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C: Surfaces, Interfaces, Porous Materials, and Catalysis

Conduction Mechanisms in Oxide-Carbonate Electrolytes for SOFC: Highlighting the Role of the Interface From First-Principles Modelling Chiara Ricca, Armelle Ringuedé, Michel Cassir, Carlo Adamo, and Frédéric Labat J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b02174 • Publication Date (Web): 16 Apr 2018 Downloaded from http://pubs.acs.org on April 16, 2018

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The Journal of Physical Chemistry

Conduction Mechanisms in Oxide-Carbonate Electrolytes for SOFC: Highlighting the Role of the Interface From First-principles Modelling Chiara Ricca1,2, Armelle Ringuedé1, Michel Cassir1, Carlo Adamo1,3, and Frédéric Labat1* 1

PSL Research University, Chimie Paristech-CNRS, Institut de Recherche de Chimie de Paris, 11 rue Pierre et Marie Curie, F-75231 Paris Cedex 05, France 2 Sorbonne Universités, UPMC Univ Paris 06, IFD, 4 place Jussieu, 75252 PARIS cedex 05, France 3 Institut Universitaire de France, 103 Bd Saint-Michel, F-75005, Paris, France Abstract A comprehensive density functional theory (DFT) investigation of conduction mechanisms of both intrinsic and extrinsic species found in oxide-carbonate composite electrolytes materials used in SOFC is presented. An interface model of YSZ-LiKCO3 is used as a case study to investigate transport properties, considering different mechanistic assumptions suggested in the literature over the years. Results clearly indicate that interfaces play an extremely important role in influencing the ionic conductivity performances of these electrolytes. In particular, redistribution of ions of both phases can lead to the formation of the so-called space charge layer at the interface, which then favourably influences the transport of intrinsic cations or extrinsic protons in particular. These results constitute an important first step towards the understanding of the electrochemical processes and transport mechanisms in these electrolytes, in order to elucidate the origin of their enhanced conductivity at low temperature when compared to that of their parent components. *Correspondence to: [email protected]

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1. Introduction Oxide-carbonate composite electrolytes are promising materials for SOFC applications thanks to their enhanced conductivity at lower temperatures (typically between 0.01 and 0.1 S/cm at 500-600°C1–6) when compared to their corresponding oxide parent pure phases (800-1000°C) conventionally used as electrolytes in these devices. Such composites are made of eutectic mixtures of alkaline carbonates (such as (Li,Na)2CO3, or (Li,K)2CO3) and an oxide phase, generally cubic fluorite oxides such as yttria-stabilized zirconia (YSZ) or ceria-based materials.2,7,8 The two phases are distributed according to a core-shell structure, with a layer of the mainly amorphous carbonate phase surrounding the crystalline oxide grains.9 So far, the origin of their enhanced performances is still not fully understood, although interfaces between the oxide and the carbonate phases are expected to play a key role.7,8,10–14 Several explanations and related conduction mechanisms have been proposed by different groups throughout the years8,10,13–16. Different species are supposed to contribute to the transport in the material: the oxygen ions, the salt cations, the carbonate ions, the protons or maybe several of them. Although the role of oxide and carbonate ions in the electrolyte performances is clear17, it has also been supposed that a space-charge layer can be formed at the interface of the composite due to the interfacial interaction: oxygen ions could accumulate at the interface, where there could also be an enrichment of the cations of the carbonate.16,18,19 The higher concentration of the defects/ions near the phase boundaries compared to the bulk phase may constitute ’superionic highways’ at the boundary between the two phases. Moreover, great emphasis was given on the supposed multi-ionic nature of these materials, which were claimed to be dual or hybrid H+/O2− or even ternary H+/O2−/CO2− conductors when operating in O2/CO2 atmosphere at the cathode side. Although the multi-ionic conduction can help in rationalizing the conductivity enhancement,

transport

mechanisms

at

play

still

remain

largely

unconfirmed.7,10,11,13,16,17,20 This short discussion underlines that there is still a huge debate concerning both the nature of the charge carrier(s) and the related conduction mechanisms. This debate is also generated by the fact that few experimental analytical tools allows for a direct study of the interface phenomena.21 For this reason, theoretical investigations can definitely help providing a picture of the properties of the composites at the atomic scale by directly describing the effect of the interfaces on the conduction properties. ACS Paragon Plus Environment

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Dealing with these two phase systems and providing an accurate and reliable description of their properties can however be challenging, even from the modelling viewpoint, since it requires a careful choice of the methodology and of the models used. In order to shed some light on these points we have carried out a detailed theoretical investigation, using first-principle methods based on Density Functional Theory (DFT) and a previously-proposed YSZ-LiKCO3 interface model29, which has been developed and validated by a combined theoretical and experimental protocol2429

. In particular, our DFT simulations of the transport mechanisms of Li, O, and

proton-related defects provide balanced descriptions which, albeit issued from a modeling study, allow for the first time for a direct comparison of all major phenomena involved. Furthermore, our study provides valuable hints to the experimentalists on the physics and chemistry underpinning the enhancement behavior observed and providing some guidelines to further improve it.

2. Computational Details All DFT calculations were performed using the periodic CRYSTAL14 code.24 Following our previous investigation on the YSZ-LiKCO3 interface25, the hybrid PBE026,27 exchange-correlation functional was selected, together with Effective Core Potentials (ECP) and associated basis sets. In particular, the relativistic Hay and Wadt28–30 small core ECP were used for Zr, Y, and K, while the Durand-Barthelat31,32 large core ECP was applied for O in ZrO2. An ECP was also used for K atoms of the carbonate together with a valence double-ζ (DZ) basis set.33 Valence electrons were described by a polarized quintuple- ζ (5Z), triple- ζ (TZ), and double- ζ (DZ) basis sets for Zr, Y, and O, respectively. All-electrons basis sets of DZ (for H, Li and C) and TZ (for O) quality were used for all the remaining atoms33. The 2D periodic interface model used to investigate the conduction mechanisms in these composite materials is taken from Ref.

25

: the unit cell of the YSZ-LiKCO3

interface has P1 symmetry with cell parameters a=7.17 Å, b=6.27 Å and γ=90.10° and contains 106 atoms, 36 (Li6K6(CO3)6) belonging to the 3 layers of the carbonate phase and 72 (Zr20Y4O46) organized in 6 O-Zr/Y-O layers along the c-axis for YSZ. It was built starting from models of the oxide and carbonate most stable surfaces that were previously found to ensure convergence of the structure. The detailed strategy


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followed to build the interface is fully detailed in mentioned references and in the supporting material.

Li defects, proton and O-related species were introduced in the optimized unit cell describing the oxide-carbonate interface, which was found to be large enough to avoid spurious interactions between the defects and their periodic images. Neutral O defects were created by adding one O atom in the unit cell, while the Li and O Frenkel pairs (LiFP and OFP) were obtained by displacing a lattice Li+ or O2- ion into an interstitial position (Lii• or Oi’’) and leaving behind a negatively charged Li vacancy (V’Li) or a doubly positive charged O vacancy (VO••). To study oxide ion or proton transport, one O2− anion together with two interstitial Li cations (Lii•), in the first case, or two H+ together with one O2- ion, in the second one, were introduced in the system, thus ensuring the neutrality of the unit cell. Several different initial positions of the defects were taken into account in all cases. To evaluate the relative thermodynamical stability of the different defects, formation energies (Ef (D)) have been computed as: = − + ∑

(1)

where Et(D) is the total energy of the defective system, Et(X) is the total energy of the system without defects, ni is the number of atoms of element i removed to create the defect (ni is negative when atoms are added instead), µi is the chemical potential of the element i removed (or added). For the oxygen chemical potential, we used the oxygen molecule in its triplet state as a reference. Finally, the transport properties of different possible species (cations of the carbonate eutectic, O2- anions, and H+) were investigated simulating several possible conduction mechanisms proposed over the years. Diffusion barriers for Li-related, O-related and H+ species were qualitatively estimated by performing relaxed scan calculations with the selected species at different contiguous positions along a selected conduction path, as already done for our previous study of Li transport on the clean LiKCO3 surface.34

3. Results and Discussion Before discussing the conduction mechanisms in oxide-carbonate composite materials, we briefly recall that the theoretical investigation of YSZ-LiKCO3 interface


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we previously performed showed that this material is characterized by a strong and favourable interaction taking place between the two phases, as indicated by the large value of the computed adhesion energy (3.19 eV25). This interaction was found to have two main origins: (i) adsorption of the ionic species of the salt on the oxide surface and (ii) minimization of the elastic strain due to the lattice mismatch between the two surfaces, resulting in a disordered amorphous-like state for the carbonate phase in agreement with the experimental observations. The YSZ-LiKCO3 interface model obtained through simulations was validated by comparison of the theoretical data with IR and Raman measurements25. In the following, we discuss results relative to the formation and the transport properties of both intrinsic and extrinsic species, starting first with Li-related defects, before moving to O-related ones and finally to the proton case. Main transport mechanisms considered are sketched in Schemes 1 and 2 for the intrinsic and extrinsic species respectively, and a summary of main results obtained is presented in Table 1. Additional details regarding the investigated conduction mechanisms are available in the Supplementary Information. Since the conductivity enhancement in carbonate-oxide composites is already observed in air atmospheres at certain temperatures lower than the melting point of the carbonate eutectic mixture,17 it is in principle necessary to take into account all intrinsic species when considering transport properties in the electrolyte, that is to say the CO32- groups, the cations M+ of the carbonate phase, and the O2- anions of the oxide. However, CO32- species can be excluded as contributing to the transport, since: (i) under experimental conditions, the carbonate phase is still in the solid state before the melting temperature with covalently bonded carbonate groups, (ii) the creation of anion vacancies through the removal of the carbonate groups or their occupation of interstitial sites is energetically highly demanding, being much larger than M+ cations. In addition, K+ ions are larger than Li+ ones and should then move more slowly, contributing much less to the transport properties in the electrolyte than Li+.35,36 They have therefore not been considered any further as participating to the transport in the following. 3.1. Intrinsic species: the case of Li In the “space charge layer” theory37–42 developed by

Maier and applied to

composites, a relevant role is played by cations present in the carbonate phase.16,18,19


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In particular, they can be stabilized by adsorption on the oxide surface, thus increasing cation vacancies in the bulk carbonate phase near the boundaries, which can be responsible for the observed conductivity enhancement in air atmosphere at low temperatures (see Scheme 1.A1). To investigate the influence of the interfaces on Li transport in the composites, neutral Li Frenkel Pairs (LiFP) where a Li+ is displaced from its lattice position to a nearby interstitial site (becoming Lii•) at the YSZ-LiKCO3 interface, thus creating a vacancy (V’Li) at its original site, were first investigated. The vacancy, V’Li, can thus contribute to the formation of the ’space charge layer’ in the bulk carbonate phase near the boundaries. Three possible configurations were generated by displacing Li+ cations found at different distances along the normal to the interface, so as to create Lii•-V’Li pairs with increasing separation distance along the c-axis (about 2.7 Å for Li(1)FP, 4.3 Å for Li(2)FP , and 6.7 Å for Li(3)FP). In all cases, Li+ was displaced from its original site to the same interstitial position at the interface, where it was allowed to relax. Table S1 shows the atomic configuration, the most important structural parameters and the formation energies (Ef) obtained for the three cases taken into account. As we can see from these data, the Ef values range between 1.07 and 1.26 eV, and are all lower than the previously computed formation energy for a LiFP created in the clean LiKCO3-(001) surface (1.79 eV) at the same level of theory.34 This result suggests that the formation of the interface can indeed modify the defect chemistry of the carbonate phase, due to the stabilizing interaction taking place in the boundary region between the two phases. Maier’s model is based on the assumption that the Lii• species are stabilized and segregated at the interfaces, while the conductivity enhancement is due to vacancy diffusion in the space charge layer.37–42 V’Li diffusion was, therefore, studied considering only the most stable Li(2)FP configuration. In the case of Li vacancies, Li diffusion corresponds to the repetitively position exchange of one lattice Li+ and of the vacancy. Lii• was then kept fixed at the interface, while relaxed scan calculations were performed for one lattice Li+ diffusing from its original position to the vacant site in Li(2)FP along a selected path according to a simple direct hopping mechanism (see Figure 1a). From Table S1 and Figure 1b, the vacancy diffusion is associated to a very small barrier (0.23 eV), so that, in the composite, Li transport can easily take place through a vacancy hopping mechanism. A comparison with previous results on the clean ACS Paragon Plus Environment

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LiKCO3-(001) surface34, clearly highlights the role of the interface. Indeed, in this last case the transport was found to be mainly due to Lii• species diffusing via directhopping along the large open channels characterizing the structure of the surface with barriers of 0.80 eV, while Li vacancies were associated to very high energy values (>10 eV) thus excluding their role as effective diffusion carriers in the pure carbonate surface.

Consequently, the interface formation in oxide-carbonate materials not only determines a reduction of the defect formation energy associated to LiFP and of the diffusion barrier, but the interaction between the two phases is also responsible for the change in the Li transport mechanism, since vacancy hopping was forbidden in the pure carbonate material, while V’Li can easily diffuse in the space charge layer in agreement with Maier’s theory in the composite. 3.2. Intrinsic species: the case of O-related carriers It is well known that oxide ion conduction can take place in the oxide phase via an oxygen vacancy mechanism, but different explanations on the oxide ion enhancement mechanism in composites have been proposed in literature. For example, it was suggested in the literature that oxygen ions can accumulate at the surface of the oxide particle, resulting in higher oxygen vacancy concentration in the bulk due to interfacial interaction17 (see Scheme 1.A2). Furthermore, according to the most common explanation based on Maier’s theory, higher concentrations of the defects/ions (negatively charged oxygen species and cations of the carbonate phase) near the phase boundaries may constitute a preferential ’high conductivity pathways’ 16,18,19

(see Scheme 1.A3).

Finally, it is also worth considering the properties of O atom migration at the oxidecarbonate interface, since carbonates may enhance the oxygen reduction process (ORR) in SOFC, through a particular mechanism for O migration in the carbonate phase. The migration species could be CO42- ions while the transport mechanism is a cooperative ’cogwheel’ (or ’paddlewheel’) mechanism, involving the breaking and reforming of O-CO32- bonds43–46 (see Scheme 1.B). It was also suggested that this process can be relevant in the case of SOFC with composite carbonate-oxide electrolytes.45 Although not strictly directly related to the conduction properties of the composite, understanding the effect of the interfaces on the oxygen migration when


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composite electrolytes are used can be very interesting.47–49 Oxygen anion related species To verify the enhancement mechanism proposed by Huang et al.17 (see Scheme 1.A2), neutral O Frenkel Pairs (OFP) were created in the oxide-carbonate composite, by displacing one O2- from its lattice position in the oxide, where a positively charged O vacancy is thus formed (VO••), towards an interstitial position at the YSZ-carbonate interface (OS’’). Different configurations, considering oxygen ions at different lattice sites, have been tested and the results for the most stable configurations are reported in Table S2. The oxide lattice slightly rearranges after the vacancy formation and the oxygen anion interacts at the interface with both the Zr cations of YSZ and the Li+ or K+ species of the carbonate. However, the value of the formation energy of 4.59 eV is too high to explain the conductivity enhancement in terms of higher oxygen vacancy concentration in the bulk oxide due to interfacial interaction with the carbonate. To simulate the mechanism based on higher concentrations of defects/ions near the phase boundaries, simultaneous adsorption of both negatively charged species and of the cations species has been simulated by introducing one O2- and two interstitial Lii• (indicated in the following as [2Li i•-O2−]) at the interface.

Formation energies

ranging between -4.20 and -3.05 eV were obtained, thus suggesting a very favorable accumulation of these species at the oxide-carbonate composite interface. Details on selected parameters for the most stable configurations obtained are reported in Table S2. In the following, only the most stable configuration ([2Lii•-O2−](1)) will be considered. Different types of mechanisms, involving the migration of the oxide ion along both the a and b axis of the unit cell or simultaneous migration of O and Li species, have been considered. In Figure 2, we report the path and corresponding energy profile for the direct hopping mechanism of O anion diffusion along the a-axis of the composite unit cell, which represents the mechanism with the lowest activation energy. More precisely, the energy profile of Figure 2b shows that direct hopping along the a-axis is associated to a barrier of 2.19 eV. Such a high barrier is not unexpected, since conductivity measurements through EIS experiments2,3,6,50 showed that at low temperatures the activation energy for the oxide-carbonate composites is higher than the values typical for the oxide ion conductors (1.6 eV in the case of YSZ-LiKCO3 composites). The barrier of 2.19 eV obtained in the [2Lii•-O2-](1) case can thus be


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partly explained considering the physical state of the carbonate phase that in our model is in an amorphous-like phase, but still close to the solid state. However, this result supports the assumption that the low temperature activation energy of the composites represents the behavior of oxide ions partially inhibited by the presence of solid carbonates.3

Oxygen atom migration Regarding O atom migration, different adsorption sites of an oxygen atom into the composite have first been considered, either at the interface (O(x)int configurations, see Table S3), or in the bulk carbonate phase (O(x)carb configurations, see Table S4). At the interface, four different sites have been taken into account: (i) onto a CO3 involved in a bidentate binding with the substrate (O(1)int), (ii) onto a CO3 involved in a monodentate binding (O(2)int), (iii) on top of a Zr site of the YSZ surface (O(3)int), (iv) into the YSZ lattice in correspondence to an oxygen vacancy site (O(4)int). From the formation energies reported in Table S3, O(1)int is found to be the most stable configuration, with a Ef value lower of at least 0.33 eV than the other ones. When adsorbing O in the carbonate phase onto CO3 groups lying at about 8 Å from the oxide surface, the two most stable configurations obtained are shown in Table S4. Structural properties similar to the O(2)int and O(3)int cases are obtained, with comparable (O(1)carb case) or higher (O(2)carb case) Ef values. A more detailed description of the structure and stability of these configurations can be found in the supporting material. Migration of the O atom from the most stable configuration O(1)int to O(2)int at the interface or from configuration O(1)carb to O(2)carb in the bulk carbonate phase according to the cooperative cogwheel mechanism (see Scheme 1B) has been subsequently studied. Figures 3a and 3b show a schematic path for these mechanisms, while Figure 3c reports the corresponding energy profiles. In the configuration O(1)int, the oxygen transfer starts with the stretching of the O-O bond between the migrating O and the O atom of the bidentate CO3 group. While this bond is elongated, the distance between the migrating O and the acceptor O of the monodentate carbonate is shortened, thus showing a concerted bond breaking and bond formation process. At the maximum of the energy profile (see Figure 3c), an O-O-O moiety is formed, as shown in Figure 3d. The same description holds for the case of oxygen transfer in the


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carbonate phase. However, while migration in the carbonate phase is associated to a barrier of 0.95 eV, a value of 0.74 eV is obtained at the interface. This clearly indicates that this process in the composite is possible, and that it can further be boosted when taking place in the boundary layers between the two phases. In addition, it should be noted that computed barriers in the carbonate phase (0.95 eV) are very close to the values already reported by Lei et al.46 in case of oxygen migration in molten Li2CO3, Na2CO3, and K2CO3 (~1 eV), but that the lowest values are obtained when migration takes place at the interface (0.74 eV), outlining again the key role of the interface in diffusion processes.

3.3. Extrinsic species: the case of H+-related species As already noted above, oxide-carbonate composite electrolytes were also claimed to be multi-ionic conductors made by both intrinsic (M+, CO32-, and O2-) and extrinsic (H+) species.7,10,13,51,16 In particular, since these materials are considered to be dual or hybrid H+/O2- conductors, a lot of efforts have been devoted to the investigation of proton transport mechanism in these materials, despite the fact that there are many contradictory results regarding the actual contribution of H+ to the overall conduction and the huge debate around the origin of the protonic conduction.3,7,52 To investigate the energetics of proton transport in these composite materials, two H+ and one O2- were randomly added at the interface or in the bulk carbonate phase. The most stable configurations are reported in Table S5 together with the corresponding formation energies. All the considered configurations are associated to negative values of the adsorption energies, indicating that this process is favorable, although less favorable than the simultaneous adsorption of [2Lii•-O2-] species described in the previous section (-2.5 vs. -4.2 eV for the most stable configurations in both cases). The first proton conduction mechanism proposed by Zhu and Mellander53 takes place in the carbonate phase through temporal bonding between H+ and CO32- (see Scheme 2.1). Consequently, this proton transfer mechanism was firstly considered, according to the path described in Figure 4a. In this case, H+ transport requires a continuous sequence of intermolecular proton transfer and rotation steps, in analogy with the Grötthus mechanism for proton diffusion in water or in polymeric membranes54,55. From the associated energy profile (Figure 4b), it is clear that the H+ transfer is associated to a small barrier of 0.18 eV, while the rotation requires a higher barrier of


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0.35 eV. These data are in good agreement with the DFT results on proton transfer in crystalline and molten alkali carbonates by Lei et al.56, especially if we consider that in our simulations we are dealing with a mixed (Li, K) carbonate phase in an amorphous-like solid state. It is also worth to notice the similarity with the proton transport in polymeric membrane for Fuel Cells applications, where the conduction is ruled by a transfer of a single proton, followed by a rotation, this latter being, like in the present case, the rate determining step54,55.

However, our data show that the formation of proton-related species is energetically more favorable at the interface, thus inducing us to investigate the energetics of the whole proton transfer at the YSZ-LiKCO3 interface. First, as suggested by Huang et al.17, a simple linear direct-hopping mechanism was considered. According to this hypothesis, the proton adsorbed at the oxygen site of the YSZ lattice becomes an interstitial species at the interface and diffuses from the initial position in the most stable configuration to a final equivalent one (see Figure 5a and Scheme 2.2). Only diffusion along the a-axis has been chosen here, due to the presence of OH groups along the b direction. From Figure 5b, the energy profile obtained for this mechanism shows high barriers (about 2 eV for the highest). Since interaction with the carbonates or with the oxygen species at the surface of the oxide could influence the energetics of the proton transport at the interface, a second path divided into six successive steps and shown in Figure 6a has been taken into account. First, according to the interfacial conducting mechanism suggested by Zhu and Mat13 (see Scheme 2.3 and Figure 6a, steps 1 and 2), H+ first migrates from its initial position on the oxide surface to an interstitial position at the interface where it interacts with the OH group that was introduced together with it, and the proton is then transferred from this interstitial position to another O surface site. Then, as in the ’Swing Model’ suggested by Wang et al.4 (see Scheme 2.4), a continuous hydrogen bond breaking and formation takes place at the oxide-carbonate interface in which the carbonate groups serve as a bridge for the proton to move from one hydrogen bond to the other (Figure 6a, steps 3 to 6). From figure 6b, the maxima of the first (inset M1) and second (inset M2) steps are associated to barriers of 0.23 and 0.33 eV, respectively, corresponding to structures in which the proton interacts with the OH group added at the interface, with O atoms of ACS Paragon Plus Environment

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the YSZ surface and/or of the carbonate groups. The low barriers obtained for these steps suggest that the presence of the oxygen species adsorbed on the surface can introduce an interfacial conducting path through a chain of hydrogen bond which favors the proton conduction. Moving to the successive ‘Swing Model’ part of the path, higher barriers of 1.06 and 1.03 eV are obtained for the rotation in the fourth step and for the transfer in the fifth one. At the corresponding maxima, the proton interacts with both one O of a carbonate group and one O site on the YSZ surface. The energy barriers encountered in the ‘Swing Model’ section of the path are thus higher than those found in the first part, but still definitely lower than the ones obtained in Figure 5b for the linear interstitial mechanism at the interface. These results suggest that the continuous hydrogen bond breaking and formation process involving the oxygen of the oxide ion or of the carbonate groups can indeed influence the energetic of the proton transport at the interface. The value of 1.06 eV is still quite high to explain the conductivity enhancement. Although in the considered interface model the carbonate phase is in a pseudo amorphous form, the ’Swing Model’ of Wang et al.4 was proposed as being particularly important when the carbonate is in the molten state, where the bending and stretching vibration of C-O bonds are enhanced as well as mobility and rotation of the CO32- groups, thus enhancing even more the proton transport. Further details on the steps of this conduction path illustrated in figure 6 can be found on the SI. To sum up, H+ migration in the bulk carbonate through temporal bonding H+-CO32- is associated to very low barriers. However, the formation of proton species appears more favorable at the interface where the proton can interact with the oxygen atoms on the oxide surface. In addition, at the interface, the presence of oxygen species adsorbed on the oxide may help in lowering the proton migration barriers at low temperature, and the migration energy can be decreased by the interaction with the carbonates, acting as a bridge for proton between different oxygen surface sites. It is especially reasonable to expect that this effect becomes more important at higher temperatures where the enhanced movements of the carbonate group in the melted state can favor the transfer.

4. Conclusions We presented a DFT study of the conduction properties of oxide-carbonate composite electrolyte for SOFC applications, considering both intrinsic (Li+ and O2-) and ACS Paragon Plus Environment

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extrinsic (H+) species. The main outcomes provided by this investigation and reported in Table 1 can be summarized as follows: • The results obtained for Li transport in the composite are in agreement with the ’space charge layer’ model proposed by Maier37–42. Indeed, the interfaces in oxide-carbonate composites can reduce the formation energy of Li-related defects because of the chemical affinity of the O atoms of the oxide for these species. Li vacancies, created in bulk carbonate as a consequence of the interfacial interaction, can diffuse through a Li vacancy hopping mechanism with very low barriers in the so formed space charge layer. • The adsorption of negatively charged oxygen species associated with the cations of the carbonate phase was found to be a favorable process, suggesting again that a space-charge layer can indeed be easily formed at the phase boundaries due to the interfacial interaction. Moreover, we were already able to show how the interfaces can boost the neutral O atom migration, in a process that can be important for the composite-based SOFCs. • Interfaces can also facilitate the formation of H-related defects through interaction of H+ with the oxygen sites of the oxide. Here, the presence of the carbonates and especially of oxygen species adsorbed on the oxide play an important role in the lowering of the barrier associated to proton migration at the interfaces even at low temperatures.

These results therefore provide important information on the electrochemical processes and transport mechanisms in these composite electrolytes, thus paving the way toward a deeper understanding of the origin of their enhanced electrochemical properties and suggesting general guidelines for their improvement.

Supporting information Additional details relative to the different defect configurations and conduction mechanisms tested for each of the considered transport species in the YSZ-LiKCO3 composite electrolyte.

Acknowledgments


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This work has been realized in the frame of IDEX, reference ANR-10-IDEX-0001-02 PSL. The authors acknowledge the HPC resources of TGCC made available by GENCI (Grand Equipement National de Calcul Intensif) through project x2016087623.


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Scheme 1: A) Space charge layer created by 1) the cations of the carbonate phase17, 2) the oxide anions of the oxide17, and 3) the adsorption of cations of the carbonate phase and oxide anions at the interface16,18,19 . B) Cogwheel mechanism for the transport of the neutral O atom in the carbonate phase. 43–46


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Scheme 2: Suggested proton transport mechanisms in the composites: 1) migration in the carbonate phase through transitory H+-CO3- species 53; 2) interstitial proton formation at the interface17; 3) conduction H-bond chain at the interface13; 4) Swing Model made by a continuous H-bond breaking and formation taking place at the interface4.


16

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Species

Ef/ads

Ebarr

Mechanism type

LiFP

1.07

0.21

Direct Hopping

OFP

4.59

-

-

Localization of the Migration Carbonate Phase (space charge layer) -

O2-

-4.20

2.19

Direct Hopping

Interface

O

0.96

0.74

Cogwheel

Interface

1.27

0.95

Cogwheel Intermolecular H+ transfer Rotation Direct Hopping (Interstitial) Non Direct Hopping (Conduction chain H---O-H, H-O-H, and O-H---O)

Carbonate

0.16 -1.76 0.35 2.19 H+ -2.52

0.33

1.06

Carbonate

Interface

Non Direct Hopping (Swing Model)

Table 1: Summary of the information relative to the conduction in the oxidecarbonate materials. Formation or adsorption energies (Ef/ads, in eV) of all the possible conductive species taken into account in this work (cations of the carbonate (Li+), oxide ions, oxygen atom, and proton) together with the values of the barrier (Ebarr, in eV) associated to the migration of these species along selected diffusion paths, the relative mechanism type and the place where the transport is taking place.


17


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Figure 1: a) Schematic representation of the path for the direct hopping diffusion of a V’Li (hashed light blue sphere) in the bulk carbonate phase when its counter species (Lii•, dark blue sphere) is trapped at the interface after the LiFP formation in YSZ-LiKCO3. b) Variation of the energy as a function of the distance travelled by the V’Li specie along the path.

Figure 2: a) Views along the a-(left) or b-axis (right) of the direct hopping along the a-axis of a O2- anion diffusing at the interface of the YSZ-LiKCO3 composite from the initial position (in orange) in configuration [2Lii• -O2-](1) to a final equivalent position (in black). The interstitial Li cations are represented in blue. b) Variation of the energy as a function of the distance travelled by O2- position along the selected path.


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Figure 3: Cogwheel mechanism for the O atom migration at the a) interface from configuration O(1)int (in yellow) to configuration O(2)int (in black) or b) in the carbonate bulk phase in YSZ-LiKCO3 from configuration O(1)carb (in yellow) to O(2)carb (in black). c) Variation of the energy as a function of the distance travelled by the O the corresponding path. d) Structure of the maximum along the path for the migration at the interface.


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Figure 4: a) Schematic representation of the proton conduction path in the bulk carbonate phase of the YSZ-LiKCO3 composite. Initial and final proton positions are indicated by light blue and black spheres, respectively, while gray spheres represent the intermediate positions along the path. b) Corresponding variation of the energy as a function of the distance travelled by the proton along the path (in Å for the intermolecular proton transfer and in ° for the rotation steps). The vertical dotted line separates these two steps detailed in the text.

Figure 5: a) Schematic representation of the linear direct-hopping of a proton, from its initial configuration (Hint(1)) to a final equivalent one. b) Corresponding variation of the energy as a function of the distance travelled by proton the path. The initial and final positions of the proton are indicated by the light blue and black spheres, respectively.


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Figure 6: a) Schematic representation of the non linear direct hopping of a proton at the interface of the YSZ-LiKCO3 composite. Blue and grey spheres represent the protons, and their intermediate positions along the path. b) Corresponding variation of the energy as a function of the distance travelled by the proton along the path. The insets show the structure of the maxima encountered along the path.


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TOC Graphic Load

Oxide-Carbonate ELECTROLYTE

e-

e-

O2-

? ELECTROLYTE

CATHODE

H2 O ANODE

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 26 of 35

H2

Li+ O2

H+

O2-

DFT


26

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