In situ formation of a melt-solid interface towards stable oxygen reduction in protonic ceramic fuel cells

doi:10.21203/rs.3.rs-3450609/v1

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In situ formation of a melt-solid interface towards stable oxygen reduction in protonic ceramic fuel cells

https://doi.org/10.21203/rs.3.rs-3450609/v1

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Protonic ceramic fuel cells (PCFCs) are one of the promising routes to generate power efficiently from various fuels at economically viable temperatures (500–700°C) due to the use of fast proton conducting oxides as electrolytes. However, the power density and durability of the PCFCs are still limited by their cathodes made from solid metal oxides, which are challenging to address the sluggish oxygen reduction reaction and susceptibility to CO₂ simultaneously. Here, we report an alternative approach to address this challenge by developing a new melt-solid interface through the in situ alkali metal surface segregation and consecutive eutectic formation at perovskite oxide surface at PCFC operating temperatures. This new approach in cathode engineering is successfully demonstrated over a lithium and sodium co-doped BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3-δ perovskite as the model material. Our experimental results unveil that the unique in situ formed melt-solid surface stabilises the catalytically active phase in bulk and promotes catalytically active sites at surface. Our novel engineered melt-solid interface enhanced the stability of the cathode against poisoning in 10% CO₂ by a factor of 1.5 in a symmetrical cell configuration and by a factor of more than two in PCFC single cells.

Physical sciences/Materials science/Materials for energy and catalysis

Physical sciences/Materials science/Materials for energy and catalysis/Electrochemistry/Fuel cells

Protonic ceramic fuel cells (PCFCs) with fast proton-conducting ceramics provide a promising cost-effective route for high power generation directly from a wide range of fuels (e.g. hydrogen, hydrocarbons, and ammonia) at intermediate temperatures (500–700°C)^1–7. However, the power density of PCFCs is limited by the lack of catalytically active and stable cathodes, where the oxygen reduction reaction (ORR) takes place and leads to the formation of water as the product. This unique oxygen reduction reaction at PCFCs necessitates the active cathode to be proton, oxygen ion, and electron (hole) conductive⁸. State-of-the-art PCFC cathodes typically have a perovskite-type structure with formula ABO₃, where the A site contains alkaline-earth elements such as Ba and Sr, and the B site contains transition metals such as Co and Fe, forming compositions such as BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZY)⁹, BaCo_0.8Ta_0.2O_3−δ¹⁰_, Sr₂Sc_0.1Nb_0.1Co_1.5Fe_0.3O_6−δ¹¹, and SrCo_0.8Fe_0.15Zr_0.05O_3−δ¹².

Despite their mixed conductivities, most of these typical PCFC cathodes are highly susceptible to poisoning by CO₂ at PCFC operating temperatures as a result of the presence of alkaline-earth elements, which react to form carbonates,¹³ leading to irreversible loss of active cathode material. Even trace amounts of CO₂ (~ 400 ppm) in ambient air can cause rapid cathode degradation and power density loss. For example, Qi et al. report that 6.4 at. % Ba in BCFZY converts to BaCO₃ after 24 h exposure to 25°C air with 380 ppm CO₂ ¹⁴, and Li et al.¹⁵ report that 1% CO₂ in air caused a 59% activity loss by the BCFZY cathode during 120 h operation at 700°C. It is well known for fuel cells operated at intermediate temperatures, such as solid oxide fuel cells, that loss of cathode activity is orders of magnitude worse at reduced operating temperature as a consequence of the relatively higher stability of carbonate phases^16–18.

Common strategies to address the susceptibility of these perovskites to CO₂ rely on modification of the cathode bulk or surface to weaken the binding strength of the perovskite with CO₂ or include additional solid phases that are not CO₂ reactive. The cathode can be modified by reducing the basicity of the A-site cation^15,19 and/or enhancing the strength of the B-O bond. However, such modifications might exert a negative impact on the material's proton conductivity, which is necessary for water formation^20,21. The inclusion of secondary phases that are not CO₂ reactive has also thus far negatively impacted ORR activity of the original cathode because the CO₂-inert phase is usually less ORR active than the primary cathode.^18,21,22 These approaches all lower the PCFC power output. Therefore, it is difficult but imperative to develop a new strategy that can realize high cathode CO₂ stability without sacrificing the ORR activity.

Herein, we pursue a new strategy to improve the CO₂ tolerance of the BCFZY material with negligible loss of ORR catalytic activity by inducing the formation of a molten-solid interface. We first introduce Li⁺ and Na⁺ to the BCFZY A-site to produce the single-phase perovskite Ba_0.95(LiNa)_0.05Co_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZYLN). In the presence of CO₂, the alkali metal cations in BCFZYLN co-segregate at the perovskite surface and react in situ with CO₂, forming a thin molten adlayer partially covering the perovskite surface. Our results show that this in situ developed interface effectively suppresses the carbonation of Ba in BCFZY and stabilises the ORR by re-liberating active sites and extending the active surface. Consequently, the unique melt-solid interface more than doubles the cycling stability of the PCFC single cell in CO₂ compared with the unmodified cathode. These findings point to a new strategy for the modulation of solid/liquid interfacial reactions to advance a broad range of high-temperature catalyst materials.

Crystal structure and chemical composition

BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZY) and Ba_0.95(LiNa)_0.05Co_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ (BCFZYLN) were synthesised via a typical sol-gel route. We then characterised the crystal structure of these two materials quenched from 600°C using synchrotron X-ray powder diffraction (SXRD) and neutron powder diffraction (NPD). Diffraction data (Fig. 1) confirm that both BCFZY and BCFZYLN are predominantly cubic perovskites with a \(\:Pm\stackrel{-}{3}m\) space group. The higher resolution SXRD data, however, reveal that this phase contains a distribution of phases with slightly different lattice parameters, and we were able to describe the data using two phases, each with a slightly different lattice parameter, termed BCFZY-1 and − 2; BCFZYLN-1 and − 2 (see Supplementary Tables 1 and 2). This roughly binary distribution of phases with differing lattice parameters suggests some compositional segregation, which is consistent with that reported for other Zr-containing perovskite oxides^23,24 and ascribed to segregation between Zr and other B-site cations such as Co and Fe²⁵. Structure refinement revealed a smaller lattice constant (4.10999(4) − 4.11833(3) Å) for the Li/Na co-doped sample than the undoped material (4.11758(5) − 4.12385(5) Å), as consistent with the smaller Li⁺ (0.90 Å) and Na⁺ (1.16 Å) dopants relative to Ba²⁺ (1.49 Å)²⁶. NPD data also reveal that the Li⁺ A site occupancy in BCFZYLN is 2(1)% and 3.9(6)% for Na⁺, which confirms that the alkali cations were doped into the perovskite lattice.

We then studied the chemical composition of the perovskite materials using inductively coupled plasma-optical emission spectrometry (ICP-OES). The results showed that BCFZY has an average formula of BaCo_0.373Fe_0.404Zr_0.122Y_0.102O_3−δ, while BCFZYLN is Ba_0.937Li_0.027Na_0.035Co_0.359Fe_0.384Zr_0.099Y_0.158O_3−δ, consistent with the nominal target compositions and powder diffraction data.

Oxygen non-stoichiometry, denoted as δ, quantifies the content of oxygen vacancy in the perovskites. Supplementary Fig. 1 compares the oxygen non-stoichiometry between both samples as a function of temperature. At 600°C, more oxygen vacancies are created in the co-doped (δ ~ 0.73) than in the undoped sample (δ ~ 0.63), which is consistent with the reported δ in BCFZY at 600°C.^27–30 The increased oxygen vacancy concentration in BCFZYLN is possibly a result of the lower valence of Li⁺ and Na⁺ compared with Ba²⁺, which requires fewer O²⁻ ions to maintain the material’s charge neutrality.

Additionally, the hydration capability of the perovskites determines the ORR electrochemical pathway of the PCFC cathode, and we therefore examined samples exposed to dry or humid air using thermogravimetric analysis (TGA). As shown in Supplementary Fig. 2, the co-doped sample exhibits a water uptake of 0.18 wt.% at 300°C, slightly higher than that of the sample (0.14 wt.%), indicating that co-doping with alkali metals enhances hydration capacity of the material.

Cathode ORR activity and CO₂ stability

We then evaluated the ORR activity of the synthesised perovskite materials in air containing 3 vol.% moisture at 450–600°C in symmetrical cells, where cathode materials are deposited onto both sides of a dense protonic conducting electrolyte disk. The ORR activity is characterised by the area-specific resistance (ASR), where a low ASR indicates a high ORR catalytic activity. The comparison of the ASR values shown in Fig. 2a shows negligible differences in ORR activities between the undoped and co-doped cathode in humid air. For example, the BCFZY cathode exhibited an ASR of 0.19 ± 0.02 Ω cm² at 600°C and 3.8 ± 0.2 Ω cm² at 450°C in humid air, in agreement with previously reported values for the same material.^9,27,31,32 Similarly, BCFZYLN exhibits an ASR of 0.20 ± 0.03 at 600°C and 3.6 ± 0.1 Ω cm² at 450°C, suggesting that Li/Na co-doping has no impact on the ORR activity of the cathode in humid air, despite its higher oxygen vacancy and water uptake.

We studied the stability of the BCFZY and BCFZYLN cathodes against CO₂ at 600°C by conditioning symmetrical cells in flowing humid air containing 10% CO₂ for 1 h and monitoring the ASR over time. Since the electrolyte material, BaZr_0.1Ce_0.7Y_0.1Yb_0.1O_3−δ, might also be CO₂ reactive and potentially influence the overall cell performance,^3,20 we first studied the effect of CO₂ on both the cathode and electrolyte in symmetrical cells by measuring their ohmic and polarization resistance during CO₂ conditioning. As shown in Supplementary Figs. 3 and 4, the ohmic resistance of symmetric cells with BCFZY and BCFZYLN cathodes shows negligible degradation, indicating that the observed performance changes are primarily related to the reaction between CO₂ and the cathode.

As shown in Fig. 2b, the pristine BCFZY cathode degrades quickly in the presence of CO₂, as evidenced by an increase in ASR from 0.19 ± 0.02 to 0.9 ± 0.1 Ω cm² at a rate of 7.4 ± 0.2 × 10^− 3 Ω cm² min^− 1. Such cathode degradation is irreversible, and the ASR of the undoped cathode only recovered to 0.32 ± 0.03 Ω cm² following the replenishment of the test chamber with humid air free from CO₂ for 30 min.

In contrast, the co-doped BCFZYLN cathode is more stable under CO₂ conditioning than the BCFZY cathode. As shown in Fig. 2b, the ASR of the co-doped cathode increases from 0.20 ± 0.03 to 0.73 ± 0.08 Ω cm² at a degradation rate as low as 4.9 ± 0.3 × 10^− 3 Ω cm² min^− 1 under the same condition, representing more than 33% improvement of stability as compared to the undoped cathode. This performance is also among the best-reported stabilities against CO₂ under similar conditions in the literature. (see Supplementary Table 3) After CO₂ removal, notably, the co-doped cathode ASR can recover back to 0.26 ± 0.03 Ω cm², 20% more active than that of the undoped cathode. Despite their minor impact on the cathode ORR activity in CO₂-free air, co-doping of alkali metals effectively enhances both the cathode ORR activity and stability in the presence of CO₂. Such an improvement in both activity and stability could not be easily achieved through previously reported methods in modifying the cathode bulk or surface. Interestingly, the co-doped cathode is also superior to the Na- and Li- single-doped counterparts (see Supplementary Figs. 5 and 6), implying a potential synergistic effect of the co-doping strategy in resisting CO₂ poisoning.

At a higher operating temperature at 700°C, the co-doped cathode further improves ORR activity by ~ 40% in humid air and stability by a factor of nearly five in the presence of CO₂ when compared with the undoped cathode (see Supplementary Fig. 7). Meanwhile, the degradation rate of the co-doped samples can be further minimised down to 0.9 ± 0.1 × 10^− 3 Ω cm² min^− 1. Such improvement is likely a consequence of the destabilised carbonate phase at elevated temperatures. In other words, a high CO₂ tolerance at reduced operating temperature is highly sought after for PCFCs that aim for low to intermediate-temperature regimes.

We then compared the performance of the undoped and co-doped materials as the cathode in a single cell configuration, including stability in the presence of CO₂. A single cell is composed of a Ni-BZCYYb anode ~ 300 µm thick, a cathode ~ 10 µm thick, and a dense ~ 20–30 µm thick electrolyte separating the fuel from air (see also Supplementary Fig. 8). We fabricated two single cells with BCFZY (cell #1) and BCFZYLN (cell #2) as cathodes, with Fig. 2c and d showing the current-voltage-power (I-V-P) curves of these cells. In the absence of CO₂ at 600°C, both cells generated similar power densities, with a peak power density of 514 mW cm^− 2 for cell #1 and 526 mW cm^− 2 for cell #2. Both cells produced relatively stable and similar cell voltages under a constant current density (300 mA cm^− 2) in the absence of CO₂ in the first 38 h. This similarity in cell performance is a result of the similar cell configurations and the ORR catalytic activity of BCFZY and BCFZYLN cathodes observed in Fig. 2c.

More importantly, when cycling the cathode reaction environment with and without the presence of 10% CO₂ (Fig. 2e), we find that cell #2, containing the co-doped cathode, outperforms cell #1 in resisting CO₂ poisoning during three cathode atmosphere changes. Particularly in the third cycle, cell #2 sustained a cell voltage above 0.8 V, more than twice that of cell #1 (< 0.4 V) in the presence of CO₂. Further, cell #2 regenerates a peak power density of 386 mW cm^− 2, far higher than that achieved by cell #1 (194 mW cm^− 2), after the 3rd cycle once CO₂ was removed. These single-cell results further demonstrate the effectiveness of Li/Na co-doping to boost the PCFC power output and stability of the cathode in the presence of CO₂.

The origin of CO₂ resistance in BCFZYLN

The role of co-doped alkali cations in suppressing carbonation

To probe the underlying mechanism behind the notable stability improvement achieved by Li/Na co-doping, we first studied the crystal structures of the samples by analysing the laboratory-based X-ray powder diffraction (XRD) data of the undoped BCFZY, Li/Na co-doped, and single-doped analogues. Before the XRD characterisation, these samples were treated under flowing air containing 10% CO₂ at 600°C for 1 h and then quenched to room temperature (Supplementary Figs. 9 and 10 with Rietveld refinement results shown in Supplementary Tables 6–9).

Lab XRD results reveal that CO₂ treatment transformed 85.8(6)% of the BCFZY phase from a cubic perovskite structure with \(\:Pm\stackrel{-}{3}m\) symmetry to orthorhombic BaCO₃ with Pmna symmetry. BaCO₃, formed from the reaction between BCFZY and CO₂, is limited in both ionic- and hole-conductivities and is primarily responsible for cathode degradation in both symmetrical and single cells in Fig. 2.

In contrast, co-doping of Li/Na significantly suppresses BaCO₃ formation, and our diffraction analyses show that the majority of the perovskite phase sustains after 1 h CO₂ exposure, with only 4.7(1)% BaCO₃ formed. Similarly, as shown in Supplementary Fig. 10, single-doping of Li or Na also limits BaCO₃ formation in CO₂, but not to the same extent as co-doping.

Since alkali metal cations replace Ba in BCFZY, it is important to investigate the impact of Ba deficiency on the structural stability of these alkali cation-doped cathodes. Hence, we applied SXRD to characterise the phase evolution of BCFZY, Ba-deficient BCFZY (B_0.95CFZY), and BCFZYLN in capillaries fed with 3 vol.% humid CO₂ at a flow rate of 2 mL min^− 1 at 600°C. The SXRD results further confirm the dominant influence of co-doping over Ba deficiency in stabilising the perovskite phase under CO₂ conditioning. For example, the co-doped sample has only 0.20(3)% BaCO₃ (see Fig. 3a and Supplementary Fig. 11), much lower than the 6.6(2) wt.% in B_0.95CFZY and 3.8(2) wt.% in BCFZY after 2h CO₂ conditioning.

The CO₂ tolerance of the samples was also evaluated by comparing the CO₂ desorption temperatures from the temperature-programmed desorption (TPD) profiles in Supplementary Fig. 12, where a high desorption temperature indicates a strong tendency for the material to interact with CO₂. The comparison of the CO₂ desorption temperatures, as shown in Fig. 3d and Supplementary Fig. 12, implies that the co-doped sample has a weaker interaction with CO₂ than either single or undoped samples, further confirming the effectiveness of co-doping in suppressing CO₂ reactivity.

Our previous studies reported that single-doping alkali metal cations can enhance perovskite oxide stability against CO₂ in dry air.^34,35 Interestingly, the Li and Na co-doped material shows enhanced CO₂ desorption and improved structure stability against CO₂ relative to its single-doped counterparts, which aligns well with symmetric cell results. The superior performance of Li and Na co-doping suggests that these two alkali metal cations might have a synergistic effect in boosting the CO₂ tolerance of BCFZY.

Surface segregation of alkali dopants in the presence of CO₂

The synergistic effect of the co-doping should be related to the segregation of alkali cations from the bulk to the surface in response to CO₂ conditioning. To examine this hypothesis, we characterised the crystal structures and compositions of the CO₂-treated BCFZYLN in detail using SXRD and NPD (Fig. 3b and c, and Supplementary Tables 10 and 11). The results show that the lattice occupancy is lowered to only 0(2)% for Li and 0(5)% for Na in the BCFZYLN lattice after CO₂ conditioning, which is an indication of loss of the alkali dopants from the bulk of the sample.

The loss of the alkali cations can only be enriched and form carbonate with CO₂ at the surface due to the low volatility of the alkali cations and their tendency to react with CO₂ at intermediate temperatures under investigation.^36,37 Therefore, we studied the chemistry of the samples using spectroscopic analyses. The Fourier transform infrared (FTIR) spectra, as shown in Fig. 3e, reveal the presence of adsorbed CO₂ and BaCO₃ formation across all samples from the two featured adsorption bands at ~ 1480 cm^− 1 and 860 cm^{− 1 38}, but only indicate the formation of Na₂CO₃ species over the CO₂-treated co-doped sample at ~ 892 cm^{− 1 39}. However, laboratory FTIR is unable to distinguish the carbonate species between Li₂CO₃ and BaCO₃ due to their overlapping features at ~ 1500 cm^− 1 and 850 cm^− 1.³⁹

We then turned to high-resolution X-ray photoelectron spectroscopy (XPS) and successfully identified the deconvoluted peaks for O 1s (289.6 eV) and C 1s (531.7 eV) at the surface(see supplementary Figs. 13 and 14), corresponding to the Li₂CO₃ species in the CO₂-treated BCFZYLN^40,41 However, the detection of Na₂CO₃ in these samples using XPS is challenging as a result of the close proximity of O 1s (~ 531.0 eV) and C 1s (~ 289.3 eV) peaks for Na₂CO₃ to those (O 1s at ~ 531.2 eV and C 1s at ~ 289.4 eV) for BaCO₃.^40,41

Therefore, to confirm our laboratory FTIR and XPS results, we applied synchrotron-based terahertz adsorption spectroscopy (THz) to characterise our samples in the far infrared region between 100 and 600 cm^− 1. The THz results, shown in Supplementary Fig. 15–18, unveil a band at ~ 435 cm^− 1 that is characteristic of Li₂CO₃ and Na₂CO₃ and only found in the spectrum of the co-doped sample. This result further supports our hypothesis that both Li and Na segregate from the bulk to the surface to form carbonates under CO₂ conditioning.

More interestingly, we could not identify any alkali metal carbonate phase even from the high-resolution SXRD data (Fig. 3b). The only plausible explanation for this finding is that these alkali metal carbonates could be in an amorphous phase, which cannot be detected by diffraction techniques. Noting that the Li₂CO₃-Na₂CO₃ binary system has an eutectic temperature of 498°C^42,43,44, these alkali metal carbonates can be present in melts instead of crystalline solids at temperatures beyond 500°C and form an amorphous phase while being quenched to room temperature.

Mixed alkali carbonate melt formation

To examine this hypothesis, we conducted differential scanning calorimetry (DSC) characterisation over CO₂-treated samples. If the melt is present in the samples, the DSC should be able to capture it by showing the change of the heat flow at the glassy transition of the carbonate phases. As shown in Fig. 3f, importantly, we identified an endothermic feature at 504°C only in the DSC data of the CO₂-treated BCFZYLN, consistent with the reported eutectic temperature (498°C) for the Li₂CO₃-Na₂CO₃ binary system.^42,43,44 This feature can be a result of the heat absorbed by the transition of alkali metal carbonates from amorphous to melt. In contrast, this endothermic feature is absent in the rest of the samples, further confirming the presence of Li/Na carbonate melts in the CO₂-treated co-doped samples.

We then compared the morphology of the cathode surface before (see Supplementary Fig. 19) and after CO₂ treatment using both scanning and transmission electron microscopy (SEM and TEM, respectively). Supplementary Fig. 19 indicates smooth surfaces for both BCFZY and BCFZYLN samples before CO₂ conditioning, suggesting that both samples are likely phase pure, as consistent with SXRD results. Further results from DSC characterisation show no evidence of a molten phase at the BCFZYLN surface before CO₂ treatment. These results illustrate that reaction with CO₂ is necessary for the formation of the melt at the co-doped sample surface at elevated temperatures. The absence of the melt could also well explain the identical ORR activities for both co-doped and pristine samples observed in Fig. 2a under CO₂-free air.

TEM (Fig. 4a and b) reveals a secondary phase ~ 5 nm thick in direct contact with the crystalline bulk of the CO₂-treated co-doped sample. The fast Fourier transform (FFT) of this (Supplementary Fig. 20c) along with selected area electron diffraction (SAED) (Supplementary Fig. 21) indicate that this secondary phase is amorphous. This secondary amorphous phase is likely formed through quenching of the Li/Na carbonate melt, with this secondary phase not present in the other samples (Fig. 4a, d, and e).

Lower magnification SEM images show that this melt partially covers the surface of the co-doped sample (Fig. 4c), and that it has a microstructure that is substantially different from that of the surface of the treated undoped material where isolated crystalline nanoparticles, potentially of BaCO₃ as identified through SXRD, are visible (Fig. 4f). A similar amorphous phase is also discernible in the BCFZYLN single cell cathode post operation, as shown in Supplementary Fig. 22. Although elemental mapping and line scans of TEM data using energy dispersive spectroscopy (EDS) were conducted of these surfaces, the results were inconclusive due to the large background signal for Na (Supplementary Fig. 23 to 25).

Our results support that co-segregation of Li/Na to the perovskite surface enables the formation of a mixed carbonate melt as identified in Figs. 3 and 4. Such A-site cation segregation is a common phenomenon in many perovskite materials at elevated temperatures,^45–51 where host/dopant cation size mismatch is widely reported as a key driving force^45,50,51 for cation surface segregation. In our material, the Shannon ionic radius is 0.92 Å for Li⁺ and 1.39 Å for Na⁺, much smaller than for Ba²⁺ (1.61 Å).²⁶ Therefore, this size mismatch likely increases the elastic energy and causes Li and Na to migrate from the bulk to the surface.

Apart from cationic size mismatch, the atmosphere also plays a role in cation segregation and melt formation. Our prior work³⁴ studying alkali metal single-doped SrFeO_3−δ perovskite oxides revealed that the presence of acidic CO₂ induces the surface segregation of basic alkali metals due to the carbonation reactions between alkali metals and CO₂. Hence, the observed co-segregation of the Li/Na surface segregation from BCFZYLN could arise from this mechanism in addition to the dopant/host size mismatch. Therefore, the segregated alkali metals make it possible for the observed formation of a Li₂CO₃-Na₂CO₃ eutectic melt adlayer in situ at the perovskite surface through their reaction with CO₂ at temperatures above 500°C. Further, the low Li/Na substitutional level limits the surface availability of segregated alkali metal species, resulting in a hybrid melt-solid surface structure as shown in Fig. 4c, which cannot be easily achieved by other conventional techniques such as wet impregnation ⁵² or mechanical mixing ⁵³. Such a melt-solid hybrid cathode surface, as identified from this study, could be the origin of the significant improvement of CO₂ tolerance by the BCFZYLN cathode at intermediate temperatures as shown in Fig. 2 and Supplementary table 1.

We envisage that the in situ formed melt-solid surface structure offers an opportunity to re-liberate ORR active sites and maximise surface oxygen-ion migration. In general, CO₂ deteriorates the ORR activity of perovskite oxides first through competitive adsorption at active oxygen adsorption sites. The absorbed CO₂ can further lead to the irreversible carbonation of the A-site alkaline-earth metal ions, resulting in the depletion of the catalytically active phase, as shown in Fig. 3a. Conventional material design strategies to address CO₂ susceptibility usually rely on either weakening the CO₂ binding strength of the material or introducing additional CO₂-insensitive solid phases^16,21,22,54. However, both these strategies fail to counter CO₂ poisoning while sustaining a high cathodic activity under 600 °C. This limitation is due to the fact that the solid surfaces resisting CO₂ adsorption or carbonation are usually compromised in oxygen adsorption, dissociation, and surface oxygen ion migration processes.

In the present work, as illustrated in Fig. 5, the in situ formed surface eutectic melt provides a pathway to transport CO₂ away from the cathode particle and also dissociate oxygen at the surface. Mixed alkali carbonate melts are known for their fast oxygen ion conduction at intermediate temperatures ^55–57 and have been widely used as electrolytes in molten carbonate fuel cells ⁵⁷. Their ionic conductivities (e.g., 1.73 S m^− 1 at 600°C for Li/Na carbonate melt) ^55,56 are orders of magnitude higher than the proton/oxygen-ion conductivity of solid BCFZY (e.g., 0.16 S m^− 1 at 600°C). The charge carriers in the carbonate melt are mobile carbonates, which undergo dynamic reactions that dissolve and release oxygen ions by absorbing and desorbing CO₂, respectively. Through this mechanism, the in situ formed Li/Na carbonate melt can suppress the formation of undesirable BaCO₃ by diverting adsorbed CO₂ away from the perovskite solid surface, re-liberating the ORR active site. The experimental results presented in Fig. 3a-3d confirm a much higher stability of the alkali-metal co-doped samples in the presence of CO₂ compared to their parent oxide and Li/Na single-doped samples that do not have an alkali carbonate melt.

Further, the melt-solid interface extends the active surface for oxygen reduction and water formation across the whole cathode. Oxygen reduction is known to occur at the boundaries between the gas phase that enables oxygen supply and water vapour release, the electron-conducting phase to supply electrons, and the proton/oxygen-ion conducting phase to transport protons and produce oxygen ions.^8,11,58 In conventional solid cathodes, CO₂ mainly poisons the electron/ion-conducting phase by out-competing for oxygen vacancies and destroys the electron- and proton-conducting phases that require a stable perovskite structure and contain alkali-earth elements. When introducing melt at the surface, instead, we can reconstruct abundant CO₂-resistant active sites made from interfaces between the stable perovskite surface and the superior ion-conducting melt. In addition, the oxygen ions can be quickly transported in the surface melt to sites where protons are available for water production. In this way, the melt-solid provides a surface structure facilitating a high density of sites for combining proton and oxygen ions, which is deemed a key step to boosting the performance of protonic fuel cell cathodes.^{31,32,59–61}

In summary, we report a cathode engineering strategy to increase the stability of the PCFC cathodes without compromising activity by developing a new melt-solid interface. By introducing Li⁺ and Na⁺ into the model BaCo_0.4Fe_0.4Zr_0.1Y_0.1O_3−δ material we show that these alkali metal cations tend to co-segregate from bulk to the cathode surface and react with CO₂ to form a nanoscale surface layer under PCFC operating conditions. The melt-solid cathode effectively preserving the perovskite structure and oxygen-reduction reaction catalytic activity in the presence of CO₂. The in situ formed melt-perovskite interface facilitates stable and fast transport of CO₂, electrons, and ions across the cathode surface, which is important for a stable oxygen reduction reaction at practical PCFC operating conditions. This new concept in cathode design achieves significant improvement in cathode stability against 10% CO₂ in both symmetrical cell and single cell configurations. The findings of this work are expected to offer an alternative route to address the challenges in developing high-temperature solid-state electrodes by decoupling the catalytic activity and durability through in-situ phase segregation and solid-melt interface formation.

Acknowledgements

We appreciate technical support from the Centre for Microscopy and Microanalysis (CMM) at the University of Queensland, the Australian Centre for Neutron Scattering (ACNS) for beamtime under proposal MI13525 and the Australian Synchrotron for beamtime under proposals PDR19011, PDR19087, M19209, M20095 and M20988. This work is financially supported by the Australian Research Council (ARC) DP200101397. DF acknowledges the financial support from the UQ Graduate School scholarship. ML acknowledges the financial support from ARC (DE230100637). SG acknowledge the financial support from the Natural Science Foundation of Jiangsu Province (No. BK20231003).

Sample preparation

All samples were prepared using a sol-gel method with stoichiometric salts. For example, BCFZYLN was prepared from salts of Ba(NO₃)₂ (99.0% Sigma-Aldrich), Co(NO₃)₂·6H₂O (99%, Sigma-Aldrich), Fe(NO₃)₃·9H₂O (98%, Sigma-Aldrich), ZrO(NO₃)₂·xH₂O (99.99% Sigma-Aldrich), Y(NO₃)₃·6H₂O (99.8%, Sigma-Aldrich), LiNO₃ (99.0%, Sigma-Aldrich), and NaNO₃ (99.0%, Sigma-Aldrich), which were dissolved in deionized water. The citric acid (99.5%, Sigma-Aldrich) was then added to the solution along with ethylenediaminetetraacetic acid (EDTA, 99.4%, Sigma-Aldrich) was first dissolved in ammonia solution (25%, Sigma-Aldrich). The molar ratio of EDTA: citric acid: total metal ions was controlled to be 3 : 3 : 2, and the final pH was increased to about 10 by adding ammonia solution.

The solution was dried using a hot plate at approximately 80°C overnight to obtain a gel phase which was then fired at 260°C using a fan-forced oven (Labec, general purpose 30 L) to obtain a powder. The powder was ground using an agate mortar before sintering at 1000°C for 2 h in a horizontal tube furnace (Labec, HTFS40-300/12). The specific surface area of the as-prepared sample was determined by gas adsorption with Brunauer, Emmett and Teller (BET) theory (Supplementary Table 11) and morphology by SEM (Supplementary Fig. 25).

Powder diffraction

X-ray diffraction (XRD) data were collected at the Centre for Microscopy and Microanalysis (CMM) at the University of Queensland (UQ) using a Bruker d8 advanced diffractometer. The silicon standard reference material (SRM) 640c from the National Institute of Standards & Technology (NIST) was used to determine the instrumental parameters for the experiment using Rietveld analysis with the GSAS II software⁶² and these were fixed in the subsequent structural refinement with a starting crystal structure of BaZr_0.5Fe_0.5O₃ ⁶³ and BaCO₃⁶⁴. Synchrotron X-ray powder diffraction (SXRD) data were collected at the Australian Synchrotron and neutron powder diffraction (NPD) data were collected using the high-resolution powder diffractometer (Echidna)⁶⁵ at the Australian Nuclear Science and Technology Organisation (ANSTO). Samples for powder diffraction analysis were heated to 600°C followed by air quenching. Samples were placed in glass capillaries with a diameter of 0.3 mm for SXRD data collection and in 6 mm diameter vanadium cans for NPD data collection and data collected under ambient conditions. For high temperature SXRD measurements, samples were loaded into 0.7 mm capillaries and heated by a Cyberstar hot gas blower with a ramping rate of 6°C min^− 1. Air and CO₂ were delivered to the samples using a XCS system.

The La¹¹B₆ SRM 660b from the National Institute of Standards & Technology (NIST) was used to determine the wavelength and instrumental parameters for the SXRD and NPD experiments using Rietveld analysis with the GSAS II software ⁶² (see Supplementary Figs. 27 and 28, respectively), and these were fixed in the subsequent structural refinement with a starting crystal structure of BaZr_0.5Fe_0.5O₃ ⁶³ and BaCO₃ ⁶⁴. In SXRD refinements, Fe and Co are considered identical because of their similar atomic number. In both SXRD and NPD refinements, Zr and Y are considered identical because of their similar atomic number and neutron scattering length ⁶⁶. In our refinement, the ICP-OES results were used to determine some atomic occupancies. A second perovskite with a different lattice parameter (~ 4.21 Å) was identified in our SXRD data, and included in the refinement using the NPD data, with the weight fraction fixed to that obtained in the SXRD refinement. Similar phase segregation was also reported in other Zr-containing perovskite oxides, such as BaCo_0.4Fe_0.4Zr_0.2O_3−δ²³ and SrCo_0.4Fe_0.6−xZr_xO_3−δ²⁴, which are attributed to the segregation between Zr and other B-site cations like Co and Fe ²⁵.

ICP-OES

Powder samples were first dissolved at 200°C with at 9:3:2 ratio of HNO₃, HF and HCl acids by microwave. 20 mL 5% boric acid was then added to neutralise the unreated HF acid. The solution was then analysed on a Thermo ICAP PRO XP ICP-OES instrument.

Phase analysis post CO ₂ reaction

To investigate the phases in sample reacted with CO₂, approximately 0.1 g of samples were treated in humid air containing 10% CO₂ for 1 h at 600°C in a horizontal tube furnace. After treatment, the sample was quenched to room temperature in air and investigated using laboratory based XRD, SXRD, and NPD, as outlined above.

Thermogravimetric analysis (TGA)

Material weight loss was investigated in air over the temperature range 200–700°C using a thermogravimetric analysis system (Perkin Elmer STA 6000). Samples were first calcined at 200°C in dry air for 1 h to remove absorbed moisture before heating to 700°C at a ramp rate of 2°C min^− 1.

Oxygen non-stoichiometry analysis using titration

A sodium thiosulfate solution (0.05 M) was first prepared by dissolving Na₂S₂O₃ (99.99%, Sigma Aldrich) in pure water (Milli-Q Type 1 Ultrapure Water System). A small amount of sodium bicarbonate was added to the solution to maintain a basic solution in which to stabilise Na₂S₂O₃. 0.05 g sample was dissolved in ~ 10 mL 2 M HCl (Sigma Aldrich) and deionized water was added to dilute to ~ 200 mL. Excess KI (99.0%, Sigma Aldrich) was added to the dilute solution to reduce transition metals. Approximately 2 mL of starch solution (Sigma Aldrich) was then added as the indicator and the solution was titrated with sodium thiosulfate solution.

Temperature-programmed desorption (TPD)

CO₂-TPD was performed using BETCAT and BETMASS systems. Approximately 0.1 g of powder was first treated in humid CO₂ at 600°C for 1 h followed by quenching to room temperature. The as-treated sample was placed in a U-shaped tube for TPD analysis. Samples were first heated to and held at 200°C for 1 h to remove moisture, before heating to 800°C at 2°C min^− 1 under Ar flow. The gas desorbed during this period was sampled and analysed by the bel-mass system.

Fourier-transform infrared (FTIR) and THz spectroscopy

Powders were first mixed with KBr (Sigma-Aldrich, 99.9%) for dilution. FTIR spectra were measured using a Spectrum 100, PerkinElmer, in the frequency range 800-2,000 cm^− 1 in attenuated total reflection mode. The resolution was 2 cm^− 1 and each sample was scanned for 4 times. THz spectra were measured in Australian Synchrotron THz/Far-IR beamline using a Bruker IFS 125/HR Fourier Transform Spectrometer. The detector is Si bolometer, which has a spectral region ranging from 10–650 cm^− 1. The resolution was 2 cm^− 1 and each sample was scanned for 3 times.

X-ray photoelectron spectroscopy (XPS)

XPS data were measured using a Kratos Axis Ultra spectrometric in the University of Queensland’s CMM with an Al Kα (1,486.8 eV) radiation source at 150 W. Fine scans were used to obtain O1s and C1s spectra and data were analysed using the CasaII software. The adventitious carbon was used for XPS calibration.

Microscopy

Samples for scanning electrom microscopy (SEM) analysis were pelletized and sintered at 1200°C for 5 h to obtain dense pellets. The dense pellets (~ 0.3 g) were treated in 10% CO₂ at 600°C for 1 h followed by quenching. The quenched samples were cracked and mounted for SEM. SEM data were obtained using a JEOL JSM-7001F at the University of Queensland’s CMM operating at 20 kV. Samples for transmission electron microscopy (TEM) analysis were ground and sized using a 100 µm sieve. The powder samples were treated in 10% CO₂ at 600°C for 1 h and were quenched to room temperature, The quenched samples were dispersed in ethanol by ultrasonic water baths. TEM images and EDS mapping were obtained using a Hitachi HT7700 at the University of Queensland’s CMM operating at 120 kV.

Differential scanning calorimetry (DSC)

DSC curves were obtained using a NETZSCH Photo-DSC 204 F1 Phoenix differential scanning calorimeter at the University of Queensland’s CMM. Phase pure samples and CO₂-treated samples were placed in Al cans and heated to 520°C, with the heat flow compared to an empty Al can.

Symmetric cell preparation

Approximately 5 g of powder cathode was mixed with 50 mL iso-propanol (99.5%, Sigma-Aldrich) and 3 mL of glycerol (99.0%, Sigma-Aldrich) and ball milled for 2 h in a planetary ball mill (Fritsch Pulverisette 5). The slurry was then coated onto an electrolyte surface by spray coating using a conventional spray gun. The thickness of the cathode layer was controlled to be ~ 10 µm by controlling the time of coating. After coating, the cell was sintered at 900°C for 2 h in a horizontal tube furnace (Labec, HTFS40-300/12) to obtain cells with a cathode|electrolyte|cathode configuration.

Electrochemical impedance spectroscopy (EIS)

The as-prepared symmetrical cells were heated in a horizontal tube furnace (Labec, HTFS40-300/120). Air was humidified via a wash bottle before being supplied to the cell. The airflow rate was approximately 200 mL min^− 1. An Autolab PGSTAT30 instrument was used to perform EIS tests in the frequency range 10⁵ to 10^− 1 Hz with a 10 mV amplitude. Post EIS test cells were also heated to 600°C and exposed to 10% CO₂ where further EIS was performed every 5 min for a 1 h period. After 1 h 10% CO₂ exposure, 200 mL min^− 1 fresh humid air was supplied to the cells for 1 h and EIS measurements again performed.

Single cell fabrication

Single cells were fabricated using a solid-state reaction method. Taking the anode supporting layer (NiO : BZCYYb : starch = 6 : 4 : 1) as an example, raw materials were mixed and ball-milled in ethanol to obtain a homogeneous powder. The powder was then pressed to form the anode support layer. The electrolyte (BZCYYb + 1% NiO) was obtained using a similar method and coated on the anode by co-pressing. The anode and electrolyte half cell were then sintered at 1350°C for 7 h to densify the electrolyte layer. The cathode was coated on the electrolyte surface by spray coating followed by sintering at 950°C for 2 h.

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In situ formation of a melt-solid interface towards stable oxygen reduction in protonic ceramic fuel cells

Status:

Version 1

Abstract

Figures

Introduction

Results

Crystal structure and chemical composition

Cathode ORR activity and CO₂ stability

The origin of CO₂ resistance in BCFZYLN

The role of co-doped alkali cations in suppressing carbonation

Surface segregation of alkali dopants in the presence of CO₂

Mixed alkali carbonate melt formation

Discussion

Conclusion

Declarations

Acknowledgements

Methods

References

Additional Declarations

Supplementary Files

Status:

Version 1

In situ formation of a melt-solid interface towards stable oxygen reduction in protonic ceramic fuel cells

Status:

Version 1

Abstract

Figures

Introduction

Results

Crystal structure and chemical composition

Cathode ORR activity and CO2 stability

The origin of CO2 resistance in BCFZYLN

The role of co-doped alkali cations in suppressing carbonation

Surface segregation of alkali dopants in the presence of CO2

Mixed alkali carbonate melt formation

Discussion

Conclusion

Declarations

Acknowledgements

Methods

References

Additional Declarations

Supplementary Files

Status:

Version 1

Cathode ORR activity and CO₂ stability

The origin of CO₂ resistance in BCFZYLN

Surface segregation of alkali dopants in the presence of CO₂