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Local structure and lithium diffusion pathways in Li4Mn2O5 high capacity cathode probed by total scattering and XANES Maria Diaz-Lopez, Melanie Freire, Yves Joly, Claire V. Colin, Henry E Fischer, Nils Blanc, Nathalie Boudet, Valerie Pralong, and Pierre Bordet Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b00827 • Publication Date (Web): 19 Apr 2018 Downloaded from http://pubs.acs.org on April 19, 2018

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Chemistry of Materials

Local structure and lithium diffusion pathways in Li4Mn2O5 high capacity cathode probed by total scattering and XANES Maria Diaz-Lopez1*, Melanie Freire2, Yves Joly1, Claire V. Colin1, Henry E. Fischer3, Nils Blanc1, Nathalie Boudet1, Valerie Pralong2 and Pierre Bordet1. 1 Université Grenoble Alpes, Institut Néel, F-38000 Grenoble, France CNRS, Institut Néel, F-38000 Grenoble, France. 2 Laboratoire de Cristallographie et Sciences des Matériaux CRISMAT, ENSICAEN, Normandie Université, CNRS, Université de Caen Normandie, 6 Bd Maréchal Juin, F-14050 Caen, France. 3 Institut Laue-Langevin, 71 avenue des Martyrs, CS 20156, 38042 Grenoble cedex 9, France. *e-mail : [email protected]

Abstract In the constant race for more efficient Li-ion batteries, extensive research has focused on the design of new, more competitive cathode materials, currently limiting the battery performance. The improvement of cathode materials demands the detailed understanding of the complex structural mechanisms at play during battery operation, i.e. when Li+ ions are inserted and extracted from the cathode. Moreover, new cathode designs involve more and more disordered/nano-sized materials for enhanced Li+ cation diffusion and larger specific surfaces. This trend poses new challenges for the structural investigation methods employed, which mostly rely on the periodic and long range ordered nature of the compounds under study. This is specially the case of the recently discovered nanostructured Li4Mn2O5 high capacity cathode material, which shows record reversible capacities superior to the state-of-the-art Li-Mn-O electrodes and displays a strongly disordered rock salt-type structure. This last feature, mainly due to its synthetic route involving high energy milling, prevented from reaching a full understanding of the lithium exchange mechanism of 1 ACS Paragon Plus Environment

Chemistry of Materials 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

particular interest in this 3D framework compound. Here, we demonstrate that a thorough description of such a disordered structure can be achieved by a combination of Near-Edge X-ray Absorption Spectroscopy and Pair Distribution Function analysis of neutron and X-ray total scattering data, which ultimately lead to the elucidation of the Li cation diffusion pathways.

Text The raise of lithium-ion rechargeable battery technology as a promising candidate to meet the current demands for clean and renewable energy sources1 has led to an extensive research on the design of new, more competitive materials with improved performance2. Amongst the several cathode materials explored, most efforts were focused on transition metal oxides intercalation compounds with 2D structures having acceptable capacities (100-250 mAhg-1) and operating voltages (2.5-5 V) 3. Recently, the focus has changed to layered Li-excess oxides with even higher capacities (250-300 mAhg-1)4–7. Cells with cathode materials based on spinel 3D structures also have adequate performances8–10 (e.g. LixMn2O4 has a reversible capacity of 120 mAhg-1 over extensive cycling)11,12, which proves that Li-conduction Is not exclusive to layered materials. The recent discovery of Li1.25Nb0.25M0.5O2 with M = V13 and Mn14, Li1.9Mn0.95O2.05F0.9515 and Li4Mn2O516 rock-salt-type compounds with high capacities constitute new additions to the less studied family of cathode materials with 3D frameworks. Although reported to be efficient, the migration of Li in the non-stoichiometric Li4Mn2O5 (Li4) rock-salt structure is not yet understood, partly due to the lack of a detailed knowledge of this disordered structure at the atomic level. In this article, we present a comprehensive study of Li4’s structure to unravel the mechanism by which lithium is able to diffuse in this new and intriguing compound. Previous Transmission Electron Microscopy (TEM) and combined neutron and X-ray powder diffraction studies of Li4-xMn2O5 (x=0 and 4)16 revealed this material shows on average an MnO-type rock-salt structure, albeit highly distorted at the local scale due to the substitution of 2/3 manganese for lithium and the presence of 1/6 oxygen vacancies. In fact, the degree of disorder is such that the PDF data up to 12 Å could no longer be described with the rock-salt model17. The full structural description of Li4 is greatly challenged by its high degree of complexity involving the presence of oxygen vacancies, substitutional disorder on the cation site and nanostructuration. In this work, we approach the structural study of Li4 by the combination of Near-Edge X-ray Absorption Spectroscopy (XANES) at the Mn K-edge and combined neutron and X-ray total scattering measurements. First, the coordination environment of manganese was extracted from XANES, and this information was used to build the starting model for the Pair Distribution Function (PDF) analysis of the total scattering data. Then, 2 ACS Paragon Plus Environment

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the PDF analysis was carried out using the RMC modelling method, capable of generating a high degree of local disorder from highly ordered starting configurations18–21. Several refinements were performed on a ‘large-box’ with dimensions similar to the structural coherence length of Li417, simultaneously fitting both neutron and X-ray reduced pair distribution G(r) and total scattering S(Q) functions. The RMC-refined model was finally validated by the comparison of simulated and experimental XANES spectra. Additionally, a quantitative analysis of the RMC outputs was performed to understand the distribution and local environments of the Mn-framework and the mobile Li ions in the validated model. Finally, we introduce the Bond Valence Site Energy (BVSE) analysis of the RMC-refined model, where accessible Li+ sites were identified by the calculation of the mismatch of the Bond Valence Sum22 (BVS) with respect to Li+ formal valence23,24 and converted to energy units25. The BVSE map of Li revealed a plausible 3D Li diffusion pathway connecting large volumes of 5-coordinated lithium clustered around oxygen vacancies.

Results The starting model In order to propose a sensible initial configuration for RMC refinement, all the available structural information on the material was gathered and thoroughly analysed. This included real and reciprocal space Rietveld refinements17 and XANES data at the Mn K-edge. Following the subtraction of the Li2O impurity from the diffraction pattern, the composition of the main phase was also corrected. To this end, the ICP results referring to the overall composition including the main and impurity phases were inputted into the Rietveld refinement. The composition of the main phase was then determined to be Li3.6(2)Mn2.4(2)O5.4(2). The initial model was built based on this corrected composition, although we continue to address Li3.6Mn2.4O5.4 as ‘Li4’, as opposed to ‘Li3.6’, for clarity. The Li3.6-xMn2.4O5.4 x-PDF data showed how Li removal is accompanied by an isotropic contraction (breathing) of the cubic Mn-framework. Surprisingly, Li extraction can be carried out without the occurrence of major structural rearrangements as evidenced by the constant value of the refined coherence length. However, no detailed local structural information, particularly on the atom coordination environments could be extracted from x-PDF data alone due to the overlapping of all cation-cation, oxygen-oxygen contributions in this highly disordered rock-salt structure and the limited sensitivity to Li scatterers. Thus, we included neutron total scattering data, with a higher sensitivity to presence of the lighter Li and O atoms, in the PDF analysis of Li4. Additionally, we have probed the coordination environment of Mn by the analysis of XANES data performed at the Mn K-edge on Li4 and delithiated Li017. 3 ACS Paragon Plus Environment


The mere visual inspection of the XANES data (see Figure 1) reveals a shift to higher energies of the Li0 spectra due to an increase in the oxidation state from Mn3+ in Li4 to Mn~4+ 16 (see Figure S1). In both Li4 and Li0 spectra very similar pre-edge and overall spectra shapes are observed, albeit the expansion of the Li0 spectra in the energy scale (i.e. the separation between Li4 and Li0 peak maxima increases with the increasing energy). These observations suggest there is a high resemblance between the local structures of the two compositions and a contraction of the lattice for Li0 according to Natoli’s rule26, which is in good agreement with the smaller cell parameters of Li0 determined by real and reciprocal space Rietveld refinements17.

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E(eV) Figure 1. (a) Mn K-edge XANES of Li4 (green) and Li0 (purple). (b) Data expanded in the energy scale showing the flattening of the spectral features of Li4 and Li0 with respect to highly ordered MnO (black) denoting the presence of structural disorder. The insets in both figures are the magnified pre-edge regions. 4 ACS Paragon Plus Environment

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Since the average structure of Li4-xMn2O5 can be described with the MnO rock-salt structure, its spectra is also introduced in Figure 1 for comparison. The significant differences between the spectra denote that the local structure is different to that of a rock-salt, in good agreement with the high mismatch observed between the PDF data and the average model up to ~12 Å17. Moreover, the high degree of local disorder in this cathode material is further supported by the flattening of the spectral features in Li4 and Li0 with respect to MnO (see Figure 1.b). Although small in amplitude, the pre-edge signal carries valuable information regarding the coordination environment of the absorber atom, i.e. Mn. For the K-edge of the first row transition metals with an open 3d shell, the 1s to 3d transition is responsible for the pre-edge signal (see Figure S2). Since this transition is forbidden by dipole selection rules its intensity is small for centrosymmetric environments, and strongly increases in intensity as progressively moving to non-centrosymmetric environments due to the hybridization between the 3d and 4p orbitals. Therefore, the intensity of the pre-edge for an octahedron is smaller than for a square pyramid and becomes rather large for a tetrahedral environment. See for example, the difference in intensities of the pre-edge in Mn2+O and in Ba3Mn5+2O8 with a relative intensity of 7 and 50 % respectively in Figure S3. The small intensity of the pre-edge signals in Li4 and Li0 confirmed the absence of tetrahedrally coordinated Mn. The presence of Mn in a tetrahedral coordination would entail rearrangements of the structure which would be detrimental for electrochemical cyclability and were neither observed on the short or long r-range PDF data, nor were they supported by the similar coherence lengths of Li4 and Li0. Alternatively, we considered the possibility of having Mn in octahedral, square pyramid and square planar environments, which may be expected in the presence of oxygen vacancies. Several models were built starting from a 2×2×2 supercell of the rock-salt type MnO cell, which is the smallest cell allowing to accommodate the Li3.6Mn2.4O5.4 stoichiometry of Li4, and their XANES spectra were simulated (see Figure 2). These model supercells include Li/Mn substitutional disorder on the cation sites and oxygen vacancy disorder on the anion sites without any atomic displacement from the average MnO structure. The best agreement with the data was obtained with the model for which Mn was found exclusively in an octahedral environment. The distance between the two pre-edge peaks, ~2 eV apart, is also in good agreement with the typical energy gap for octahedral configurations. The broadening of the preedge peaks was attributed to the distribution in energy of the t2g and eg states given by the 8 non-equivalent Mn sites in the initial model.

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Figure 2 (a) Pre-edge region of several simulations with Mn in different configurations: square pyramid (pink), square planar (blue), a 1:1 ratio mixture of octahedral and square pyramid (orange) and purely octahedral (black) compared to the Li4 data (green). (b) Partial Density of States (DOS) of the 8 nonequivalent Mn sites in the octahedral model showing a wide distribution of t2g and eg energy levels responsible for the broad pre-edge peaks. In both figures the energy of the Fermi level (EF) are denoted by vertical dashed lines. As a consequence of the octahedral coordination of manganese, lithium is clustered around the oxygen vacancies where it is at most 5-coordinated. In the starting model, the oxygen vacancies were distributed on the anion network so as to avoid the formation of large clusters of lithium. Such configuration is probable to allow the reversible lithium insertion/extraction for several cycles, as experimentally observed. The resulting initial model with the composition Li19Mn13O29, a = 8.332 Å with a Pm space group was expanded to form a large box for subsequent RMC refinement. This box with a = b = c = 58.324 Å (close to the coherence length


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of ~ 6 nm) contained a total of 20 923 atoms. It should be noted that the lattice parameters used here were determined by the combined Rietveld refinement of neutron and x-ray data with a very high accuracy.

The RMC Refinement In order to lead the RMC modelling algorithm into outputting physically and chemically sensible refined models minimum distance cut off constraints and soft BVS constraints were included in the refinement. The choice of cation-oxygen cut-off constraints was performed on the basis of the shortest distance values extracted from the G(r) function (i.e. 1.7 Å). For other atom pairs with overlapped signals, the choice of cutoff distances was performed on distances observed for related Li-Mn-O systems, subtracted by ~0.2 Å. BVS constraints were used as implemented in RMCProfile27 with R0 and b bond valence parameters for Mn3+, Li+ and O2- taken from 28. The refinement was further constrained by the simultaneous fitting of all the available data i.e. neutron and x-ray G(r) and S(Q), allowing for the short, medium and long range structures to be simultaneously modelled. Several tests were carried out before the final refinement to determine the correct weighting of each data sets. Once the refinement parameters were optimized and after ~40 ·106 atom movements were tested the overall goodness of fit of the refinement χ2 of ~8.5 did not decrease any further. Thereby refined models produced excellent fits to all data sets (see Figure 3).



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Figure 3. RMC fits of Li4 neutron (top) and X-ray (bottom) G(r) (left) and S(Q) (right) at 297 K.

Quantitative evaluation of the local structure In this section we introduce the quantitative analysis of the geometry of the constituent polyhedra of the outputted RMC refined models. The reported values account for the study of a total of ten refined models providing partial correlation functions with good statistical quality. It should be noted that each refinement produced consistent and reproducible results (to verify this, the interested reader is referred to Figure S4). We first evaluate the reliability of the refined models by computing the BVS of each atom (see Figure 4). A narrow distribution (~±0.5 v.u.) around the expected BVS values was observed for oxygen (1.9 v.u.), 6coordinated manganese (2.8 v.u.) and 5-coordinated lithium (1.0 v.u.). The BVS of the minority 6coordinated lithium (~5 %) of 1.1 v.u. was further away from its ideal value, yet reasonable. The calculated BVS plots as a function of the coordination number indicated the quasi-absence of 4-coordinated manganese and lithium, as expected by the absence of a peak in the PDF at these shorter cation-oxygen distances. The appearance of 4-coordinated sites and 5-coordinated manganese in the refined model, which are small in number and have broad BVS distributions, could be related to the thermal vibrations of the atoms around their ideal positions. Although the formation of such sites with a lower coordination number is captured in the RMC model, its presence is neither statistically meaningful nor relevant for the interpretation of the refined model. Therefore, a quantitative evaluation of only 6-coordinated manganese and 6 and 5coordinated lithium was carried out. 8 ACS Paragon Plus Environment

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Figure 4. BVS plots for (a) oxygen in grey and for (b) manganese (c) lithium in grey, orange and blue for 6, 5 and 4-coordinated sites respectively. The vertical red lines indicate the initial BVS values in the starting model. For the interpretation of the RMC outputted model the partial atomic pair distribution function gij(r)s is given in Figure 5.a, where we observe the contribution of the different pairs of atoms to the PDF. Here, we can clearly identify the sharpness of the Mn contributing peaks with respect to Li and O, denoting a higher degree of ordering within the cubic Mn-framework, corroborating our expectations17. This observation is in good agreement with the smaller shifts between the refined atomic positions and the initial positions within the rock-salt structure calculated for Mn with respect to the larger shifts observed for Li and O (see Figure 6.b).



Figure 5. (a) Partial atomic g(r)s of the Li4 RMC refined model at 297 K. (b) Lognormal distribution of shifts between the atomic positions in the refined and starting models. The calculated modes are 0.48, 0.41 and 0.26 Å and the mean values 0.67, 0.65 and 0.57 Å for O, Li and Mn respectively. The first two peaks in Figure 5.a centred at ~2 Å account to the nearest Li-O and Mn-O neighbours’ distances (see also Figure 6.a and b), where once more the sharpness of the Mn-O peak is in clear contrast with the broader Li-O peak. For Mn-O, a narrow distribution centred at 2.01 Å and a smaller peak at a longer 2.31 Å is observed. For Li-O, a broader and asymmetric peak with a maximum at 1.87 and a shoulder at 2.15 Å is observed for 5 and 6-coordinated sites (see Figure 6.b).


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0.4

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σ M-O ( )

Figure 6. Histograms of all Mn-O (a) and Li-O (b) distances in the refined model. From left to right: distribution of M-O distances, M-O distances averaged per M site and corresponding standard deviations (σMO)

given as a function of the coordination number.

We now proceed to evaluate the Mn/Li-O distribution of distances as a function of the coordination number. To quantify the degree of distortion of the octahedrally coordinated Mn and Li sites, we calculated the Minimum Bounding Ellipsoid (MBE)29 of every octahedron as follows. First, the ellipsoid principal axis lengths radii were calculated and ordered as R1≥R2≥R3. Then, the S parameter, where S = R3/R2 – R2/R130, was used to address the shape of the octahedra in the refined model as: axially compressed or prolate (S0) or regular (S=0). For a representation of the ellipsoidal fits of distorted octahedral the interested reader is referred to Figure S5. In Figure 7, the plot of S for Mn shows a narrow and symmetric distribution around S = 0, which indicates the longer Mn-O distances of ~2.31 Å are randomly distributed and do not give rise to prolate or oblates distortions in Li4. This observation was also supported by the smooth distribution of average Mn-O distances in Figure 6.a, and BVS in Figure 4.b. For Li, as well as for Mn, a narrow and symmetric distribution of the S value around zero is observed, which suggest that Ohcoordinated Li sites also tend to have spherical shapes.



Figure 7. Histograms of calculated SMn (a) and SLi(b). The red vertical line indicates the initial value of S = 0 in the starting model The majority of Li (95 %) was 5-coordinated with square pyramidal geometry. To better understand the distortions in the 5-coordinated Li sites in the refined model, we evaluated the relaxation of Li and O atoms surrounding the oxygen vacancies (□), whose positions were fixed during the refinement. The distribution of refined □-O and □-Li distances is introduced in Figure 9, which shows that while the average refined values of 2.95 and 2.11 Å are identical/close to the initial values of 2.95 and 2.08 Å for □-O and □-Li respectively, the distribution of □-O distances is not symmetrical. Thus, from the inspection of the distribution of □-O in Figure 9, we can conclude that oxygen within the pyramids base are drawn in towards the vacancy as given by the shift of the maxima from 2.95 Å in the starting configuration to 2.81 Å. On the other hand, the refined □-Li distances displayed an almost symmetrical distribution around the initial value of 2.08 Å.


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Number of -Li

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distance (Å) Figure 8. Histograms of □-Li (a) and □-O (b) distances. The red vertical lines indicate the initial values on the starting model. The small deviations of the □-Li distances from their initial values does not explain the splitting of Li-O distances in the refined model. To find the origin for the distribution of short and long Li-O distances centred at 1.87 and 2.15 Å in Figures 5.a and 6.b, we computed the projection of the Li shift along and perpendicular to the □-Li direction (See Figures 9.a and b). The projection of the refined lithium position along the □-Li direction shows a symmetrical distribution centred at zero, which is in contrast with the larger log-normally distributed shifts (mode = 0.25 Å, mean = 0.62 Å) perpendicular to this direction. In order to compensate its BVS, the 5-coordinatd Li was found to move towards some of the oxygen within the square pyramid’s base and away from others, resulting in the formation of largely distorted square pyramids with a wide distribution of Li-O distances. We quantified the direction of the observed shift perpendicular to the □-Li direction by projecting the refined lithium position in the pyramid’s base and computing the angle of the shift in Figure 9.c, where 0° accounts for displacement of Li along one oxygen atom and 45° is perpendicular to the pyramid’s base edge. Figure 9 shows a homogeneous distribution of Li over the surface of the pyramid’s base which points to a random displacement of Li in every direction.



Figure 9. The displacement of lithium was calculated with respect to its initial configuration at the centre of the base of the square-based pyramid of oxygen as illustrated in the top right figure. Red, shaded and full green spheres denote O, Liinitial and Lirefined respectively.

The displacement vector from Liinitial to Lirefined (in

black) was projected along and perpendicular to the vacancy-Liinit direction. (a) Projection of the refined lithium position along the □-Li direction, where positive values denote the displacement of lithium towards the oxygen vacancy, and (b) perpendicular to □-Li. (c) Histogram of calculated angles for the projection of the refined Li position in the pyramid’s base (where 0° accounts for a displacement along one oxygen atom and 45° is perpendicular to the pyramid’s base edge.

Confirmation of the validity of the RMC refined model In the XANES region the photoelectron has a low kinetic energy and is able to travel a long mean free path interacting with a large number of atoms. For this reason, the XANES data at the Mn K-edge are not only


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useful to determine the local symmetry and change in oxidation state of Mn from Li4, but can also serve us as a whole-structure probe. We have used approximately a quarter of the volume of the box for a single configuration probing a total of a thousand manganese sites to reproduce the experimental XANES spectra introduced in Figure 10 (see model A). The good agreement between the simulated and experimental XANES and the significant improvement of the simulation using the refined model with respect to the initial one (see Figure S6) are a robust proof of the validity of the RMC refined model.

Figure 10. (a) Experimental and simulated XANES spectra using the refinement strategies A-D. (b) 3 × difference calculated for the simulated spectra with respect to the experiment, where the vertical dashed line indicates the position of the white line. Note that the changes in the simulated spectra are more significant to the left of this line corresponding to the pre-edge region. The creation of a sensible starting configuration for RMC modelling was key to the success of the refinement. Although in principle, the RMC method is capable of fitting a given set of data from any starting configuration19; in practice, this approach could either be too time consuming to be feasible or, as in this case, it could be that the refinement of the diffraction data alone is not able to discriminate the correct 15 ACS Paragon Plus Environment


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coordination environment of manganese. We introduce here the two refinements of the initial models in Figure 2.a with 50 % and 100 % of 5-coordinated manganese (namely B and C in Figure 10 and Table 1) using the same refinement approach as in model A; and a third refinement where we have used model A as the starting configuration, but where the swapping between lithium and manganese was allowed during the refinement (refinement approach D). From the results shown in table 1 we conclude that in this particular case, the RMC modelling approach can successfully generate disorder within the model to fit the data regardless of the coordination environment of manganese in the starting model. When the swapping between cations was allowed in refinement D, the refinement tends to homogeneously distribute Mn and Li amongst the 5-coordinated sites, also providing good fits to the data. Even though all the refinements in A-D produced good fits (notice the very similar overall χ2 in Table 1), the refined model produced in the refinement strategy A was the only approach which successfully described the XANES pre-edge and generated chemically plausible models (see refined BVS values in Table 1). Table 1 Goodness of fit and calculated BVS for the RMC refinement using different initial configurations.

Refinement Initial % of

Cations

ID

swapping

5-

χ2

Initial/Refined BVS Mn

Li

O

coordinated (Y/N) Mn

A

0

N

8.5(1) 2.5062/2.719(3) 0.9353/1.041(3) 1.7482/1.901(1)

B

50

N

8.6(1) 2.2813/2.609(3) 1.0230/1.078(2) 1.6929/1.875(1)

C

100

N

8.5(2) 2.1207/2.529(4) 1.0726/1.110(2) 1.6534/1.861(1)

D

0

Y

8.1(7) 2.5062/2.644(9) 0.9353/1.062(4) 1.7482/1.881(3)

3D Diffusion pathway of lithium The clustering of lithium in a square pyramid coordination around the oxygen vacancies together with the systematic displacement of these atoms away from the centre of the pyramid’s base lead to the formation of highly distorted □Li6 environments (i.e. the octahedron formed by the 6 Li atoms surrounding a vacant site). Moreover, the direction of the lithium shifts for given □Li6 environments seems random, as illustrated in Figure 9.c. Therefore, the better suited BVSE map calculation approach introduced in 2.4 was undertaken for the prediction the most plausible lithium diffusion pathways in Li4. This approach exploits the analogy between the squared BV mismatch and the Morse-type potential31, thereby transforming valence units into energy units. As for BV mismatch maps32, the graphical


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representation of isosurfaces with constant BVSE can provide a clear image of the lithium diffusion pathways in the material studied33–35.




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Figure 11. Graphical representations of isosurfaces with BVSE of 2.5 eV along a, b and c axes (from top to bottom respectively) in a Li4 4ax4ax4a supercell section of the RMC refined ‘big box’. Note that in the graphical representation of BVSE maps the energy threshold was lowered from 3 eV used in the calculations in Table 1 to illustrate the percolation more clearly. Green, purple and red spheres denote Li, Mn and O atoms respectively. The BVSE map in Figure 12 was performed in a 4ax4ax4a (where a is the lattice parameter of the ‘MnO’ cell) section of the RMC refined model introduced in 3.3 containing 488 atoms. The yellow isosurfaces represent regions in which Li+ could diffuse for a 2.5 eV energy threshold. This choice of energy threshold revealed an extensive 3D pathway network for Li+ transport where all the 5-coordinated lithium, constituting 95 % of the total lithium content, becomes accessible. The accessible volume elements belonging to the same ‘pathway cluster’ consisted of ‘…-Li-□-Li-Li-□-Li-…’ intercrossing chains, which percolate along the three axes. The pathway clusters shared common faces and edges with vacancy-lithium octahedra ‘□Li6’. The diffusion mechanism in Li4 can be better understood when looking at □Li6 environments individually. In Figure 13.b it can be clearly noticed how the Li+ isosurfaces describe a sphere around the oxygen atoms bridging the different Li in □Li6. The perspective side view of □Li6 in Figure 13.d shows how such spheres intersect in the empty tetrahedral interstices. Thus, the transport of Li+ between three lithium sites within the □Li6’s faces could occur via the momentary occupation of the neighbouring vacant tetrahedral interstices. The Li+ diffusion mechanism in Li4 bears a resemblance to that of other rock-salt type high capacity cathodes lacking the presence of oxygen vacancies (e.g. α-LiFeO2 or spinel-like LiCoO236,37). The crystal structure of these layered and spinel transitional metal oxides consist of a close-packed oxygen sublattice with octahedral and tetrahedral interstitial sites occupied by TM /Li with an ABC stacking sequence. In these structures , the lithium migration is reported to occur between two octahedral sites via a tetrahedral activated state38, mediated by the presence of cation vacancy clusters (divacancies for layered materials and di- and triple vacancies for spinels39). However, due to presence of oxygen vacancies in Li4, lithium migration rather occurs between two square-based pyramidal sites via a trigonal planar activated state, and contrarily to the aforementioned crystal structures, the presence of cation vacancies does not play a role in the lithium diffusion mechanism for this particular compound. The existence of pathway clusters alone is not sufficient to sustain Li+ migration throughout the material. Such clusters need to be arranged in a percolating network40,41 to enable long-range lithium transport. To fully address the lithium migration in Li4 we calculated the percolation energy, as the minimum energy above the ground state connecting the BVSE isosurfaces along a particular direction in the unit cell, and the accessible volume fraction (F), as the fraction of the unit cell that belongs to the infinite pathway33 for a 3 eV energy threshold. The promising percolation thresholds and F-values for Li4, comparable to those for 19 ACS Paragon Plus Environment


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representative state-of-the-art cathode materials introduced in table 1, point out the high potential of Li4 warranting further investigations into this system.

Figure 12. (a, c) □Li6 connectivity and (b, d) superimposed BVSE map, Li-O connectivity and tetrahedral interstices (black spheres). The black dashed lines in d denote the intersection of three spheres centred at the neighbouring O. Table 2. Calculated volume fraction for ion mobility, minimum and percolation energies for several representative lithium transition metal oxides from models reported in the literature (column Ref.). The calculated BVSE maps are given in Figure S7.

Percolation E (eV) along

Ref.

Material

F (%)

Minimum E (eV)

a

b

c

Li4

14(3)

-5.3(4)

2.0(6)

2.0(6)

2.0(6)

This work

LiMn2O4

13.28

-4.4175

0.54

0.54

0.54

42


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LiNi0.5Mn0.5O2

17.13

-3.1193

0.30

0.30

0.30

43

α-MnO2

19.44

-3.0518

No

No

0.18

44

LiMnBO3

4.5

-8.9629

No

No

1.81

45

LiCoO2

11.17

-5.1456

1.96

1.96

No

46

LiFePO4 olivine

9.86

-4.586

No

1.04

No

47

V2O5

10.48

-4.349

No

1.38

2.26

48

Li2FeSiO4

19.44

-3.8476

0.96

No

0.96

49

LiV3O8

11.67

-5.0514

No

1.34

2.46

50

Conclusion The complete understanding of the local disorder in nanostructured Li4 was reached via the RMC modelling of combined neutron and X-ray total scattering data. The success of the refinement relied on building a sensible starting model based on the knowledge of the average structure from diffraction experiments, the local environment of Mn from XANES and the implementation of an intelligent set of geometrical and chemical constraints. The analysis of the RMC refined model indicated that Mn is only present in an octahedral coordination that while distorted, is on average spherically shaped. Li is clustered around the oxygen vacancies and it is predominantly 5-coordinated with an unusual and highly distorted square pyramidal geometry. The 5coordinated Li is displaced away from the centre of the pyramid’s base leading to a wide distribution of Li-O distances. BVSE landscapes revealed large pathways for Li-diffusion surrounding all 5-coordinated Li sites which amounted to 95 % of the Li content constituting ~60 % of the cations in Li4. Moreover, these large isosurfaces percolated along the three axes providing Li4 with a unique 3D Li-ion diffusion network only ever reported for spinel electrodes.

Methods Synthesis. 5 g of nanostructured Li4 were synthesized by a two-step route, the first of which involved the synthesis of HT-LiMnO2 by the solid-state reaction method. Stoichiometric amounts of MnO, MnO2 and LiOH (with 5 wt% excess to compensate for evaporation at high temperatures) were mixed and heated at 1000 °C under an argon flow for 8 h. The second step involved the mechanical grinding of stoichiometric amounts of 21 ACS Paragon Plus Environment


HT-LiMnO2 and Li2O in a 20 ml capacity WC bowl with four WC grinding balls (with a 10 mm diameter) using a Fritsch Planetary Micro Mill PULVERISETTE 7 premium at 700 rpm for 10 h. Li4 was electrochemically oxidized versus lithium in a Swagelok type cell. After the first electrochemical charge to 4.8 V all lithium was subtracted from Li4. The battery was then paused and 20 mg of delithiated material (Li0) powder were recuperated and washed with acetonitrile for XAS experiment. The Li0 sample studied here corresponds to that in 17 , the Li:Mn ratio in ‘Li0’ is equal to 0.2±0.05:1 as determined by ICPOES. Near-edge X-ray Absorption Spectroscopy and FDMNES modelling. XANES data at the Mn K-edge was acquired at the French CRG-FAME beamline (BM30) at the European Synchrotron Radiation Facility (ESRF). The beam energy was selected using a cryo-cooled double crystal monochromator whose energy value was calibrated with a Mn-foil standard. A total of seven samples including five Mn standards in different oxidation states (Mn2+O, Mn3+2O3, Mn4+O2 and Ba3Mn5+O8), oxidized Li0 and pristine Li4 were diluted (~5 wt%) with boron nitride and measured in quartz capillaries. The capillaries for the air sensitive Li4 and Li0 were carefully prepared under inert conditions and epoxy sealed. The spectra were collected at room temperature using a 30-element Canberra Ge solid state detector. All the spectra herein introduced correspond to at least six scans that were averaged, background and self-absorption corrected and normalized with the ATHENA software51. Our data processing protocol is supported by the good agreement between our spectra of the standards (see Figure S3) and others found in the literature52. Experimental XANES spectra were simulated with the FDMNES software53 using the multiple-scattering theory under the muffin-tin approximation on the potential shape, i.e. the potential is spherically averaged in the atoms and constant between them. The code runs in direct space and the calculations were limited to a cluster embedded in a sphere centred on the absorbing atom. When several non-equivalent atoms were present in the unit cell, the corresponding number of simulations were performed and summed up to get the total spectra. We have first checked the convergence using increasing radius and found that for cluster size of 6.1 Å around the absorber atom corresponding to ~90 atoms, the spectra of the Mn3+2O3 standard was successfully reproduced (see Figure S8). Simultaneous RMC modelling of n-PDF and x-PDF. High-Energy Synchrotron X-ray powder diffraction (HESXRPD) patterns with high spatial resolution and counting statistics were acquired at the French CRG-D2AM beamline (BM02-ESRF) using λ = 0.4959 Å radiation for subsequent Rietveld and PDF analysis. The samples were prepared under inert conditions in epoxy-sealed borosilicate glass capillaries with an outside diameter of 0.7 mm. The data were collected at room temperature in the 0.8-128.0º 2θº range with 2 h of data


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collection time using a 2D pixel detector XPAD354. In order to collect this wide angular range, 120 diffraction images were taken at each degree 2θº. They were assembled, averaged and rebinned using the PyFAI55 software to produce a single 1D pattern. In addition, the diffraction pattern of an empty capillary was also measured. The data were calibrated using a LaB6 standard collected for 4 h. A Ni standard sample was also used to calibrate the experimental effects on the PDF data. The HE-SXRPD data were corrected and Fourier Transformed using PDFgetX3 software56 to yield PDFs with Qmax = 21.4 Å-1. Neutron powder diffraction patterns57 were acquired at the D4c diffractometer58 at the Institut LaueLangevin (ILL) with a monochromatized wavelength of 0.4990 Å. Approximately 1.5 g of powders were prepared in cylindrical vanadium cans of 6 mm of diameter sealed with Helicoflex in an inert Ar atmosphere. The data were collected at room temperature in the 1.7-135.0º 2θº range with 3 h of data collection time per sample. The data were corrected for background scattering, vanadium can and self-absorption, and normalized with the Correct software59. The corrected data was Fourier Transformed using the STOG software embedded in the RMCProfile suite 60 using Qmin = 0.4 Å-1, Qmax = 24 Å-1. The combined refinements of the neutron and x-ray G(r) and S(Q) data sets were carried out with the RMCProfile software60. To account for the finite size of the simulation box, S(Q) was convoluted with a box function. Note that although it is desirable to include the Bragg data in the refinement to further constrain the long-range average structure, this was not possible in the present study due to the current version of the software lacking the profile shape of constant wavelength neutrons and anisotropic size broadening. Due to the impossibility to perform a multiphase refinement with this piece of software, the contribution of a minor crystalline Li2O impurity phase was removed from the diffraction patterns as follows: the diffraction signals from the impurity phase were first well described by Rietveld refinements performed with the Fullprof software61 and then the simulated patterns using the refined parameters were subtracted from the full diffraction patterns (see Figure S9). For the interpretation of the RMC outputted models, a quantitative analysis of the refined atomic positions inside the large box was carried out with a specifically written code using MatLab62. Li+ diffusion in Li4. BVSE pathway models (also known as energy-scaled bond-valence mismatch landscapes or energy landscapes)34 were derived from the RMC refined static structure model for the elucidation of the Li migration pathways in Li4. BVSE maps were constructed by summing the BVSE contributions to a hypothetical Li+ ion at each point with a 7·10-4 Å3 volume up to a distance of 8 Å. SoftBV parameters24 were used to account for the influence of higher coordination shells (i.e. to avoid abrupt changes in the BVSE by cut-off effects at the boundaries of coordination shells), which could lead to the formation of artefacts in the BVSE maps.



Additionally, the available volume fraction of a percolating pathway with a 3 eV threshold, and the percolation thresholds energy values along each axes were calculated from Monte-Carlo simulations using the BondStr software embedded in Fullprof-Suite61. The figures illustrating the local environments and diffusion pathways in this publication were produced with the VESTA software63.

Acknowledgements The authors acknowledge Dr Olivier Proux and Dr Jean-Louis Hazemann for their assistance during the experiment at FAME (CRG-ESRF, France). The refinements introduced in this publication were performed with the RMC-Profile software modified by Dr Wojciech Slawinski to allow for the simultaneous refinement of two S(Q) data sets, and his contribution is gratefully acknowledged. This work was supported by the ANR grant ANR15-CE05-0006-01 DAME.

Author contribution M.D., P.B. and C.V.C. were the experimentalists in the neutron and X-ray diffraction and XANES experiments, M.D. and P.B. performed the RMC refinement and model interpretation, M.D. and Y.J. performed the XANES simulations, M.F. and V.P. contributed to the synthesis of the material, H.F. lead the experiment at D4 and contributed to the data processing and correction for total scattering analysis, N.B. and N.B. designed the experimental set-up at D2AM, M.D and N.Boudet contributed to the data processing of X-ray data. M.D, C.V.C and P.B. contributed to writing the paper. P.B. and V.P. conceived and designed the project.

Competing interest The authors declare no competing financial interests.

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TOC graphic:

Supplementary information: Figure S1: Energy tresholds and voltage composition profile Figure S2: Quadrupole and dipole transitions Figure S3: Mn K-edge of standards Figure S4: Quantitative analysis of a single refinement Figure S5: Ellipsoidal fits of distorted octahedra Figure S6: XANES simulation of initial and refined models Figure S7: BVSE maps of several lithium transition metal oxides Figure S8: XANES simulation of Mn2O3 Figure S9: Data correction 28 ACS Paragon Plus Environment

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