# Origin of voltage decay in high-capacity layered oxide electrodes

M. Sathiya, Artem M. Abakumov, Dominique Foix, G. Rousse, G Vantendeloo, Danielle Gonbeau, C. P. Laisa, Marie-Liesse Doublet, K. Ramesha, A. Prakash, et al.

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M. Sathiya, Artem M. Abakumov, Dominique Foix, G. Rousse, G Vantendeloo, et al. "Origin of voltage decay in high-capacity layered oxide electrodes." Nature Materials, 2015, pp.1-9. 10.1038/nmat4137. hal-01091051

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Origin of voltage decay in high-capacity layered oxide electrodes

**Reference:**

Sathiya M., Abakumov Artem M., Foix D., van Tendeloo Gustaaf, et al..- Origin of voltage decay in high-capacity layered

oxide electrodes

Nature materials - ISSN 1476-1122 - (2014), p. 1-9

DOI: http://dx.doi.org/doi:10.1038/nmat4137

Handle: http://hdl.handle.net/10067/1210020151162165141
# Origin of voltage decay in high-capacity layered oxide electrodes

M. Sathiya<sup>1,2,3</sup>, A. M. Abakumov<sup>4</sup>, D. Foix<sup>3,5,6</sup>, G. Rousse<sup>1,6,7</sup>, K. Ramesha<sup>8</sup>, D. Gonbeau<sup>3,5,6</sup>, M. Saubanère<sup>3,6,9</sup>, M. L. Doublet<sup>3,6,9</sup>, H. Vezin<sup>10</sup>, C. P. Laisa<sup>8</sup>, A. S. Prakash<sup>1,8</sup>, G. VanTendeloo<sup>4</sup> and J-M. Tarascon<sup>1,3,6*</sup>

Although Li-rich layered oxides (Li<sub>1+x</sub>Ni<sub>y</sub>Co<sub>2</sub>Mn<sub>1-x-y</sub>O<sub>2r</sub> > 250 mAh g<sup>-1</sup>) are attractive electrode materials providing energy densities more than 15% higher than today's commercial Li-ion cells, they suffer from voltage decay on cycling. To elucidate the origin of this phenomenon, we employ chemical substitution in structurally related Li<sub>2</sub>RuO<sub>3</sub> compounds. Li-rich layered Li<sub>2</sub>Ru<sub>1-y</sub>Ti<sub>y</sub>O<sub>3</sub> phases with capacities of ~240 mAh g<sup>-1</sup> exhibit the characteristic voltage decay on cycling. A combination of transmission electron microscopy and X-ray photoelectron spectroscopy studies reveals that the migration of cations between metal layers and Li layers is a congenital feature of the charge-discharge process that increases the trapping of metal ions in interstitial tetrahedral sites. A correlation between these trapped ions and the voltage decay is established by expanding the study to both Li<sub>2</sub>Ru<sub>1-x</sub>Sn<sub>x</sub>O<sub>3</sub> and Li<sub>2</sub>RuO<sub>3</sub>; the slowest decay occurs for the cations with the largest ionic radii. This effect is robust, and the finding provides insights into new chemistry to be explored for developing high-capacity layered electrodes that evade voltage decay.

Lithium-ion battery technology, having the highest gravimetric and volumetric energy densities of commercialized batteries, has conquered the portable electronics market<sup>1,2</sup>. Moreover, it is the battery of choice for powering electric vehicles and has great potential for grid energy storage<sup>3</sup>. Success in these fields will depend on our ability to further increase their energy density<sup>4</sup> to ensure longer cruising autonomy for electric vehicle and lower cost (per kilowatt-hour) to compete with pumped hydroelectric storage for grid applications. To reach such targets, current research mainly focuses on the synthesis of new Fe-based poly-anionic compounds<sup>5,6</sup> and improving the layered oxides used at present<sup>7</sup>, the latter being the topic of this paper.

Layered oxides of general formula LiNi<sub>1/3</sub>Co<sub>1/3</sub>Mn<sub>1/3</sub>O<sub>2</sub> (NMC) exhibit the highest capacity (≈200 mAh g<sup>-1</sup>) of any positive-electrode materials used at present<sup>8</sup>. The crystal structure of these materials can be considered as an ordered rock-salt derivative, where octahedrally coordinated Li and transition-metal cations ideally form alternating layers confined to the (111) planes of cubic oxygen close packing<sup>9</sup>. By modifying the chemical composition through substituting M with Li in the metal layers, Thackeray's and Dahn's groups have fabricated materials, termed as Li-rich layered oxides or Li-rich NMC (Li<sub>1+x</sub>Ni<sub>y</sub>Co<sub>x</sub>Mn<sub>1-x-y-z</sub>O<sub>2</sub>), exhibiting reversible capacities exceeding 250 mAh g<sup>-1</sup> (refs 10,11). The extra capacity was recently shown by complementary X-ray photoelectron spectroscopy (XPS) and electron paramagnetic resonance (EPR) experiments to be nested in the contribution of reversible anionic (O<sup>2-</sup> → O<sub>2</sub><sup>2-</sup>) redox processes to the overall material capacity<sup>12,13</sup>. Nevertheless, it remains that these materials have poor electrode kinetics and undergo large voltage decays (that is, change in slope of the characteristic voltage-composition curves) on cycling, both

of which are plaguing their implementation in practical cells<sup>14-16</sup>. Present scenarios, although not fully conclusive, tend to relate this voltage decay to structural effects and more specifically to the formation of spinel-like domains<sup>17-23</sup>. We recently found that the replacement of Mn<sup>4+</sup> in Li<sub>2</sub>Ru<sub>1-y</sub>Mn<sub>y</sub>O<sub>3</sub> by the larger and less electropositive cation Sn<sup>4+</sup> reduces voltage fading<sup>1,24</sup>. To assess the importance of substituent's chemical-physical characteristics on the voltage fade on cycling, we herein extend our work to the study of Li<sub>2</sub>Ru<sub>1-y</sub>Ti<sub>y</sub>O<sub>3</sub>. The Ti<sup>4+</sup> (d<sup>0</sup>) substitution was selected owing to its zero crystal field splitting, similar to d<sup>10</sup> Sn<sup>4+</sup>, its smaller size (0.60 Å) and its presumed ability to show accelerated cation migration<sup>25,26</sup> for direct visualization of migration paths.

Members of the Li<sub>2</sub>Ru<sub>1-y</sub>Ti<sub>y</sub>O<sub>3</sub> series were prepared by thoroughly mixing stoichiometric amounts of RuO<sub>2</sub>, TiO<sub>2</sub> and Li<sub>2</sub>CO<sub>3</sub> and annealing at 800 °C for 24 h. The resulting single-phase samples form a complete solid solution, demonstrated by the continuous variation of lattice parameters as a function of Ti content (Supplementary Fig. 1). The crystal structure is monoclinic with an atomic distribution corresponding to the Li<sub>2</sub>RuO<sub>3</sub> model with some Ru<sup>4+</sup> (ionic size 0.62 Å) being substituted by Ti<sup>4+</sup> (ionic size 0.60 Å).

For concision, our study focuses mainly on Li<sub>2</sub>Ru<sub>0.75</sub>Ti<sub>0.25</sub>O<sub>3</sub>. A combination of Rietveld refinement from powder X-ray diffraction (XRD) data (Fig. 1a), electron diffraction (Supplementary Fig. 2) and high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) imaging (Supplementary Fig. 3) reveals that its structure can be viewed (Fig. 1a, inset) as stacking of Li<sub>3</sub> layers and LiM<sub>2</sub> layers (M = Ru, Ti) along the *c* axis. In Li<sub>2</sub>Ru<sub>0.75</sub>Ti<sub>0.25</sub>O<sub>3</sub>, the Li and (Ru, Ti) cations demonstrate perfect honeycomb-like ordering within the LiM<sub>2</sub> layers; no cation inter-

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<footnote n="1">¹FRE 3677 'Chimie du Solide et de l'Energie', Collège de France, 11, Place Marcelin Berthelot, 75231 Paris, France, ²LRCS, CNRS UMR 7314, Université de Picardie Jules Verne, 80039 Amiens, France, ³ALISTORE-European Research Institute, 80039 Amiens, France, ⁴EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium, ⁵IPREM/ECP (UMR 5254), University of Pau, 2 av. Pierre Angot, 64053 Pau Cedex 9, France, ⁶Réseau sur le Stockage Electrochimique de l'Energie (RS2E), FR CNRS 3459, France, ⁷Sorbonne Universités UPMC Univ Paris 06, 4 place Jussieu, 75252 Paris Cedex 05, France, ⁸CSIR-CECRI- Chennai Centre, CSIR Campus, Taramani, Chennai-600 113, India, ⁹Institut Charles Gerhardt, CNRS UMR5253, Université Montpellier 2, 34 095 Montpellier, France, ¹⁰Univ. Lille Nord de France, CNRS, UMR 8516 - LASIR, Univ. Lille 1, F-59655 Villeneuve d'Ascq, France.</footnote>

<footnote>*e-mail: jean-marie.tarascon@u-picardie.fr</footnote>
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**Figure 1** | Structural and electrochemical aspects of Li₂Ru₀.₇₅ Ti₀.₂₅O₃. a, Powder XRD pattern (black points) and the Rietveld fit (red lines) for Li₂Ru₀.₇₅Ti₀.₂₅O₃; refinement was carried out using the monoclinic C2/m space group (lattice parameters: *a*=5.037(7) Å, *b*=8.717(1) Å, *c*=5.126(6) Å and *β*=109.137(3)°). Bragg positions are indicated as green vertical tick marks, and the blue line is the difference between the observed and calculated patterns. *R*<sub>Bragg</sub> = 7.88%. The peaks marked by stars are the reflections originating from the difference in scattering density between octahedral positions in the Li₃ and LiM₂ layers (interlayer ordering) and can serve as a mean to weigh the interlayer mixing, that is, the amount of transition-metal atoms (Ru/Ti) moving from the octahedral sites of the LiM₂ layers (denoted OcLIM₂) to the octahedral sites of the Li layers (denoted OcLi₃). The inset shows the structure of Li₂Ru₀.₇₅Ti₀.₂₅O₃ with stacking of LiM₂ (*M*=Ru, Ti) and Li₃ layers; Li is yellow, Ru/Ti is blue, possible tetrahedral sites for the occupation of metal ions are shown by green (in the Li₃ and LiM₂ layer) tetrahedra. b, Voltage versus composition profile for Li₂Ru₀.₇₅Ti₀.₂₅O₃ versus lithium as deduced from galvanostatic cycling at a 0.2 C rate. c, Percentage capacity retention plot of Li₂Ru₀.₇₅Ti₀.₂₅O₃ in comparison with that of Li₂RuO₃ and Li₂Ru₀.₇₅Sn₀.₂₅O₃; the as-measured capacities are presented in the inset. d, From cycle 1 to 100, the galvanostatic discharge profile of Li₂Ru₀.₇₅Ti₀.₂₅O₃ versus Li in the 2–4.6 V potential region; every tenth cycle is shown and the capacity values were normalized such that the maximum capacity in each cycle is taken as unity (240 and 151 mAh g⁻¹ is taken as unity for the first and hundredth cycles respectively). The inset shows the derivative plot of the same and the evolution of the new redox process at less than 3 V is marked by blue arrow.

mixing between the Li₃ and LiM₂ layers was observed. Thus, two types of cation ordering coexist in the material. The interlayer ordering between the Li₃ and LiM₂ layer drives the structure from the cubic rock salt *Fm*3*m* aristotype towards an *R*3*m* sublattice, whereas the honeycomb intralayer ordering decreases the symmetry even further to monoclinic C2/m. Stacking faults are therefore abundant (Supplementary Fig. 3) because of lateral shifts and/or 60° rotations of the LiM₂ layers.

The voltage–composition trace for a Li₂Ru₁₋₃Ti₃O₃ electrode (y = 0.25) shows the feasibility to remove 1.8 Li via two plateaux, of which 1.5 Li can be reinserted via an S-type discharge curve (Fig. 1b). This leads to a reversible capacity of nearly 240 mAh g⁻¹ that decreases with increasing y and reaches 110 mAh g⁻¹ for y = 0.75 (Supplementary Fig. 4).

A common feature of all members of the Li₂Ru₁₋₃Ti₃O₃ solid solution series is the large first-cycle irreversibility that increases by increasing the cutoff potential (Supplementary Fig. 5). Such irreversibility may occur partly owing to electrolyte oxidation but is mainly caused by irreversible loss of oxygen from the material²⁴,²⁷,²⁸ as confirmed by *in situ* gas and pressure analysis, with the onset of an O₂ loss at 4.1 V that then increases continuously as the charging

voltage proceeds to 4.5 V (not shown here). Last, similar to all Li-rich NMC or Li₂Ru₁₋₃M₃O₃ (*M*=Mn, Sn) phases, the first charge voltage trace differs from the first discharge trace, suggesting the contribution of a reversible anionic (O²⁻ → O₂²⁻) redox process to the overall capacity of Li₂Ru₁₋₃Ti₃O₃ samples. However, a minor difference from the Mn- and Sn-based layered electrodes¹²,²⁴ is the amplitude of the first-charge 3.7 V plateau for Li₂Ru₁₋₃Ti₃O₃, which no longer equals the initial amount of Ru. This suggests another redox process in competition with the Ru⁴⁺ → Ru⁵⁺ process, as is discussed later.

Long cycling tests performed on Li/Li₂Ru₀.₇₅Ti₀.₂₅O₃ cells (Fig. 1c) were compared to those of Li/Li₂Ru₀.₇₅Sn₀.₂₅O₃ and un-substituted Li₂RuO₃ cells cycled under various cycling rates between 2 V and 4.6 V. All samples exhibit initial capacities exceeding 200 mAh g⁻¹ (inset of Fig. 1c) with the Sn- and Ti-based samples showing the highest and lowest capacity retention, respectively (90% versus 62% after 100 cycles). This trend is preserved irrespective of the cycling rate (data not shown here).

In addition, the normalized discharge voltage traces for a Li/Li₂Ru₀.₇₅Ti₀.₂₅O₃ cell show that the voltage decays by nearly 500 mV between the first and hundredth cycle (Fig. 1d). This voltage

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<figure><img src="image_8.png" /><figcaption>**Figure 2** | Cycling behaviour of Li₂Ru(1-y)M_yO₃ (M = Ru/Sn/Ti) in the view towards voltage decay. **a–c**, Normalized capacity discharge profiles for Li₂RuO₃ (**a**), Li₂Ru₀.₇₅Sn₀.₂₅O₃ (**b**), Li₂Ru₀.₅Sn₀.₅O₃ and Li₂Ru₀.₂₅Sn₀.₇₅O₃ (**c**) from 1 to 100 cycles; the data are reported only for every 10 cycle. The capacity retention for Li₂Ru₀.₅Sn₀.₅O₃ and Li₂Ru₀.₂₅Sn₀.₇₅O₃ samples is shown as an inset in **c**. The capacity normalization was performed by taking the maximum discharge capacity in the corresponding cycle as unity.</figcaption></figure>

decay occurs mainly over the low potential region as confirmed by derivative curves (Fig. 1d inset), which show a progressive dumping of high to low voltage peak. When comparing the different M-based (M = Ti, Sn, Ru) samples, the voltage decay on cycling mirrors the capacity decay and is the smallest (~150 mV after 100 cycles) for Li₂Ru₀.₇₅Sn₀.₂₅O₃ (Fig. 2b). Such a trend persists whatever the amount of substituent (y; Supplementary Fig. 6). Note that the voltage decay is gradually reduced with increasing amounts of Sn in Li₂Ru₁⁻₃Sn₃O₃ (Fig. 2a–c) with almost no voltage fading observed after 100 cycles and nearly 100% capacity retention for y = 0.75 (Fig. 2c inset). This clearly confirms the positive role played by larger Sn⁴⁺ ions in reducing the voltage decay over many cycles.

Cycling tests were also done as a function of charge cutoff voltage (Supplementary Figs 7 and 8). When cycling is limited to the Ru⁴⁺/Ru⁵⁺ redox couple (<4 V) 100% capacity retention is observed for all systems, whereas the associated voltage decays vary from nearly zero for Sn to 80 mV for Ti. This suggests that capacity and voltage decay are not systematically related, as it may appear by raising the charge cutoff voltage to 4.2 V and 4.6 V where we note that both the capacity and voltage decay decrease in a similar manner on cycling.

In situ XRD patterns were collected at different states of charge/discharge for the Li/Li₂Ru₀.₇₅Ti₀.₂₅O₃ cell cycled at C/20 rate (Fig. 3a). As the cell is being charged to 4 V, an anisotropic variation of the lattice parameters is observed with c/a ratio decreasing from 4.99 to 4.92 suggesting a gradual suppression of Li₃–LiM₂ interlayer

ordering. When further charging the cell to 4.6 V, many Bragg peaks broaden and experience changes in relative intensities, indicative of a possible migration of Ru/Ti from LiM₂ to Li₃ layers. On subsequent discharge, the structure nearly converts back to the pristine layered ordered structure with however an increasing cation disorder within the LiM₂ layers as deduced from the reduction in intensity of superstructure peaks arising from the honeycomb LiM₂ pattern.

To assess structural changes on prolonged cycling, high-resolution synchrotron powder diffraction patterns were collected for the pristine Li₂Ru₀.₇₅Ti₀.₂₅O₃ and for electrodes recovered in the discharged state after 1 and 50 cycles, respectively (Fig. 3b). After 50 cycles, the XRD pattern significantly differs from that of the first cycle by an intense peak broadening and a progressive disappearance of the superstructure peaks. Owing to such broadening, the patterns were refined in the R̃3m subcell using an X* tan θ dependence accounting for strain effects; we observed a sharp increase of the strain broadening parameter X with increasing cycle numbers (Fig. 3c). Interestingly, the XRD pattern of the Li₂Ru₀.₇₅Sn₀.₂₅O₃ electrode cycled in similar conditions (C/10 rate, cycle numbers 1, 20, 50 and 100) shows a constant X parameter (Fig. 3d), suggesting a correlation between the voltage decay and micro-strains induced by local disorder (cation migration and/or stacking fault).

XPS analyses were carried out on pristine Li₂Ru₀.₇₅Ti₀.₂₅O₃ (Fig. 4a) and on similar electrodes charged to ~3.9 V (Fig. 4b) and 4.6 V (Fig. 4c), and charge discharged to 2 V (Fig. 4d). The

<figure><img src="image_9.png" /></figure>
<figure><img src="image_10.png" /><figcaption>**Figure 3** | Li-driven structural changes of Li₂Ru₀.₇₅M₀.₂₅O₃ (M = Sn/Ti). **a**, *In situ* XRD pattern of Li₂Ru₀.₇₅Ti₀.₂₅O₃ versus Li cells; the patterns collected on pristine electrode, during charge at 4 V, 4.6 V and then on completely discharged sample (2 V). On charging the Li/Li₂-xRu₀.₇₅Ti₀.₂₅O₃ cell to 4 V, there is a shift in the (003)R peak towards higher 2θ and merging of (006)R and (102)R Bragg peaks (indexed in R3m) indicating an anisotropic variation of lattice parameters with the removal of lithium. On discharge, most of the peaks regained their intensity and reappearance of (006)R and (102)R peak splitting shows that the structure converts back to the pristine layered ordered structure. **b**, Synchrotron powder XRD pattern of pristine Li₂Ru₀.₇₅Ti₀.₂₅O₃ and samples collected at the discharged state (2 V) after 1 cycle and 50 cycles. **c**, Variation of the Lorentzian broadening parameter X* tan θ with cycling derived from refinement of the synchrotron XRD patterns. **d**, Synchrotron powder XRD pattern of pristine Li₂Ru₀.₇₅Sn₀.₂₅O₃ and samples cycled several times (1, 25, 50 and 100 cycles) between 4.6 and 2 V and collected at 2 V. The green region in **b** and **c** shows the reflections from honeycomb Li-M ordering within the LiM₂ layers.</figcaption></figure>

evolution of Ru 3d<sub>5/2</sub> and O 1s core spectra along charge/discharge is very similar to what was previously observed in Sn-based electrodes<sup>22</sup>. The shift to higher binding energy of the Ru 3d<sub>5/2</sub> core peak on moving from OCV to 4 V identifies the Ru<sup>4+</sup> → Ru<sup>5+</sup> oxidation process<sup>29</sup> (Fig. 4, left), while the appearance of the 530.4-eV O 1s component at higher voltages is characteristic of the formation of peroxo-like O<sub>2</sub><sup>2-</sup> species<sup>30</sup>. A peculiarity here is the appearance of an O<sub>2</sub><sup>2-</sup> peak in the 4 V sample before complete oxidation of Ru<sup>4+</sup>–Ru<sup>5+</sup>, for reasons to be discussed later (Supplementary Fig. 9). Repeated discharge–charge (Supplementary Fig. 10) demonstrates the reversibility of these Ru<sup>4+</sup> ⇌ Ru<sup>5+</sup> and O<sup>2-</sup>/O<sub>2</sub><sup>2-</sup> processes, the latter being also confirmed by parallel EPR measurements (Supplementary Fig. 11) as previously described for the Li<sub>2</sub>Ru<sub>1-y</sub>Sn<sub>y</sub>O<sub>3</sub> phases<sup>12</sup>.

The true novelty here comes from the evolution of Ti 2p<sub>3/2</sub> core spectra. In the pristine electrode, the Ti 2p<sub>3/2</sub> component observed at 458.0 eV and the binding energy difference (O 1s–Ti 2p<sub>3/2</sub>) of 71.5 eV are characteristic of Ti<sup>4+</sup> ions in an octahedral environment<sup>31</sup>. For samples charged to 3.9 V a small shoulder appears at 459.1 eV, which becomes prominent at 4.6 V. Such a high binding energy (that is, decrease in electron density) is characteristic of Ti<sup>4+</sup> in a tetrahedral environment as reported earlier<sup>32</sup> and can be explained by weaker Ti 3d–O 2p overlap in the tetrahedral environment. The amount of tetrahedral Ti<sup>4+</sup>, as deduced from integration of the spectra, is ~15% and ~30%

for the 3.9 V and 4.6 V samples, respectively. This result is consistent with the relative ease of Ti<sup>4+</sup> ions migrating from octahedral to tetrahedral coordination due to the absence of ligand field stabilization<sup>25,33</sup>. At this point, it is tempting to suspect a correlation between the formation of O<sub>2</sub><sup>2-</sup> entities and the Ti<sup>4+</sup> migration, but such a correlation cannot be ascertained owing to the large contamination by surface oxide species that render the spectral deconvolution speculative given the small amount of O<sub>2</sub><sup>2-</sup>. After discharge, there is appearance of an additional component at 457.2 eV which can be attributed to Ti<sup>3+</sup> ions (reduction of Ti<sup>4+</sup>). From quantitative analysis we could deduce that the amount of Ti<sup>3+</sup> is about 35% and that octahedral Ti<sup>4+</sup> ions are mainly involved in the reduction process. The spectra obtained after 30 and 50 cycles clearly indicate an increase in the amount of tetrahedral titanium from ~25% after the first cycle to ~40% after 50 cycles. Those values are further refined by the investigation of Ti 3p core peaks (Supplementary Fig. 12), which offer less surface sensitivity, indicating a similar evolution over many cycles with slightly lower amounts of titanium in tetrahedral sites (1 cycle: ~20%, 50 cycles: ~34%). Unfortunately, a similar approach is not feasible with Ru 3d because of the less resolved core spectra associated with the Ru<sup>4+</sup> → Ru<sup>5+</sup> oxidation process. Last, the invariance of Sn 3d core spectra when charging/discharging Li₂Ru₁₋₃Sn₃O₃ implies the absence of Sn<sup>4+</sup> migration to tetrahedral interstices<sup>12</sup>.

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**Figure 4** | XPS of Li₂Ru₀.₇₅Ti₀.₂₅O₃. a-d, XPS spectra of the pristine Li₂Ru₀.₇₅Ti₀.₂₅O₃ sample (a), 3.9 V (b), sample charged to 4.6 V (c) and charge discharged to 2 V (d). C 1s-Ru 3d spectra showing the Ru 3d₅/₂ peak (blue region) together with a satellite peak (dashed blue region). The red and green arrows and lines are guides to the eye to show the shifting of the Ru 3d₅/₂ peak with charging and discharging. The O 1s core spectrum consists of two peaks corresponding to crystalline O²⁻ (529.5 eV) and surface oxidized species (531.6 eV). The new component appearing at 530.5 eV for the sample charged to 4.6 V is due to peroxo-like species associated with anion contribution for the redox process. Ti 2p₃/₂ spectra, as a function of the cell voltage, are shown in the second panel from the right to highlight the partial migration of Ti⁴⁺ from the octahedral site together with its partial reduction on discharge. The right panel shows the evolution of the Ti 2p₃/₂ spectra with increasing number of cycles (1, 30 and 50 cycles).

Overall, XPS results show: that the Li-uptake/removal process in Li₂Ru₁⁻ₓTi_yO₃ involves cumulative contributions of cationic (Ti⁴⁺ → Ti³⁺/Ru⁵⁺ → Ru⁴⁺) and anionic (O₂²⁻ → O²⁻) reversible redox processes; and a significant amount of Ti⁴⁺ migration from octahedral to tetrahedral interstices that increases on cycling (Fig. 4, right). To address the preferentially reduced Ti⁴⁺ sites, EPR was performed on the Li₂Ru₀.₇₅Sn₀.₂₅O₃ electrode, discharged to 2 V. The 4.2 K EPR spectrum (Supplementary Fig. 11) shows a set of two signals, with a weak *g* anisotropy (*g* = 1.92 and 1.89) for the first one and a lower *g* factor (1.79) for the second, which we ascribe to octahedral and tetrahedral Ti³⁺ centres, respectively³⁴,³⁵. This result confirms the Ti⁴⁺/Ti³⁺ reduction process detected by XPS, and provides evidence for the feasibility to partially reduce Ti⁴⁺ in both tetrahedral and octahedral sites, although quantification could not be done.

Another feature that cannot be deduced by XPS is the degree of cation migration between octahedral sites within the LiM₂ layers and the Li₃ layers, hereafter denoted as OcM and OcL₁, respectively, owing to nearly identical binding energies of both environments. The most natural technique to visualize cation migration at the local scale is HAADF-STEM. Thus, HAADF-STEM images were collected for the Li₂Ru₀.₇₅M₀.₂₅O₃ (M = Ti, Sn, Ru) electrodes cycled in similar conditions for various states of charge and discharge. Details on interpretation of the HAADF-STEM images are provided in Supplementary Figs 13 and 14.

We first studied the pristine Li₂Ru₀.₇₅M₀.₂₅O₃ material for which HAADF-STEM images revealed a two-dimensionally ordered structure with perfect Li and (Ru, Ti) ordering within the LiM₂ layers and no antisite defects between the Li₃ and LiM₂ layers (Supplementary Fig. 3). This contrasts with the overview [100] HAADF-STEM image (Fig. 5a) collected for a fully charged sample that shows the removal of Li to cause massive micro-structural changes, specifically: substantial disorder of the Li and (Ru, Ti) cations within the LiM₂ layers; migration of significant fractions of

(Ru, Ti) cations from the octahedral OcM positions in the LiM₂ layers to the octahedral positions OcL₁ in the Li₃ layers; and the presence of local areas with partial occupation of the tetrahedral interstices between the LiM₂ and Li₃ layers by the (Ru, Ti) cations. These structural changes are not homogeneous throughout the sample, as highlighted by the three typical HAADF-STEM images in Fig. 5b-d. In the portion of the sample probed in Fig. 5b, Li and (Ru, Ti) ordering within the LiM₂ layers is mainly preserved, but a fraction of (Ru, Ti) cations migrate to the tetrahedral interstices and a marginal amount of cations move to the OcL₁ positions. In the area shown in Fig. 5c, the occupation of OcL₁ positions is greater, with nevertheless some (Ru, Ti) cations remaining at tetrahedral interstices. Last, the image in Fig. 5d shows a part of the sample in which a large fraction of (Ru, Ti) cations have moved to the OcL₁ site and virtually no M cations remain in the tetrahedral sites.

Turning to the discharged Li₂Ru₀.₇₅Ti₀.₂₅O₃ sample (Supplementary Fig. 15), the HAADF-STEM pattern nearly converges back to that of a pristine sample, implying that most of the cations migrate back to their initial positions. Nevertheless, some remaining traces of cations in tetrahedral interstices are observable at the edge of the crystallite, and also at the distances of 15–25 nm towards the crystallite core, where the image resolution is not yet deteriorated markedly by increasing the crystal thickness (Supplementary Fig. 16).

Next, we collected HAADF-STEM images of a Li₂Ru₀.₇₅Ti₀.₂₅O₃ electrode cycled 50 times. The images significantly differ from those collected after one cycle, showing an intense disorder of Li and (Ru, Ti) cations within the LiM₂ layers and migration of (Ru/Ti) atoms away from the LiM₂ planes towards neighbouring interstitial positions (Fig. 6a,c, left). Fitting of the HAADF-STEM signal profile reveals a dominance of peaks corresponding to the migration of (Ru/Ti) cations towards tetrahedral sites, with peaks reminiscent of OcL₁ site occupancy being barely visible. Overall it clearly indicates, in agreement with XPS results, that the occupancy of interstitial
<figure><img src="image_12.png" /></figure>

**Figure 5** | HAADF-STEM images of $Li_{2-x}Ru_{0.75}Ti_{0.25}O_3$ electrodes. **a**, [100] overview HAADF-STEM image of the $Li_2Ru_{0.75}Ti_{0.25}O_3$ sample charged to 4.6 V. The insets show (from top to bottom): HAADF-STEM image of the atomic columns in the $LiM_2$ layer of the pristine $Li_2Ru_{0.75}Ti_{0.25}O_3$ sample along with the corresponding intensity profile demonstrating pair of dots separated by ~0.14 nm (this corresponds to perfect honeycomb ordering of the Li and M cations within the $LiM_2$ layers); HAADF-STEM image of the atomic columns in the $LiM_2$ layer of the charged $Li_2Ru_{0.75}Ti_{0.25}O_3$ sample along with the corresponding intensity profile demonstrating significant disorder of the M cations within the $LiM_2$ layers; selected-area electron diffraction pattern of the charged $Li_2Ru_{0.75}Ti_{0.25}O_3$ sample showing lines of diffuse modulated intensity at the $k \neq 3n$ positions due to local remaining order of the M cations in the $LiM_2$ layers. **b-d**, Images and corresponding HAADF-STEM signal profiles illustrating the different stages of the cation migration. **b**, The pristine structure is mostly preserved, most of the (Ru, Ti) cations are in the octahedral $Oc_{LiM_2}$ positions (blue peaks on the profiles), but a fraction of the (Ru, Ti) cations are moved to the tetrahedral interstices (green peaks on the profile marked with arrowheads). Occupation of the octahedral positions in the $Li_3$ layer remains marginal (yellow peaks on the profile marked with asterisks). **c**, Occupation of the octahedral positions in the $Li_3$ layer increases, some (Ru, Ti) cations remain at tetrahedral interstices. **d**, A large fraction of the (Ru, Ti) cations are moved to the octahedral position in the $Li_3$ layer. For simplification, in the aforementioned description, the $Li_3$ and $LiM_2$ layers of the pristine structure are termed as $Li_3$ and $LiM_2$ although the cation composition of these layers is changed substantially owing to cation migration.

tetrahedral sites increases with the number of cycles. A simple explanation is that the cations are progressively trapped within the tetrahedral sites as they migrate back and forth between $Oc_M$ and $Oc_{Li}$ sites, suggesting a possible correlation between the voltage decay and the trapping degree on cycling.

To test such a possible correlation we explored $Li_2Ru_{0.75}Sn_{0.25}O_3$ (Fig. 6) electrodes that had been cycled 100 times and showed the lowest voltage decay. HAADF-STEM images provide direct evidence for cation migration towards $Oc_{Li}$ sites in this system as well (Fig. 6b,c, right). More important is the quasi absence of cations in tetrahedral sites for this sample, further supporting our tentative correlation.

Finally, to clarify the outcome of Ru ions on cycling, we also studied $Li_2RuO_3$ electrodes cycled once, and after 100 cycles (Supplementary Figs 17 and 18). Comparing the HAADF-STEM

image of $Li_2RuO_3$ and $Li_2Ru_{0.75}Ti_{0.25}O_3$ electrodes after one charge-discharge cycle (Supplementary Figs 16 and 17) reveals virtually no population of tetrahedral sites in the pure Ru sample. This contrasts with the HAADF-STEM images collected after 100 cycles, which indicate the presence of Ru in tetrahedral sites. However, a comparison of the HAADF-STEM signal profiles of $Li_2Ru_{0.75}Ti_{0.25}O_3$ and $Li_2RuO_3$ samples cycled 50 and 100 times, respectively (Fig. 6c and Supplementary Fig. 18), demonstrates that even after twice as many cycles the population of tetrahedral interstices in $Li_2RuO_3$ does not reach the level of $Li_2Ru_{0.75}Ti_{0.25}O_3$. Taking into account the lower rate of voltage decay for $Li_2RuO_3$ compared with Ti-doped samples, this further supports the relationship between cation trapping at the tetrahedral interstices and voltage decay.

With respect to practical application, the migration of cations within the van der Waals gap (namely M in the $Li_3$ layer) has
<figure><img src="image_13.png" /><figcaption>**Figure 6** | Comparison of HAADF-STEM images of $Li_2Ru_{0.75}M_{0.25}O_3$ ($M = Sn/ Ti$) on long cycling. **a**, [11̄0] HAADF-STEM image of the $Li_2Ru_{0.75}Ti_{0.25}O_3$ sample after 50 charge-discharge cycles. Extra cations at the tetrahedral interstices next to the $LiM_2$ layers appear locally (some marked with arrowheads). **b**, [11̄0] HAADF-STEM image of the $Li_2Ru_{0.75}Sn_{0.25}O_3$ sample after 100 charge-discharge cycles. On both images the intense cation disorder is visible within the $LiM_2$ layers (appear brightest). Areas of the remaining order are also seen where pairs of bright dots alternate along the layers. **c**, HAADF-STEM signal profiles from areas outlined in the HAADF-STEM images. Green peaks associated with cation columns at the tetrahedral interstices are marked with arrowheads; yellow peaks of cation columns at the $Oc_{Li}$ positions are marked with asterisks; blue peaks represent the octahedral $Oc_{LiM_2}$ positions.</figcaption></figure>

been proved to be detrimental for cycling and rate capability performances³⁶,³⁷. Here we believe that such migration is not pernicious, because it is a cooperative migration probably initiated by the formation of the peroxo-like groups that creates free space thereby opening pathways for easy in- and out-of-plane cation migration. Although the performances of such disordered systems are surprising at first, it should be recalled that numerous redox systems such as conversion or displacement reactions enlist marked structural rearrangements, and even certain amorphous compounds can also work efficiently³⁸⁻⁴¹.

Together, our studies have shown $Li_2Ru_{1-y}Ti_yO_3$ to be a model system for exploring the complex interactions between cation

migration, structural rearrangements, creation of peroxo-like species and voltage decay, and reveal attractive features as discussed next in comparison with Sn-based or pure $Li_2RuO_3$ phases.

We found for $Li_2Ru_{1-y}Ti_yO_3$ phases an incomplete oxidation of $Ru^{4+}-Ru^{5+}$ during the first 3.6 V plateau suggesting a competitive anionic redox process that starts at a 200–300 mV lower voltage than for the similar $Li_2Ru_{0.75}Sn_{0.25}O_3$ electrode. Lowering of potential has to be correlated with the significantly smaller volume of $Li_{2-x}Ru_{1-y}Ti_yO_3$ compared with $Li_{2-x}Ru_{1-y}Sn_yO_3$, which increases the $Ru(4d)-O(2p)$ band overlap and therefore the covalent character of Ru–O bonds. Moreover, Ti and Sn substitutions, in spite of the fact that neither $Sn^{4+}$ nor $Ti^{4+}$ has free *d* electrons,

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Acknowledgements

The authors acknowledge both ALISTORE-ERI and RS2E institutions for fully supporting this work. G.V.T. acknowledges the European Research Council, ERC grant No. 246791-COUNTATOMS. The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483-ESTEEM2 (Integrated Infrastructure Initiative-13). C.P.L. thanks the CSIR, New Delhi for granting a Senior Research Fellowship. Use of the APS (Advanced Photon Source) at Argonne National Laboratory (ANL) supported by the US Department of Energy under Contract No. DE-AC02-06CH11357 is greatly acknowledged.

Author contributions

M.Sathiya, K.R., C.P.L. and A.S.P. carried out the synthesis, M.Sathiya, and J-M.T. conducted the electrochemical work and J-M.T. designed the research approach; G.R. analysed the crystal structures and diffraction patterns; A.M.A. and G.V.T. carried out, analysed and exploited the electron diffraction and HAADF-STEM studies; D.F. and D.G. collected and analysed the XPS spectra; M.Saubanère and M.L.D. conducted the DFT calculations and developed the theoretical framework; M.Sathiya, A.M.A. and J-M.T. wrote the manuscript and all authors discussed the experiments and final manuscript.

Additional information

Supplementary information is available in the online version of the paper. Reprints and permissions information is available online at www.nature.com/reprints. Correspondence and requests for materials should be addressed to J-M.T.

Competing financial interests

The authors declare no competing financial interests.