1 3\nBSCFY with XRD profile shown in SI-12), before and after the full hydration (800°C, 10 vol.%-\nH O in air, 5 h), which were shown in Fig. 4d. As it presented, Y-doping effectively enhanced the\n2\nfracture strength of BSCF, both before and after full hydration, which means hydration is not limited\nto the composite BCC system but also applies to single-phase cubic perovskites. However, this\nimprovement is less significant compared to that observed in BCC. There might be two reasons for\nthat. First, the high fraction of Co and Fe in BSCF already results in high fracture of oxygen vacancy\nthat the further doping of Y does not show significant improvement of hydration capacity, Second,\nthe high reducibility of Co3/4+ and Fe3/4+ ions (absent in valence-fixed Y3+) at high temperature is\nthe origin of its large TEC (>20 ×10-6 K-1), and thus large thermal stress between grains. That implies\n22\nthe high fraction of Co/Fe in BSCF than BCC could dominate the fracture strength, other than the\nhydration capacity. Besides, we also doped the Yb element into BCC (nominally\nBaCo Ce Yb O , noted as BCCYb with XRD profile shown in SI-12), which is also a trivalent\n0.7 0.15 0.15 3\ndopant, creating oxygen vacancies for hydration, similar to the Y element. The BCCYb also showed\nimproved strength than BCC (see Fig. 4d), further confirming that the enhanced hydration due to\ntrivalent doping can enhance fracture strength.”\nFig 4. d Comparison of fracture strength before and after hydration for BCC, s-BCC-Y, BSCF,\nBSCFY, and BCCYb cathode materials.\n23\n[Comment 5]: In Figure 4e, the comparison of the two samples using FIB-SEM does\nnot specify whether they underwent any particular treatment prior to analysis. Cracks\nmay have formed during the sampling process, so it is essential to confirm if the\nsamples were exposed to treatments such as rapid heating-cooling cycles. Including\ndata from treated samples would more effectively demonstrate their superior thermal\nshrinkage resistance and provide a more thorough comparison.\n[Response 5]: Thank you very much for pointing out this issue. In the previous manuscript, we\ndid not clearly explain the sample treatment before FIB-SEM testing. We would like to provide\nadditional details in the manuscript:\nFor the sample in Fig. 4e, after sintering, the sample was naturally cooled to room temperature\nand then carefully transferred to the electron microscope for observation. No rapid heating or\ncooling cycles were applied; the process involved only heating followed by natural cooling, and the\nmilling process for sampling by ions beam in FIB-SEM was believed to has no significant effect to\nthe original cracks as the space resolution of beam milling can be down to 1 nm scale, that is much\nsmaller than the crack scale(~micro) and good Pt deposition was also prepared on the sample surface\nto reduce the undesired damage by ion milling.\nTo provide a thorough comparison as suggested by reviewer, we conducted supplementary\nexperiments during this revision period. We quenched the dense BCC and s-BCC-Y samples from\nhigh temperature (1150°C) to RT, by direct taking the samples from furnace to the water (RT). which\nmay cause rapid shrinkage and possible thermal-induced cracks. As shown in the following surface\nof quenched sample (Fig. S17), the cracks formed in the BCC sample were much more severe than\nthose under natural cooling and the quenched Y-doped samples as well. This demonstrates the\nsuperior thermal shrinkage resistance and stiffness of s-BCC-Y.\nFig. S17 SEM observations of BCC and s-BCC-Y after quenched from 1150°C to RT in water.\nBased on these results, we added details of quenching experiments in SI-13 and made revisions\nin the main text as follows on Lines 316-321:\n“Besides, a rapid quenching method was conducted to examine their crack resistance. Dense BCC\nand s-BCC-Y samples were quenched from high temperature (1150 °C) to RT, by direct taking the\nsamples from furnace to the water (RT). which subjects the samples to rapid shrinkage and induces\npossible thermal cracks. As shown in the following surface (SI-13, Fig. S17), the cracks formed in\nthe BCC sample were much more severe than the Y-doped samples, demonstrating the superior\nthermal shrinkage resistance of s-BCC-Y.”\n24\n[Comment 6]: To strengthen the authors’ claims, SEM images of BCC and s-BCC-Y\nbefore operation should be included alongside the images presented in Figure 5b.\n[Response 6]: Thank you very much for your reminder. We had previously placed the pre-\noperation SEM images in the supplementary section, which may not serve the readers well. We have\nnow updated Fig. 5b by placing the pre- and post-test SEM images together, making it easier for\nreaders to compare.\nFig 5. b Cross-sectional SEM and schematic diagram of a symmetric cell before and after cycling:\ncrack, loss of the cathode, and delamination from electrolyte.\nAccordingly, we also added comments on Lines 364-370 for the comparison:\n“We examined the SEM morphology of BCC and s-BCC-Y before and post-cycling (Fig. 5b).\nAs anticipated, not only delamination but also bulk crack (leads to the loss of partial cathode in the\nbulk) were observed in the BCC sample post-cycling. As illustrated in the schematic diagram (Fig.\n5b), we attribute the extensive cathode loss observed after long-term cycling to fractures caused by\nthermal stress at the interface of a low-stiffness cathode. By contrast, the improved stiffness and\nfracture strength enabled s-BCC-Y to retain its intact porous framework without any delamination\nand bulk cracks, which is key to superior durability.”\n”\nThe manuscript details a study on improving the mechanical durability of high-\ntemperature fuel cell cathode materials to prevent degradation. It proposes doping\nBaCo Ce O perovskite with yttrium (Y) to enhance thermal expansion tolerance and\n0.7 0.3 3\nmechanical properties. This approach has been validated through structural, mechanical,\nand electrochemical evaluations, showing increased cathodic stiffness and fracture\nstrength and resulting in better fuel cell performance and stability. The innovation is\nsignificant for advancing SOFCs and PCFCs.\n[Comment 1]: The manuscript is dense with technical details and data, which is\ncommendable, but it may benefit from a more streamlined presentation to enhance\nclarity and readability. Simplifying the more complex sections, or providing clearer\nexplanations, might make the document more accessible to a broader audience.\n[Response 1]: Thank you for your suggestion. To enhance the clarity of our study, we have\nincluded many details and data in the manuscript. Your kind comment made us realize that the\nstructure and flow of our content must be improved. The extensive experimental data presented may\nhave made it difficult for readers to quickly extract key information.\nUpon re-examining the manuscript’s structure, we found that some section titles lacked clarity.\nIn response, we have revised these titles to improve readability and the overall presentation of our\nwork. As shown below and also in table.\nConcept of chemical induced expansion for stiffness enhancement\nEvidence in dilatometry and mitigated interface stress by simulation\nImproved bulk stiffness and fracture strength\nElectrochemical evaluation of s-BCC-Y upon thermal cycling\n(Revised)\nHydration-induced chemical expansion concept (at line 93)\nReduction of TEC and interface optimization (at line 162)\nHydration-enhanced bulk stiffness at a micro-scale (at line 223)\nHydration-improved fracture stiffness at a macro-scale (at line 273)\nElectrochemical evaluation of hydrated s-BCC-Y during thermal cycling (at line 340)\nIn addition, we have made several modifications to the detailed descriptions throughout the\nmanuscript. We hope that the revised section titles and improved details could provide clearer\nexplanations and a more streamlined presentation for the audience. For example,\nAt lines 163-165: “The hydration-induced expansion is expected to have a mitigating impact on\nthe dilatometry, which was evidenced by the experimental curves of bar-shaped samples in ambient\nair with a slight amount of wet, as shown in Fig. 2a.”\nWas revised to\n“The hydration-induced expansion is expected to mitigate the dilatometric response, as evidenced\nby the experimental curves of bar-shaped samples in ambient air with a slight amount of moisture,\nshown in Fig. 2a.”\nthermal stress is concentrated, as shown in Fig. 2d. The reduction in TEC value by 34.1% results\nin a 49.8% decrease in von Mise stress at interface edges (from 4.94 to 2.48GPa). The above\nsimulated results indicate the risk of interfacial delamination between the cathode and electrolyte\ncan be significantly reduced by the hydration-induced TEC mitigation.”\nwas revised to\n“This reduction is particularly pronounced at the cathode edges, where thermal stress is\nconcentrated, as shown in Fig. 2d. The 34.1% reduction in TEC leads to a 49.8% decrease in von\nMises stress at the interface edges (from 4.94 to 2.48 GPa). These simulated results suggest that\nhydration-induced TEC mitigation can significantly reduce the risk of interfacial delamination\nbetween the cathode and electrolyte.”\nAt lines 280-285: “Interestingly, the enhancement of fracture strength was obvious at low\ntemperatures, while it became non-significant at higher temperatures (>600 °C). We attribute this\nphenomenon to the chemical hydration that was only significantly below the turning point of\nhydration temperature since as known that perovskite-based cathode materials undergo hydration\nduring the cooling process24, which could be beneficial to optimize grain contact interfaces and thus\nincrease material stiffness.”\nwas revised to\n“Interestingly, the enhancement in fracture strength was more pronounced at lower temperatures,\nwhile it became negligible at higher temperatures (>600 °C). We attribute this phenomenon to\nchemical hydration, which occurs predominantly below the hydration temperature turning point. As\nis well known, perovskite-based cathode materials undergo hydration during the cooling process,\nwhich can optimize grain contact interfaces and, in turn, increase material stiffness.”\nAt lines 374-376: “To fully demonstrate the concept of chemical-induced expansion offset as a\ngenerally effective method for thermal-mechanical enhancement, we further adopted s-BCC-Y as\nthe cathode for both oxygen ion and proton conducting single cells (OSOFC and PCFC) tested in\nambient air using H as fuel.”\n2\nrevised to\n“To fully demonstrate the concept of chemically induced expansion offset as a general method\nfor thermal-mechanical enhancement, we further tested s-BCC-Y as the cathode in both oxygen ion\nand proton-conducting single cells (SOFC and PCFC), using H as the fuel.”\n2\n27\n[Comment 2]: While the mechanical improvements are clear, the paper could discuss\nany potential drawbacks of the Y-doping strategy more extensively. For instance, does\nthe doping affect the long-term stability of the material or introduce any other trade-\noffs in terms of cost or material availability?\n[Response 2]: Thank you for this point. While Y-doping is indeed a promising strategy\nfor improving material mechanical properties, it is potential to have limitations and drawbacks.\nRegarding your suggestion on the cost and material availability, we have addressed it together in\nthe economic analysis in Comment 4.\nFirst of all, B-site Y doping may possibly reduce the surface oxygen exchange coefficient (and\ndecrease electronic conductivity) compared to Co and Fe in the B-site. The electrical conductivities\nof three common Y-doped cathodes, s-BCC-Y, BCFY and BCFYN [Liu, Z. et al. Applied Catalysis\nB: Environmental 346, 123678 (2024)], and the classic cathodes BCC and BSCF were measured,\nas shown in Figure below (Considering the conductivity is unrelated to the core idea of this paper,\nwe did not include it in our manuscript). The three Y-doped cathodes exhibited a significant\ndecrease in conductivity. For example, at 700°C, the electrical conductivity of the Y-doped cathodes\nranged from 2 to 4 S/cm, while BCC and BSCF already reached 40 S/cm.\nFigure: Electronic conductivity of s-BCC-Y, BCC, BCFY, BCFYN and BSCF tested by 4-probe\nmethod from 300 to 600 °C.\nIn addition, Excessive hydration can lead to competition with oxygen adsorption and\ndesorption [Jing, Y., Chemistry of Materials 30, 138-144 (2018)], which can reduce the activity of\nthe OER or ORR processes. Although no such phenomenon was observed in this study, this\ndrawback could potentially arise when excessive Y doping is used. Additionally, proper amounts of\ndoping are beneficial for proton conductivity, which is well known (the introduction of vacancies\nby trivalent cation doping enhances hydration and creates more proton defects). However, excessive\nY doping can also lead to a decrease in proton conductivity, primarily due to the elongation of the\nproton migration path, thus increasing the difficulty of proton hopping [Duffy, J. H. et al. Journal\nof Materials Chemistry A 11, 8929-8938 (2023)]. Therefore, precise control of Y doping is required\nand needs to be traded-off.\nBy the way, it is important to note that in this article, the purpose of Y-doping is to enhance\ncathode hydration, which is the primary mechanism behind the observed mechanical improvements.\ndoping. (the hydrated were much higher than the untreated in all cathodes)\nFig 4. (d) Comparison of fracture strength before and after hydration for BCC, s-BCC-Y, BSCF,\nBSCFY, and s-BCC-Y cathode materials.\n29\n[Comment 3]: It would be interesting to include a comparison to other strategies aimed\nat solving the same problem. How does this approach compare to others in terms of\neffectiveness and practicality?\n[Response 3]: Thank you for your insightful suggestion. Indeed, there are several alternative\nmethods for addressing these thermo-mechanical issues. We have compared and evaluated three\ncommonly used approaches: reducing the Co amount, compositing electrolyte phases, and\nincorporating negative expansion (NTE) materials in the cathode.\nIn contrast to these methods, hydration enhancement through Y-doping offers a new, simple, and\ncost-effective approach. It holds great potential for simultaneously achieving high ORR activity and\nsuperior thermomechanical performance. On the other hand, synthesizing NTE materials is\nchallenging, and excessive incorporation of electrolyte phases or eliminating Co can compromise\nORR/OER activity.\nTo provide a comprehensive comparison, we have included further details in the manuscript, and\nadditional information can be found in SI-1, considering the length of the article:\nAt line 48 -52\n“Early studies have been aimed at solving the interface delamination issue, through the reduction\nof the thermal expansion coefficient (TEC) of cathodes to match with the rigid electrolyte4. For\nexample, developing low Co or Co-free cathode5, compositing with the electrolyte powder6-10, or\nnegative thermal expansive addictive3,11 were all proven to be effective (description in details were\nshown in SI-1).”\nSI-1 A short review of methods for optimizing thermo-mechanical issues\n1.1 Reducing Co amount\nGenerally, it can be observed that classic cathode materials mostly contain Co element. The\ndoping of Co can substantially reduce the cathode's polarization resistance. However, this comes at\nthe cost of reduced Young's modulus and increased thermal expansion coefficient (TEC). Cathodes\nwith lower Young's modulus are more prone to experiencing mass loss and intricate deformation,\nbecause they must withstand a complex flow environment including ambient air, external thermal\ngas and internal oxygen partial pressure gradients. An increased TEC of cathode can indeed result\nin problems such as fracture or delamination from electrolyte, especially when there is a substantial\ndisparity in TEC between the cathode and electrolyte (TEC of 10×10-6 for SDC, 11.2-12.3×10-6 for\nYSZ, 9.3-12.4×10-6 for BZCYYb). Some classic SOFC cathode’s Young’s modulus, TEC and area-\nspecific resistance (ASR) for Co-cathode material were shown in Table S1.\nMany studies are actively addressing these thermos-mechanical issues. Developing of cathode\nmaterials without Co elements is recognized as an effective approach to enhance mechanical\nstrength. LaMn(Fe)O , a first-generation SOFC cathode material, already exhibits high electrical\n3\nconductivity. Furthermore, through doping with elements like Sr, Pr and Nd, materials like Ln\n1-\nSr Mn(Fe)O (Ln=Pr, Nd) 17, 18have achieved even higher electrical conductivity, catalytic activity,\nx x 3\nstructure stability and a lower TEC matching with the electrolyte. Among them, Pr Sr FeO\n0.8 0.2 3\nstands out with polarization resistance of only 0.2038 Ω cm2 at 800℃ and a TEC of 12.1×10-6 K-1,\nmaking it a successful endeavor in the realm of non-Co-based perovskite cathodes. However, in\ngeneral, cathodes doped with Co tend to exhibit better chemical performance compared to those\nwithout Co element. High-performance durable Cathodes like BaCo Fe Zr Y O 19,\n0.4 0.4 0.1 0.1 3\n0.9 0.1 0.9 0.1 3 0.5 0.5 1.5 0.5 5\ncathode to a new level.\nTable S1. Young’s modulus, TEC and area-specific resistance (ASR) for Co-cathode material\nTEC\nE ASR\nMaterial (GPa) (×10-6 K-1) (Ω cm2)\nLa Sr Mn LSM 124(a) 1 12.5 2 0.4 3\n0.8 0.2\nLa Sr Mn Co LSMC 109(b) 4\n0.8 0.2 0.94 0.04\nLa Sr Mn Co LSMC 0.12 3 GDC\n0.8 0.2 0.6 0.4\nLa Sr Mn Co LSMC 13.9 2 900℃\n0.8 0.2 0.5 0.5\nLa Sr Mn Co LSMC 0.063 3\n0.8 0.2 0.2 0.8\nLaMn Co LSC 14.2 2\n0.4 0.6\nBa Sr Fe BSF 110(c) 5 12.4 6 0.613 7\n0.5 0.5 0.5\nBa Sr Fe Co BSFC 63(a) 8 19.95 9 0.361 7\n0.5 0.5 0.8 0.2 SDC\nBa Sr Fe Co BSFC 0.166 7\n0.5 0.5 0.6 0.4 600℃\nBa Sr Fe Co BSFC 0.13 7\n0.5 0.5 0.4 0.6\nBa Sr Fe Co BSFC 0.085 7\n0.5 0.5 0.2 0.8\nLa Sr Fe LSF 165(c) 5 16.3 10 0.25 11\n0.6 0.4\nLa Sr Fe Co LSFC 162 12 (152 13) (c) 17.5 10 0.064 14\n0.6 0.4 0.8 0.2 SDC\nLa Sr Fe Co LSFC 20.3 15\n0.6 0.4 0.5 0.5 600℃\nLa Sr Fe Co LSFC 21.4 15\n0.6 0.4 0.2 0.8\nLa Sr Co LSC 144(c) 5 20.5 10 0.12 16\n0.6 0.4\n1. Young’s modulus test method. (a) 4-point bending. (b) Ultrasonic. (c) Resonance.\n1.2 Compositing electrolyte phase in cathode\nAnother feasible approach is to incorporate electrolyte materials into the cathode material. This\nmethod can enhance the cathode's Young's modulus, increase the oxygen reduction reaction (ORR)\nsites, and further match the TEC between the cathode and electrolyte. Jae-Dong Kim22 and\ncolleagues achieved a successful improvement of the classic LSM cathode by mixing YSZ\nelectrolyte. After adding 40% YSZ, the polarization resistance of LSM-YSZ composite decreased\nfrom 0.38 Ω cm2 to 0.05 Ω cm2 (YSZ, 950°C). Subsequent research revealed that this approach also\noptimized the mechanical performance of the cathode. Under the condition of 40% YSZ mixing,\nthe TEC decreased from 12.3×10-6 K-1 to 11.2×10-6 K-1, and the Young's modulus significantly\nincreased 23. There have been several similar cathode-electrolyte composite enhancement methods,\nsuch as BSCF-SDC24, 25 and LSC, SSC, PBSCF-BZCYYb 26, 27. All of them show significant\nimprovements in performance. However, this approach has limitations and is not suitable for all\ncathode and electrolyte materials. For this method to be effective, the cathode material must exhibit\ngood electronic and oxygen ion conductivity, allowing for the expansion of the triple phase\nboundary (TPB) to facilitate electrochemical reactions. Additionally, the electrolyte material should\nhave high oxide ion conductivity while possessing little to no electronic conductivity to prevent\ninternal cell short circuits 28.\n1.3 Compositing Negative Expansion material in Cathode\nIn addition, Yuan Zhang and Bin Chen 29 have used negative expansion materials instead of\nelectrolyte in the cathode. The resulting SYNC composite (c-SYNC) exhibits excellent\nelectrochemical performance and remarkable thermomechanical stability, providing a new\nperspective for improving the mechanical properties of cathode materials and solve the potential\nrisk of short circuits. Infiltration is another method used for cathode improvement, optimizing the\ncathode's microstructure, and playing a crucial role in enhancing cathode electrochemical\n31\nperformance, especially catalytic activity and long-term stability.30 There is a wide variety of\nsolutions used for infiltration, and Vohs JM31 has discussed the opportunities and challenges of\ncathode infiltration. It points out that the fundamental science behind the infiltration process and the\nprinciples of catalytic integration between materials are still not well understood, making it difficult\nto apply universally across all cathodes.\n32\n[Comment 4]: While the findings are promising, the scalability of the solution\npresented is not extensively discussed. Insights into manufacturing challenges and the\neconomic viability of scaling up the doping process could strengthen the manuscripts\napplicability to the industry.\n[Response 4]: Thanks for the reviewer’s valuable suggestions. In our manuscript, we have\nprimarily emphasized the performance enhancement and stability of our solution, while overlooking\nthe aspects of economic viability and industrial applicability. To solve this problem, we have revised\nthe manuscript and added relevant descriptions, particularly in the discussion section.\nFirstly, the Y-doping strategy proposed in this work will not significantly increase the\nmanufacturing cost or complexity of air electrodes, and its scale-up application can be considered\nstraightforward. Traditional perovskite oxide catalysts are synthesized using solid-state or liquid-\nphase methods (such as co-precipitation, sol-gel, etc.), and the sol-gel method used in this study is\none such liquid-phase approach. It simply involves replacing a portion of Ce nitrate with Y nitrate\nduring the preparation stage, without introducing major changes in processing difficulty or\ncomplexity. Furthermore, no significant adjustments are required for the subsequent heat treatment\ntemperature.\nSimilarly, for large-scale catalyst synthesis methods like co-precipitation, Y-doping can easily be\nimplemented by adjusting the proportions of soluble salts in the raw materials. Therefore, from a\npreparation standpoint, the Y-doping strategy proposed here adds no significant complexity and can\nbe considered a straightforward and scalable approach.\nFrom a cost perspective, Y (a rare earth element) is priced at approximately 40 CNY per 1kg,\nwhich is comparable to other B-site elements such as Co and Ce. As a result, the economic impact\nof incorporating Y-doping is minimal.\nAt Lines 412-417：\n“This work highlights the critical role of cathode fracture in long-term cycling stability and\nproposes a simple, feasible hydration method to enhance the material’s fracture stiffness. We believe\nthis approach will not only improve the long-term stability of SOFCs but also offer an economically\nviable strategy for developing high-durability, high-stability, and high-performance cathodes.\nFurthermore, the scalability of this method is straightforward, with minimal impact on\nmanufacturing costs, making it a promising solution for large-scale industrial applications.”\n33\n[Comment 5]: The longevity of the cathode material is crucial for commercial\napplication. The paper could explore the long-term performance of the cathodes under\ncontinuous operation, beyond cyclic testing.\n[Response 5]: Thank you very much for pointing out this issue. We believe that cyclic stability\nin SOFCs is a more thorough evaluation of stability compared to continuous stability, as the interface\nstress caused by temperature fluctuations presents a more severe challenge to its durability. This is\nwhy we focused more on cyclic stability in our previous manuscript and overlooked continuous\nstability.\nHowever, we also recognize that continuous stability is essential for a high-performance cathode,\nso we have added tests to evaluate the isothermal stability of the material. We monitored the ASR\nchange of the cathode under 600°C for 300 hours. The Y-doped cathode exhibits improved stability\nthan BCC. As seen in Fig. S19, after 300 hours, the ASR of BCC increased from 0.325 Ω cm² to\n0.71 Ω cm², showing a 1.28 mΩ cm²/h degradation, while s-BCC-Y’s polarization resistance was\nrelatively smaller, increasing from 0.249 Ω cm² to 0.42 Ω cm², a 0.58 mΩ cm²/h degradation.\nNotably, we would like to emphasize to reviewers and readers that the durability test over 300\nhours is not negligible in s-BCC-Y, this could be caused by many factors, including stability of the\nelectrolyte (BZCYYb) the test environment from lab to lab. To make fair comparison, we also\ncompare to other classic cathode candidates that are tested in the same condition. that are viewed as\nvery promising cathode in the reports, such as BSCF shown a 1.3 mΩ cm²/h degradation at 600 °C\n[Han, L. et al. Chemistry of Materials (2024)], C/H-BSCF with a decay rate of 0.45 mΩ cm2/h at a\nrelatively low temperature 550 °C [Liu, Z. et al. Nature Communications 15, 472 (2024)]. This\nindicates that the cathode stability of s-BCC-Y has development potential.\nWe have also included the result in SI-14.\nFig. S19 Time-dependent ASR values of BCC and s-BCC-Y composite electrode based on\nBZCYYb electrolyte at 600°C in 3% H O-air.\n“After 300 hours, the Rp of BCC increased from 0.325 Ω cm² to 0.71 Ω cm², showing a\ndegradation of 1.28 mΩ cm²/h, while the decay of s-BCC-Y’ s is only 0.58 mΩ cm²/h (SI-14, Fig.\nS19).” Details were shown in SI-14.\n34\n[Comment 6]: In the method section, the terminology \"atropurpureus gel\" appears\nundefined or nonstandard. Should this refer to a distinct variety of gel, it would be\nprudent to elucidate its composition or an appropriate amendment to its designation.\n[Response 6]: We sincerely apologize for the improper use of the term “atropurpureus gel”, which\nwas unclear and nonstandard. We meant to refer to a gel with a dark purple color. We have corrected\nthis term in the manuscript and clarified that it refers to a sol-gel precursor used in the synthesis\nprocess.\nAt lines 438：\n“A sol-gel precursor was obtained after the removal of water in the solution by evaporation,\nwhich was pre-decomposed at 180°C for 8 hours to obtain the precursors.”\n35\n[Comment 7]: Given the complications associated with the performance of 8YSZ\nelectrolytes at intermediate temperatures, how do the introduction of chemical\nhydration and the improvement in mechanical stiffness and fracture strength, as\ndemonstrated by the stiffness-enhanced air electrode s-BCC-Y, impact the thermo-\nmechanical durability of cathodes in solid oxide fuel cells (SOFCs), particularly with\nproton ceramic fuel cells (PCFCs)? Can these strategies be applied to optimize other\nintermediate-temperature SOFC materials, such as doped ceria?\n[Response 7]: We sincerely appreciate your careful review of our work. The point you raised is\nexactly what we intended to convey. Regarding the impact of hydration on mechanical stiffness and\nfracture strength, we have highlighted this in the schematic diagrams in Fig. 1a and Fig. 5b. We\nrealize, however, that our explanation might not have been clear enough.\nRegarding the first question, we would like to clarify that the increased stiffness of s-BCC-Y can\nstill be maintained under SOFC operating conditions. Like PCFCs, SOFCs are exposed to a water\nvapor environment on the air electrode side, which supports hydration. This hydration-induced\nstiffness enhancement is expected to improve the thermomechanical durability of SOFCs, as\nconfirmed by durability tests conducted on the cathodes in both SOFCs and PCFCs (see Fig. 5e, f).\nFor the second question, we have previously tested Y doping in various classic air electrode\nmaterials, i.e. BSCF perovskite, which shows the same positive results. Regarding doped CeO ,\n2\nprevious studies such as [Fu, Y.-P. Ceramics International 35, 653-659 (2009)] have shown that Y\ndoping improves fracture toughness, attributed to a reduced grain size. We agree with this finding,\nas grain size and grain boundary bonding are very key to mechanical properties in dense CeO\n2\nelectrolytes. However, it is unclear whether hydration will significantly affect the mechanical\nstrength of CeO . Due to the high density and tight grain boundaries of CeO , we believe that\n2 2\nsignificant hydration and chemical expansion may have a minimal effect. As this falls outside the\nscope of our current work, we have decided not to fully introduce this discussion in the main text to\nmaintain focus.\nAt Lines 382-388：\n“Both cells were able to sustain a consistent power output over an extended duration. In Fig. 5e,\nthe cell output shows a voltage drop from 0.797 V to 0.758 V after 380 hours, corresponding to a\n0.103mV/h degradation. In Fig. 5f, the voltage decreases from 0.753 V to 0.71 V, representing a\n0.153 mV/h degradation. This strongly suggests that the s-BCC-Y composite cathode, improved by\nstiffness enhancement, exhibits remarkable performance in single fuel cells. This also indicates that\nthe cathode, with enhanced hydration demonstrates good durability in both SOFC and PCFC\nconfigurations.”\n36"}, {"round": 2, "reviewer_comments": {"reviewer_1": "Revisions were reasonably done and now the manusript is suitable for publication.", "reviewer_2": "In this revision, authors addressed most issues raised by the reviewer, thus it is believed that this manuscript can be published in this journal without further revision.", "reviewer_3": "COMMENTS TO THE AUTHOR(S)\nTitle: Hydration-induced stiffness enabling robust thermal cycling of high temperature fuel cells cathode.\nManuscript No.# NCOMMS-24-47288A\nThe manuscript presents an important study focusing on the enhancement of mechanical durability in high-temperature fuel cells through yttrium doping of BaCo0.7Ce0.3O3 perovskite. The authors have carried out a comprehensive set of evaluations demonstrating notable improvements in cathodic stiffness, fracture strength, and overall fuel cell performance. Such advancements are significant for the future of SOFCs and PCFCs. However, certain aspects of the manuscript could be improved to increase accessibility, address potential drawbacks, provide comparative analysis, discuss scalability issues, consider long-term operational performance, clarify terminology, and explore the broader applicability of the research findings.\nAuthor's Responses to Comments:\nThe authors have made a commendable effort to restructure the manuscript by modifying section titles for clarity and revising the content for conciseness and readability. This response is satisfactory, showing that they took the reviewer's feedback seriously and worked to make the manuscript more accessible.\nThe authors have systematically addressed all the reviewer comments, improving both the substance and presentation of the manuscript. No further major issues are evident, and it appears that the manuscript is suitable for publication upon the completion of these revisions."}, "author_response": ""}]

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