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Boosting the stability of perovskites with exsolved nanoparticles by B-site supplement mechanism - Nature Communications
Results .
DFT calculations and B-site supplement .
To reoccupy the B-site vacancies of reduced SFNM scaffold during exsolution, the appropriate filling agent should be selected to initiate the TIE process. For this purpose, the co-segregation energies of Ni, Fe, and Mo at B-site of SFNM were calculated by DFT simulation, which are ?2.26?eV, ?1.46?eV, and ?1.06?eV, respectively (Fig.? 1a ). It indicates that Fe exsolves more favorably than Mo but less favorably than Ni 30 . Considering that TIE is driven by the difference in co-segregation between guest ion and host ion 25 , the exchanges of guest Fe \({{\leftrightarrow }}\) host Ni (?0.79?eV) and guest Fe \({{\leftrightarrow }}\) host Fe (0?eV) are thermodynamically more favorable than that of guest Fe \({{\leftrightarrow }}\) host Mo (0.40?eV) when external Fe ion is introduced on the surface of SFNM substrate (Fig.? 1b ), suggesting the feasibility of refilling of guest Fe into the B-site vacancies and exsolution of Fe-Ni alloy.
Fig. 1: Feasibility of B-site supplement calculated by DFT calculations and schematic illustrations of two exsolution process. a Co-segregation energy and schematic illustrations of the DFT models for co-segregation by conventional exsolution. b Exchange energy comparison of B-site cations of SFNM with guest Fe. c Schematic illustration of exsolution with and without B-site supplement on SFNM.
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With the guidance of DFT calculations, two different exsolution routes are recapitulated in Fig.? 1c for schematic illustrations. Route i depicts the TIE-assisted exsolution initiated on the guest Fe-SFNM complex while the conventional counterpart is shown as route ii. In the TIE-assisted approach, the foreign Fe ions are initially deposited on the surface of perovskite by freeze drying (Supplementary Fig.? 1 ), which will then be incorporated into the remaining B-site vacancies, accompanied by the egression of host Ni and Fe upon reduction (Eq.? 4 ).
TIE-assisted exsolution:
$${{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{Fe}}}}}}}^{{{{{{\rm{X}}}}}}}+{{{{{{\rm{Ni}}}}}}}_{{{{{{\rm{Ni}}}}}}}^{{{{{{\rm{X}}}}}}}+2{{{{{{\rm{O}}}}}}}_{{{{{{\rm{O}}}}}}}^{{{{{{\rm{X}}}}}}}+2{{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{guest}}}}}}}\to {{{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{guest}}}}}}}}_{{{{{{\rm{Fe}}}}}}}^{{{{{{\rm{X}}}}}}}+{{{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{guest}}}}}}}}_{{{{{{\rm{Ni}}}}}}}^{{{{{{\rm{X}}}}}}}+2{{{{{{\rm{V}}}}}}}_{{{{{{\rm{O}}}}}}}^{{\cdot} {\cdot} }+4{{{{{{\rm{e}}}}}}}^{{\prime} }+{{{{{\rm{Fe}}}}}}-{{{{{\rm{Ni}}}}}}+{{{{{{\rm{O}}}}}}}_{2}$$
(4)
For the conventional exsolution, Fe- and Ni-site vacancies emerge in perovskite matrix with the formation of Fe-Ni alloy nanoparticles at the surface (Eq.? 5 ).
Conventional exsolution:
$${{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{Fe}}}}}}}^{{{{{{\rm{X}}}}}}}+{{{{{{\rm{Ni}}}}}}}_{{{{{{\rm{Ni}}}}}}}^{{{{{{\rm{X}}}}}}}+2{{{{{{\rm{O}}}}}}}_{{{{{{\rm{O}}}}}}}^{{{{{{\rm{X}}}}}}}\to {{{{{{\rm{V}}}}}}}_{{{{{{\rm{Fe}}}}}}}^{{\prime} {\prime} }+{{{{{{\rm{V}}}}}}}_{{{{{{\rm{Ni}}}}}}}^{{\prime} {\prime} }+2{{{{{{\rm{V}}}}}}}_{{{{{{\rm{O}}}}}}}^{\cdot \cdot }+{{{{{\rm{Fe}}}}}}-{{{{{\rm{Ni}}}}}}+{{{{{{\rm{O}}}}}}}_{2}$$
(5)
Examination of the exsolved nanoparticles and perovskite scaffold .
To determine the optimal Fe loading amount on surface of SFNM, a series of SFNM+ \(x\) Fe-red samples ( x ?=?0.0, 0.5, 0.8, 1.2, which refer to the molar ratios of guest Fe to host Ni) were prepared. X-ray diffraction (XRD) analysis confirms that the Fe-Ni alloy phases have emerged in all SFNM+ \(x\) Fe-red samples with the well-preserved perovskite phase (Fig.? 2a 1 ). Nevertheless, further increasing \(x\) to 1.5 gives rise to metallic Fe phase (PDF# 06-0696) (Supplementary Fig.? 2 ), suggesting that the excessive surface Fe seeds trigger the formation of Fe clusters at this deposition level. Therefore, the interplay between the nanoparticle formation and perovskite structure evolution among SFNM+ \(x\) Fe-red ( x ?=?0.0, 0.5, 0.8, 1.2) are further elucidated. The characteristic Fe-Ni alloy peaks in SFNM+0.0Fe-red can be well indexed to FeNi 3 phase (PDF#03-065-3244) (Supplementary Fig.? 3 ), while the peak located at 44.1° among SFNM+ \(x\) Fe-red ( x ?=?0.5, 0.8, 1.2) slightly shifts leftward (Fig.? 2a 2 ). Transmission electron microscopy (TEM) with energy-dispersive X-ray spectroscopy (EDS) element mappings on randomly selected nanoparticles of SFNM+0.0Fe-red and SFNM+1.2Fe-red reveal the significant increase of Fe proportion in the exsolved nanoparticles of SFNM+1.2Fe-red (Supplementary Fig.? 4 and Supplementary Table? 1 ). It may be presumably explained by the involvement of guest Fe in the nanoparticle growth, and the slightly smaller electronegativity of Fe than that of Ni causes the lattice expansion of Fe-Ni alloy 31 , 32 , 33 . More importantly, distinct peak-shifting of parent perovskites can be observed from the XRD patterns of SFNM+ \(x\) Fe-red. As shown in Fig.? 2a 3 , the magnified diffraction peak at about 32° of SFNM+0.5Fe-red shifts slightly to a higher angle with regards to SFNM+0.0Fe-red, which may be ascribed to the partial filling of B-site vacancies with the guest Fe. Oppositely, the peak shifts leftward when further increasing the deposited Fe content to x ?=?0.8 and 1.2, mainly due to the loss of lattice oxygen after the accelerated exsolution 34 . It in turn indicates the promotion of exsolution and substantial occupation of B-site vacancies by the guest Fe in SFNM+0.8Fe-red and SFNM+1.2Fe-red.
Fig. 2: Structure characterizations after reduction and reoxidation. a 1 XRD patterns of SFNM+ \(x\) Fe ( x ?=?0.0, 0.5, 0.8, 1.2) reduced at 800?°C for 2?h in 5% H 2 /N 2 from 20 to 80°. Amplified XRD patterns in the range of a 2 41–45° and a 3 31.5–33°. b 1 XRD patterns of re-oxidized SFNM+ \(x\) Fe-red at 1000?°C in air for 10?h. Amplified XRD patterns in the range of b 2 35–45° and b 3 26–31°.
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Field emission-scanning electron microscopy (FE-SEM) results demonstrate that the nanoparticle exsolution among SFNM+ \(x\) Fe-red ( x ?=?0.5, 0.8, 1.2) is more efficient compared with that on SFNM+0.0Fe-red, as illustrated by the larger average size, wider size distribution and larger population (Supplementary Figs.? 5a –d, 6 ). SFNM+ \(x\) Fe-red samples ( x ?=?0.0, 0.5, 0.8, 1.2) were subsequently subjected to O 1? s X-ray photoelectron spectroscopy (XPS) analysis to inspect the surface oxygen species. As shown in Supplementary Fig.? 7 , the concentration of surface-adsorbed oxygen gradually increases as the surface Fe deposition increases, which is consistent with the promoted exsolution observed by SEM.
To further verify the reoccupation of B-site by the guest Fe, SFNM+ \(x\) Fe-red samples ( x ?=?0.0, 0.5, 0.8, 1.2) were re-oxidized to inspect the ability of exsolved nanoparticles to reintegrate into the parent perovskite lattice 35 , 36 , 37 . After reoxidation at 1000?°C for 10?h, XRD analysis shows that the additional NiO peaks have clearly emerged in re-oxidized SFNM+0.0Fe-red, while NiFe 2 O 4 or SrMoO 4 can be detected on re-oxidized SFNM+ \(x\) Fe-red ( x ?=?0.5, 0.8, 1.2) (Figs.? 2b 1 –b 3 ). It can be concluded that the redissolution of the exsolved FeNi 3 nanoparticles is partially reversible on the SFNM+0.0Fe-red. Moreover, Fe atoms in FeNi 3 nanoparticles preferentially dissolve into perovskite compared to Ni atoms, which is in accord with the fact that Ni exsolves more favorably than Fe (Fig.? 1a ). As a result, the residual Ni on the surface transforms into NiO after heating in air 38 . However, for SFNM+ \(x\) Fe-red ( x ?=?0.5, 0.8, 1.2), the B-site vacancies of parent perovskite have been occupied by the guest Fe partially/entirely; the exsolved Fe and Ni on the surface are prone to self-assembly into binary oxide NiFe 2 O 4 in the air rather than dissolve into perovskite lattice. In addition, the secondary phase SrMoO 4 detected on re-oxidized SFNM+0.8Fe-red and SFNM+1.2Fe-red is due to the fact that B-sites were almost completely occupied by the high-valence Fe and Mo, exceeding the tolerance of perovskite structure. This is consistent with the emergence of SrMoO 4 in the air-sintered SFM 39 (Supplementary Fig.? 8 ). Therefore, it suggests that the target P-eNs have been successfully synthesized where Fe-Ni alloy nanoparticles are on the SFM substrate with decreased B-site vacancies and Ni incorporation.
Electrocatalytic performance .
The electrocatalytic capacities of SFNM+ \(x\) Fe-red for CO 2 reduction were subsequently examined by current density-voltage ( \(j\) –V) curves at 800 and 850?°C, as shown in Supplementary Fig.? 9 . The cathode electrodes were prepared by mixing SFNM+ \(x\) Fe-red with gadolinium doped ceria (GDC), and the full cells for the measurements of ((La 0.6 Sr 0.4 ) 0.95 Co 0.2 Fe 0.8 O 3- δ -GDCGDCYSZGDCSFNM+ \(x\) Fe-red-GDC) are denoted as SFNM+ \(x\) Fe-red-GDC for simplicity. Apparently, SFNM+1.2Fe-red-GDC exhibits superior electrocatalysis performance compared with others, indicative of its better catalytic activity towards CO 2 conversion. Specifically, the \(j\) reaches 0.97 and 1.13?A?cm ?2 for SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC, respectively, at 1.6?V and 850?°C (Fig.? 3a ).
Fig. 3: Electrocatalysis performances and surface activity characterization. a Current density-voltage curves and b EIS of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC for pure CO 2 electrolysis at 850?°C. c DRT analyses of the EIS for SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC at 1.6 and 1.8?V. d FTIR spectra of CO 2 chemisorption and physisorption for SFNM+0.0Fe-red and SFNM+1.2Fe-red at 600?°C. e Overlap of Sr- and O- EDS signals for SFNM+0.0Fe-red and SFNM+1.2Fe-red.
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To further gain insights into the CO 2 electrolysis in SOEC, SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC were subjected to the electrochemical impedance spectra (EIS) analysis. Figure? 3b shows the Nyquist plots of both cells at the applied potentials of 1.0, 1.2, 1.4, 1.6 and 1.8?V at 850?°C and the corresponding equivalent circuit model LR(Q H R H )(Q L R L ) (inserted image). The simulated ohmic resistance (R S ), polarization resistances at high and low frequencies (R H and R L ) are summarized in Supplementary Table? 2 . Both cells show the similar R S due to their identical cell assembly (Supplementary Fig.? 10 ), while the R P values of SFNM+1.2Fe-red-GDC are comparably smaller than that of the SFNM+0.0Fe-red-GDC at all the monitored potentials, suggesting the faster cathode kinetics for CO 2 reduction over SFNM+1.2Fe-red-GDC. Interestingly, R P of SFNM+0.0Fe-red-GDC firstly drops with the increase of voltage (from 1.0 to 1.6?V), following an increase at 1.8?V, while R P drops monotonously as the voltage increases over SFNM+1.2Fe-red-GDC. This confirms that the exceptional kinetic performance of SFNM+1.2Fe-red-GDC is well maintained even under harsh poling conditions. Remarkably, the R P at 1.8?V for SFNM+1.2Fe-red-GDC reaches an appreciably low value of 0.09?Ω?cm 2 , almost half the value of that for SFNM+0.0Fe-red-GDC (0.17?Ω?cm 2 ). In addition, the Nyquist plots and the fitted R p of SFNM+1.2Fe-red-GDC at 800?°C are obtained (Supplementary Fig.? 11 and Supplementary Table? 3 ). It shows the lowest R p at high voltages (≥1.6?V) among the state-of-art P-eNs-based cathode candidates as listed in Supplementary Table? 4 , indicative of the higher catalytic activities of SFNM+1.2Fe-red-GDC at higher negative potentials.
In conjunction with the EIS data, distribution of relaxation times analysis was used to discern the contributions of underlying kinetic processes to the polarization resistances over these two cells 40 . As shown in Supplementary Fig.? 12 , the electrode process can be deconvoluted into four peaks, which are denoted as (1) oxygen ion transfer through electrolyte and oxygen evolution at anode (P 1 ), (2) charge transfer (P 2 ), (3) surface CO 2 adsorption and activation (P 3 ), (4) gas diffusion process (P 4 ) at cathode from high to low frequency 37 , 41 . Notably, the integral areas of P 2 and P 3 peaks on SFNM+0.0Fe-red increase as voltage rises from 1.6 to 1.8?V, as opposite to that on the SFNM+1.2Fe-red (Fig.? 3c ). It can be inferred that the higher surface reactivity, the faster charge transfer on SFNM+1.2Fe-red, whereas the passivated surface would deteriorate the charge transfer on SFNM+0.0Fe-red at 1.8?V.
The origin of the enhanced catalytic activity .
To understand the origin of the difference in catalytic performances between these two cells for CO 2 reduction at high voltages, further characterizations were carried out on SFNM+0.0Fe-red and SFNM+1.2Fe-red. First, Fourier transform infrared spectroscopy (FTIR) measurement was performed to compare their CO 2 adsorption capacities 14 , 42 . As shown in Fig.? 3d and Supplementary Fig.? 13e–f , SFNM+1.2Fe-red exhibits a higher chemical CO 2 adsorption signal with a quicker response time at 600?°C than the SFNM+0.0Fe-red, implying the superiority of SFNM+1.2Fe-red in terms of surface CO 2 adsorption and reactivity. However, SFNM+0.0Fe-red presents more intensive physical adsorption peaks compared with SFNM+1.2Fe-red at all the examined temperatures (Fig.? 3d and Supplementary Fig.? 13b ). It may be ascribed to its higher surface alkalinity caused by more severe Sr-segregation. As shown in Eq.? 6 , the Fe-Ni nanoparticle exsolution, together with the exhaustion of B-site cations and lattice oxygen, in a conventional way would trigger the stripping out of Sr cations to re-establish stoichiometry.
Conventional exsolution:
$${{{{{{\rm{Sr}}}}}}}_{2}{{{{{{\rm{Fe}}}}}}}_{1.3}{{{{{{\rm{Ni}}}}}}}_{0.2}{{{{{{\rm{Mo}}}}}}}_{0.5}{{{{{{\rm{O}}}}}}}_{6-\delta }\mathop{\longrightarrow}^{{{{{\mathrm{Reduction}}}}}}{{{{{{\rm{Sr}}}}}}}_{2-x-y}{{{{{{\rm{Fe}}}}}}}_{1.3-x}{{{{{{\rm{Ni}}}}}}}_{0.2-y}{{{{{{\rm{Mo}}}}}}}_{0.5}{{{{{{\rm{O}}}}}}}_{6-{\delta }^{{\prime} }}+\left(x+y\right){{{{{\rm{SrO}}}}}}+x{{{{{\rm{Fe}}}}}}-y{{{{{\rm{Ni}}}}}}$$
(6)
The Sr-rich surface would reorganize into SrCO 3 at a high CO 2 concentration, which has a detrimental effect on the surface catalytic activity for CO 2 reduction 43 . Furthermore, high voltages would impose high reducing potentials at the cathode, the concomitant exsolution and Sr segregation would be sustained by the voltage driving force. On the contrary, the undesired Sr segregation could be efficiently alleviated when the B-site vacancies are occupied by the guest Fe; the active sites can be well exposed even at high potentials (Eq.? 7 ).
TIE-assisted exsolution:
$${{{{{{\rm{Sr}}}}}}}_{2}{{{{{{\rm{Fe}}}}}}}_{1.3}{{{{{{\rm{Ni}}}}}}}_{0.2}{{{{{{\rm{Mo}}}}}}}_{0.5}{{{{{{\rm{O}}}}}}}_{6-\delta }+(x+y){{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{guest}}}}}}}\mathop{\longrightarrow}^{{{{{\mathrm{Reduction}}}}}} {{{{{{\rm{Sr}}}}}}}_{2}{{{{{{\rm{Fe}}}}}}}_{1.3-x}{{{{{{\rm{Ni}}}}}}}_{0.2-y}{{{{{{{\rm{Fe}}}}}}}_{{{{{{\rm{guest}}}}}}}}_{x+y} {{{{{{\rm{Mo}}}}}}}_{0.5}{{{{{{\rm{O}}}}}}}_{6-{\delta }^{{\prime} }} \\ + x{{{{{\rm{Fe}}}}}}-y{{{{{\rm{Ni}}}}}}$$
(7)
The difference in Sr segregation between SFNM+0.0Fe-red and SFNM+1.2Fe-red has also been verified by comparable EDS signal intensity of Sr in the form of thin shell around the exsolved nanoparticles 19 (Fig.? 3e ). Combined with the facilitated formation of Fe-Ni alloy nanoparticles and surface-active oxygen vacancies, it in turn persuasively confirms the evolution of initial surface of SFNM+1.2Fe-red into a more catalytically active surface by B-site filling, which contributes to enhanced CO 2 adsorption and activation at high operating voltages 14 .
Thermogravimetric analysis reveals that the lattice oxygen loss increases with the guest Fe deposition amount among SFNM+ \(x\) Fe-red ( x ?=?0.0, 0.5, 0.8, 1.2), indicative of higher oxygen vacancy concentration within perovskite scaffold of SFNM+1.2Fe-red than that of SFNM+0.0Fe-red (Supplementary Fig.? 14 ). Since the oxygen ion transfer in perovskite is realized by reverse jumping of oxygen vacancies, the oxygen ion conductivity is also dependent on the mobility of oxygen vacancies in addition to oxygen vacancy concentration 20 . Apparently, the association between neighboring B-site vacancy and oxygen vacancy on SFNM+0.0Fe-red is likely to trap the oxygen vacancy. The calculated binding energy of the \({V}_{{Fe}}^{{\prime} {\prime} }-{V}_{O}^{\cdot \cdot }\) defect pair is ?1.73?eV (Supplementary Fig.? 15 and Supplementary Table? 5 ), indicating that the additional association barrier needs to be overcome to achieve the jumping of oxygen vacancies 20 . In contrast, the full occupation of B-sites on SFNM+1.2Fe-red allows the elimination of constraint of B-site defect to surrounding oxygen vacancies. The electrical conductivity relaxation (ECR) experiments for SFNM+0.0Fe-red and SFNM+1.2Fe-red were carried out to study the oxygen ion bulk diffusion 44 , 45 (Fig.? 4a ). The bulk diffusion constant ( \({D}_{{chem}}\) ) of SFNM+1.2Fe-red obtained by fitting the ECR curve is \(3.747\times {10}^{-4}\) ?cm 2 s ?1 , much higher than \(1.439\times {10}^{-5}\) ?cm 2 ?s ?1 of the SFNM+0.0Fe-red 46 . It suggests that there are more available oxygen ion transfer pathways in SFNM+1.2Fe-red, which lead to the higher oxygen ion conductivity and lower charge transport resistance.
Fig. 4: Charge transport characterizations and X-ray photoelectron spectroscopy for SFNM+0.0Fe-red and SFNM+1.2Fe-red. a Normalized electrical conductivity relaxation curves for SFNM+0.0Fe-red and SFNM+1.2Fe-red at 850?°C with the switching of the gas stream from 2:1 CO?CO 2 to 1:1 CO?CO 2 . XPS of b Fe 2 p and c Mo 3 d for SFNM+0.0Fe-red and SFNM+1.2Fe-red. d Temperature-dependent electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red under 50% CO 2 /50% CO atmosphere.
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Since the valence states of B-site cations have a great influence on the electronic transfer within perovskite bulk, the electronic environments of B-site cations on SFNM+0.0Fe-red and SFNM+1.2Fe-red were examined by the XPS, X-ray absorption near edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) spectra 47 , 48 , 49 (Fig.? 4b–c and Supplementary Figs.? 16 – 17 ). As shown in Fig.? 4b , higher proportion of Fe 3+ (23.30%) in SFNM+1.2Fe-red can be observed than that in SFNM+0.0Fe-red (16.29%). In parallel, there is a distinct reduction in the average oxidation state of the Mo atom in SFNM+1.2Fe-red (+5.67) with respect to that in SFNM+0.0Fe-red (+5.89), which was derived from partial reduction of Mo 6+ to Mo 5+ (Fig.? 4c ), indicating the coexistence of Fe 3+ -Mo 5+ and Fe 2+ -Mo 6+ electronic configurations to achieve charge neutrality (Fe 3+ +Mo 5+ ?=?Fe 2+ +Mo 6+ ) 50 . More Fe 2+ -Fe 3+ and Mo 5+ -Mo 6+ charge pairs should endow SFNM+1.2Fe-red with the improved electronic conductivity owing to the sufficient B (n?1)+ -O-B n+ conduction pathways in SFNM+1.2Fe-red. The temperature-dependent electrical conductivity results show that SFNM+1.2Fe-red performs better than SFNM+0.0Fe-red in the 50% CO 2 /50% CO and 5% H 2 /95% Ar atmospheres (Fig.? 4d and Supplementary Fig.? 18 ). In the 50% CO 2 /50% CO atmosphere, the electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red increase as the temperature increases, reaching 16.5 and 24.6?S?cm ?1 at 850?°C, respectively (Fig.? 4d ). SFNM+1.2Fe-red shows the higher conductivity in the prospective operation atmosphere of CO 2 electrolysis.
In summary, the B-site supplement of the reduced SFNM by external Fe source plays a crucial role in preserving the high surface activity, high oxygen ion conductivity and electronic conductivity at high voltages, thus leading to the faster cathode reaction kinetics.
Short-term/long-term stability performances .
The cathode stability under operational conditions is a vital criterion to evaluate the CO 2 electrocatalysis performances in SOEC. The 15?min potential step chronoamperometry was firstly performed at 850?°C to gain initial indication of the stability of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC. Online gas chromatography (GC) was employed to monitor the CO formation during the short-term stability tests. These two cells deliver quite similar \(j\) at negative potentials of 1.0, 1.2, and 1.4?V (Fig.? 5a ). Upon increasing the potential to 1.6?V and then to 1.8?V stepwise, stark differences in \(j\) have emerged. The significant decay of \(j\) can be observed over SFNM+0.0Fe-red-GDC, whereas there is negligible \(j\) decline over SFNM+1.2Fe-red-GDC under both polarization conditions. SFNM+1.2Fe-red-GDC shows the competitive short-term stability at the voltages ≥1.6?V in comparison with the SFNM+0.0Fe-red-GDC and other state-of-art P-eNs-based SOECs (Supplementary Table? 6 ). Additionally, both cells could steadily generate CO with appreciable Faraday efficiency ( \({{{{{{\rm{FE}}}}}}}_{{{{{{\rm{CO}}}}}}}\) ) (~90%) at all the potentials (Fig.? 5b ). As expected, the CO production rate for SFNM+1.2Fe-red-GDC gradually increased with increasing the external voltages, while the production rate for SFNM+0.0Fe-red-GDC peaked at 1.6?V and then decreased upon further raising the potential to 1.8?V. Considering that the nanoparticles closely socketed on the surface of SFNM+0.0Fe-red and SFNM+1.2Fe-red can retain the virtue of high resistance to agglomeration 51 (Supplementary Fig.? 19 ), the discrepancies in \(j\) decay and CO productivity appear to be raising from the differences in the structure evolution of perovskite matrix at higher negative potentials.
Fig. 5: Stability performances and post-mortem surface microstructural characterizations. a The current density response curves for the 15?min potential step chronoamperometry and b corresponded CO productivity and FE CO for SFNM+0.0Fe-red-GDC and SFNM+ 1.2Fe-red-GDC at 850?°C with pure CO 2 (Gray shaded areas indicate the applied voltages). c 100?h long-term stability testing for SFNM+0.0Fe-red-GDC and SFNM+ 1.2Fe-red-GDC at 1.6?V and 850?°C in pure CO 2 . SEM images of cathode surface microstructure of d 1 – d 2 SFNM+0.0Fe-red-GDC and e 1 – e 2 SFNM+1.2Fe-red-GDC after long-term stability at 1.6?V and 850?°C. SE-STEM and STEM-EDS images of cathode on f SFNM+0.0Fe-red-GDC, g SFNM+1.2Fe-red-GDC after long-term stability at 1.6?V and 850?°C.
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It should be worth noting that the stable SOEC operation at high voltages has significant implications on improving the energy efficiency and industrial scale of applications 18 . Furthermore, high cathode potentials are applied to intentionally accelerate electrode degradation to help us understand the cathode evolution and further shed lights on the degradation mechanisms. This in turn is of great significance for designing P-eNs materials in line with the goal for industrialization. Figure? 5c presents the prolonged stability performances of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC at 850?°C and 1.6?V. Interestingly, SFNM+0.0Fe-red-GDC experiences the current density output instability at around 55?h. Moreover, despite the apparent \(j\) decrease over the entire operation, a trend of “degradation?reactivation?degradation” emerges that can be roughly categorized to Phase I (0–8?h), Phase II (8–35?h), and Phase III (35–55?h). Within Phase I, \(j\) experiences dramatic decrease with a degradation rate of 44?mA?cm ?2 ?h ?1 . As the reaction proceeds and enters Phase II, the \(j\) slowly climbs up, peaking at 35?h with a maximum \(j\) of 0.47?A?cm ?2 . In the following Phase III, the \(j\) drops again until it becomes unstable at around 55?h. Especially, \(j\) declines at an attenuation rate of 6?mA?cm ?2 ?h ?1 from 50 to 55?h. In sharp contrast to the complicated \(j-t\) profile on SFNM+0.0Fe-red-GDC, SFNM+1.2Fe-red-GDC manifests a rather simple and mild degradation profile throughout the long-term stability measurement without observable \(j\) fluctuations, only having an average degradation rate of 3?mA?cm ?2 ?h ?1 . The \(j\) remains at a high value of 0.75?A?cm ?2 at 100?h. In addition, the recorded \({{{{{{\rm{FE}}}}}}}_{{{{{{\rm{CO}}}}}}}\) profiles show that SFNM+0.0Fe-red-GDC experienced a significant \({{{{{{\rm{FE}}}}}}}_{{{{{{\rm{CO}}}}}}}\) decline, only having 75% at 48?h. In contrast, the \({{{{{{\rm{FE}}}}}}}_{{{{{{\rm{CO}}}}}}}\) of SFNM+1.2Fe-red-GDC has remained at above 90% throughout the 100-h stability test, indicating its superiority in preventing the yielding of by-products.
Post-mortem characterization and degradation mechanism .
After the long-term stability tests, the significant morphology changes and coarsening of exsolved particles can be observed on the both cathode surfaces, and the coarsening of surface nanoparticles on SFNM+1.2Fe-red-GDC is less severe than that on SFNM+0.0Fe-red-GDC (Fig.? 5d–e ). Both cathodes were scratched from the cells for the element distribution detection by secondary electron-scanning TEM (SE-STEM) and STEM-EDS. More pronounced Sr- and Mo-derived phase separation in the form of irregular particles appear on the surfaces of SFNM+0.0Fe-red-GDC, (Fig.? 5f–g ), suggesting that SFNM+0.0Fe-red-GDC undergoes the more significant phase decomposition than SFNM+1.2Fe-red-GDC under the synergistic effect of the high voltage and CO 2 /CO mixed atmosphere during the CO 2 electrolysis. The Raman spectra from cathode surface of both cells were collected (Fig.? 6a ). As shown, the distinct Raman feature peaks of SrCO 3 (located at 701?cm ?1 ), SrMoO 4 (located at 795, 842, and 887?cm ?1 ), carbon with lattice defect ( D band, located at 1319?cm ?1 ) can be found on the cathode surface of SFNM+0.0Fe-red-GDC 39 , 52 , 53 , whereas only trace amounts of SrMoO 4 peaks were detected on the cathode surface of SFNM+1.2Fe-red-GDC. The emergence of Sr-containing impurities on both spectra confirms the structure decomposition and surface reconstruction on SFNM+0.0Fe-red and SFNM+1.2Fe-red. Obviously, the phase decomposition of SFNM+0.0Fe-red is more thorough. Furthermore, a small amount of carbon deposition on SFNM+0.0Fe-red-GDC may be ascribed to the newly born Ni nanoparticles driven by the applied voltage. Subsequently, SFNM+0.0Fe-red-GDC was subjected to the stability test again and was interrupted before entering the Phase II, and the Raman spectrum of the cathode surface was collected. As shown in Fig.? 6b , only the Raman feature peaks of SrCO 3 and SrMoO 4 were detected but no peaks of carbon, which confirms that the perovskite scaffold of SFNM+0.0Fe-red primarily underwent the bulk structural decomposition before the exposure of newly grown nanoparticles and carbon deposition. This also explains the \(j\) evolution over SFNM+0.0Fe-red-GDC.
Fig. 6: Post-mortem Raman spectra characterizations and illustrations of degradation mechanisms. a Raman spectra collected from cathode surface of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC after the long-term stability test. b Raman spectra collected from cathode surface of SFNM+0.0Fe-red-GDC whose stability test was interrupted before entering Phase II. c Structure evolution-driven degradation mechanisms for SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC at high negative potentials.
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To verify the exsolution scenario and the structural decomposition during the electrolysis process, the SFNM without pre-reduction treatment was fabricated as the composite cathode material (SFNM-oxi-GDC) for the electrochemical testing (Supplementary Fig.? 20 ) and the post-mortem characterizations (Supplementary Figs.? 21 – 23 ). As shown in Supplementary Fig.? 20d , the SFNM-oxi-GDC shows the similar long-term stability profile as the SFNM+0.0Fe-red-GDC. The SEM, SE-STEM, and STEM-EDS results confirm that partial Ni elements segregate from the SFNM bulk and the newly born fibrous phases appear on the cathode surface (Supplementary Figs.? 21 – 22 ). Meanwhile, the Raman feature peaks of SrCO 3 , carbon with lattice defect (D band) can be observed (Supplementary Fig.? 23 ), suggesting that the Ni element in situ exsolves from the SFNM bulk and subsequently results in the carbon fiber growth during CO 2 electrolysis. This can be ascribed to the lower co-segregation energy of Ni in SFNM, which causes the structural decomposition under the synergistic effect of 1.6?V and 850?°C.
Accordingly, the severe degradation may be effectively mitigated by preserving the structural integrity and enhancing the resistance to exsolution of perovskite scaffold of SFNM+0.0Fe-red. In SFNM+1.2Fe-red, the robustness of perovskite scaffold has been significantly enhanced by B-site supplement using redox-stable Fe ions, which postpones the decomposition of perovskite bulk and reassembly of Sr-based insulators on surface. Furthermore, the continuous exsolution has been greatly suppressed due to the decreased Ni content in perovskite bulk. This also well accounts for the mild degradation rate, high \({{{{{{\rm{FE}}}}}}}_{{{{{{\rm{CO}}}}}}}\) , no intuitive reactivation process, decreased Sr-based matters and negligible deposited carbon over SFNM+1.2Fe-red-GDC.
In view of the above analyses, a perspective on the degradation mechanism of P-eNs from the structure stability of perovskite substrate can be proposed. Since the proper negative potential acts as the driving force to drain out B-site reducible cations from perovskite, the high voltage on cathode will not only accelerate CO 2 reduction reaction, but also give rise to the continuous and slow exsolution during the electrolysis 23 , starting from breaking of weak B–O bonds, diffusion of B-site cations, followed by nucleation and growth of nanoparticles 54 . In parallel, A-site segregation occurs, leading to surface reassembly and perovskite bulk reconstruction. As below, we can tentatively discuss the differences in structure evolutions of SFNM+0Fe-red-GDC and SFNM+1.2Fe-red-GDC in light of their degradation processes (Fig.? 6c ). For SFNM+0Fe-red-GDC, the profile presents a high \(j\) but a faster decay rate in the Phase I (Fig.? 5c ). This is presumably ascribed to the breaking of B–O bonds and subsequent diffusion of ions from bulk to surface driven by the external voltage 55 , which causes deteriorated charge transfer capability of perovskite scaffold and surface assembly of detrimental Sr-based matters (Fig.? 6b ). This stage is named as “Structure decomposition” regime. However, such a declining profile has been gradually counterbalanced by the newly exposed active nanoparticles derived from diffusion of active cations in the deeper region of perovskite (Phase II in Fig.? 5c ). It can be seen that this process is very mild, which means that exsolution is limited by the concentration of exsolved cations in the perovskite and proceeds slowly 54 . This stage is termed as the “Reactivation” regime. As electrolysis proceeds, the structural decomposition progresses slowly, and a balance between surface reorganization and bulk reconstruction has gradually reached. Meanwhile, carbon deposition around the nanoparticles becomes increasingly distinct (Fig.? 6a ). Consequently, \(j\) drops again or even becomes unstable (Phase III in Fig.? 5c ), this period can be named as the “Passivation” zone. Nevertheless, for SFNM+1.2Fe-red-GDC, the “structure decomposition” has been significantly alleviated by the reoccupation of relatively redox-stable Fe ions at B-site vacancies of perovskite. The high applied potential offers a higher initial transient \(j\) , however, the amplitude of \(j\) attenuation has been remarkably reduced at the initial stage compared with that for SFNM+0.0Fe-red-GDC, which proves that the structural decomposition has been greatly alleviated (Fig.? 5c ). Although the “reactivation” stage is not visible from the \(j\) curve, the change in \({{FE}}_{{CO}}\) and phase analysis from Raman results suggest the existence of a small amount of exsolution and surface reconstruction (Figs.? 5 c and 6a ). Moreover, carbon deposition near active sites has been circumvented (Fig.? 6a ), resulting in a moderate degradation process in the following reaction and a longer operating life. In addition, the deteriorations of the electrolyte and the anode under the harsh conditions seem to be inevitable (Supplementary Fig.? 24 ), which are also responsible for the significant attenuation of the current density during CO 2 electrocatalysis at 1.6?V and 850?°C. .
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