Metal-to-metal charge transfer driven by electropositive species for stable anionic redox

Apart from conventional redox chemistries, exploring high-voltage redox processes, such as high-valent redox of a transition metal (TM) or novel anionic redox, is challenging due to unstable energy states correlated with destructive structural disorder in intercalation-type cathode materials. Here, we show a new strategy to design high-energy-density 4d-based Li-excess oxides with strong structural reversibility by substituting electropositive species. It is known that TMs in 4d-based Li-excess oxides, in contrast to 3d-based oxides, actively interact with neighboring oxygen ligands, inducing b 1 /b 1 * energy levels. Metal-to-metal charge transfer, driven by covalency competition within the asymmetric TM 3d -O-TM 4d backbone, induces a larger overlap between electronegative species and oxygen ligands, leading to stable high-voltage redox. Furthermore, we reveal that redox-inactive dopants control the reversibility of cation migration, thereby extending the battery lifetime. These insights open new perspectives for the control of intrinsic redox chemistry and enable rational designs for high-energy-density Li-excess cathodes.

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Abstract
Apart from conventional redox chemistries, exploring high-voltage redox processes, such as high-valent redox of a transition metal (TM) or novel anionic redox, is challenging due to unstable energy states correlated with destructive structural disorder in intercalation-type cathode materials.Here, we show a new strategy to design high-energy-density 4d-based Li-excess oxides with strong structural reversibility by substituting electropositive species.It is known that TMs in 4d-based Li-excess oxides, in contrast to 3d-based oxides, actively interact with neighboring oxygen ligands, inducing b 1 /b 1 * energy levels.Metalto-metal charge transfer, driven by covalency competition within the asymmetric TM 3d -O-TM 4d backbone, induces a larger overlap between electronegative species and oxygen ligands, leading to stable highvoltage redox.Furthermore, we reveal that redox-inactive dopants control the reversibility of cation migration, thereby extending the battery lifetime.These insights open new perspectives for the control of intrinsic redox chemistry and enable rational designs for high-energy-density Li-excess cathodes.

Full Text
With the sustained growth of the electric vehicle and electric power industries, the global energy storage market has entered a new phase where the innovative design of cathode materials is inevitably desired beyond commercial energy materials 1 .As the oxidation number of redox-active transition metals (TMs) limits the theoretical speci c capacity in a conventional charge compensation system [2][3][4][5][6] , the possibility of incorporating novel anionic redox chemistry into the conventional charge compensation system has been extensively explored.The cumulative cationic and anionic redox processes enable discharge capacities exceeding 250 mAh g −1 with a high upper cut-off potential of 4.6 V, delivering high-energy densities of up to 1,000 Wh kg −1 . 2,7However, such a high-voltage redox (hereafter denoted as HV-redox), commonly known as oxygen-redox (hereafter denoted as O-redox), triggers remarkable voltage hysteresis, voltage fading, and sluggish kinetics accompanied by irreversible structural transformations such as oxygen evolution and phase transition [8][9][10] .Previous studies have addressed these issues by (1) involving 4d-TM Ru and 5d-TM Ir to stabilize the oxygen state via stronger M-O bonding than that in 3d-TM Mn 11,12 and (2) doping inactive TMs to enhance the anionic redox reversibility [13][14][15] .Despite the positive outcomes, the interpretation of the origin of O-redox still remains challenging in these new types of cathode materials for high-energy Li-ion batteries.Typical Li-excess layered oxides (xLi 2 TMO 3 •(1-x)LiTMO 2 ) possess honeycomb-type atomic arrangements of Li 1/3 TM 2/3 , where four Li-ions and two TM ions coordinate oxygen to form an OTM 2 Li 4 octahedron (Fig. 1a).Considering the orbital orientation of TMt 2g and one of the O 2p orbitals along the direction where oxygen ions linearly bond to two Li-ions (Li-O-Li con guration), they do not hybridize with each other via sigma bonding, as shown in the inset in Fig. 1a, thus retaining nonbonding electrons around the oxygen anions 2 .From a quantitative molecular orbital perspective (Fig. 1b), the O 2p nonbonding (denoted as O NB ) energy states are located close to the original energy levels of oxygen and higher than the energy levels of bonding states from the hybridized bond.The general perception is that these O NB states enable anionic oxidation, O 2− /O − , prior to the depopulation of the M-O bonding state, causing a permanent occurrence of highly reactive oxygen radicals 16 .We previously de ned this type of oxidation as the oxidation of the orphaned oxygen O 2p state along the Li-O-Li con gurations 3 .First-principles calculations demonstrated that the hole density around the O anions resided along with the Li-O-Li local con guration after extracting labile electrons from O NB states in the highly charged state.However, this theoretical modeling should be discussed assuming that the O NB orbital does not overlap with adjacent TM orbitals such as Ni-O, Co-O, Mn-O, etc. Considering this, in the case of Ru (4d) and Ir (5d) with more spatially expanded orbitals than Mn (3d) 11,17 , the t 2g orbitals of TM are more likely to interact with O NB orbitals via π-type interactions 15 .Therefore, a deeper study of the "O-redox" process is necessary for 4d-and 5dbased Li-excess materials.The orbital hybridizations between the t 2g and O NB orbitals lead to energy splitting into the b 1 * antibonding state and b 1 bonding state; the splitting degree is affected by the covalence of the TM-O bond (Fig. 1c) 18 .Li-excess oxides with hybridization than one without hybridization facilitate the oxidation of hybridized t 2g and O NB orbitals, allowing further oxidation beyond conventional TM oxidation.However, oxidation of the O-dominant b 1 * state causes structural distortion followed by the formation of peroxo-like O-O species by unstable HV-redox 19,20 .Although highly covalent antibonding states, thereby enabling reversible HV-redox (Fig. 1f).We explore the intrinsic HV-redox chemistry in Li Electrochemical evaluation and phase transition in L(N)RO LRO and LNRO were synthesized by a conventional solid-state reaction 26 described in the Methods section.The structures were con rmed by synchrotron XRD combined with Rietveld re nement analysis of LRO and LNRO (Supplementary Fig. 1), con rming that both are well tted to the C2/c space group with a slight peak broadening due to defects, which are similar to already reported Ru-based Li-excess layered oxides 22,27 .
The effects of Ni substitution on the redox chemistry in Ru-based Li-excess oxides were investigated using galvanostatic electrochemical testing at a current density of 0.1 C (20 mA g -1 ) between 2.0 and 4.6 V. The voltage pro les exhibit remarkable gravimetric capacity beyond its theoretical capacity for both LRO and LNRO, as shown in Fig. 2a,d.Theoretically, it is well-known that the stable redox couple of Ru is Ru 4+/5+ , which corresponds to 164.5 mAh g −1 for LRO. 22Fig. 2a demonstrates that LRO delivers 299.6 mAh g −1 reversible capacity, which corresponds to 1.82 Li + per formula unit, indicating that Ru 4+ is overoxidized to Ru >5.5+ .The following cycle supplies 1.71 Li + , i.e., 94.4% of reversible Li + is supplied on the rst cycle.As differential electrochemical mass spectrometry (DEMS) analysis (bar chart in Fig. 2a,d) runs simultaneously to detect gas evolution, a signi cant amount of oxygen evolution was detected beyond Ru 5+ , attributed to lattice-oxygen dimerization.Meanwhile, in Fig. 2d, LNRO exhibits high reversible capacities (228.1 mAh g −1 , corresponding to 0.96 Li + ) beyond the theoretical capacities (226.7 mAh g −1 , corresponding to 0.95 Li + ), calculated from Ni 2+/4+ and Ru 4+/5+ redox (Fig. 2e).Moreover, LNRO exhibits 226.2 mAh g −1 reversible capacity (corresponding to 0.95 Li + ) on the second cycle with 99.2% capacity loss after the formation process.Noticeably, the DEMS result for LNRO contrasted with that for LRO, showing almost no oxygen evolution.The comparison of synchrotron XRD patterns also supports the results before and after the formation process (Supplementary Fig. 2).The superstructure peaks ranging from 19° to 34° represent the typical Li 1/3 TM 2/3 honeycomb-type ordering mentioned earlier 8,9,15 .
LRO shows overall peak broadening and dissipation, as estimated by a 36.8%peak area decrease, whereas LNRO maintains its original peak shapes with only a 10.8% peak area decrease.The results apparently show that Ni substitution effectively relieves the structural disorder.
In the same context, galvanostatic intermittent titration technique (GITT) analysis was performed to compare the voltage hysteresis of LRO and LNRO for the 2 nd cycle (Fig. 2b,e).GITT measurements show that the overpotential of each relaxation step gradually increases in the early stage of charging but rapidly increases during HV-redox.Relaxation behaviour at a high voltage shows that both samples are kinetically hindered by different parameters (insets in Fig. 2b,e).The relaxation step of LRO proceeds during the steady-state voltage change, whereas that of LNRO is largely contributed by the IR drop, which corresponds to the ohmic resistance and charge transfer.Recently, it was reported that the steady-state voltage change is affected by phase transition or local atomic rearrangement 28 .Despite similar voltage relaxation, our results show that HV-redox species would be structurally stabilized due to Ni substitution, consistent with signi cantly reduced voltage hysteresis as con rmed by equilibrium voltage curves (dashed curves in Fig. 2b,e).Fig. 2c,f shows the comparison of voltage pro les at 25 and 60 ℃, demonstrating that the kinetic hindrance during HV-redox is temperature-dependent, and that reduced delivery capacities are su ciently compensated for the addition of the CV step.

Redox mechanism during the rst cycle
To understand the underlying redox mechanism, we employed operando synchrotron transmission X-ray absorption near edge structure (XANES) at the Ru and Ni K-edges during the formation cycle, as shown in Fig. 3a-c.Ru K-edge XANES spectra consist of three representative peaks: i) pre-edge: 1s → hybridized 4d-5p transition, whose peak intensity stands out as the TMO 6 symmetry is more distorted, ii) whiteline (WL) edge: 1s → 5p transition, and iii) back peak: 1s → 5p transition by single scattering, following the dipole transition rules [29][30][31] .At the voltage increase to 3.9 V, as shown in Fig. 3a, the Ru WL energy in LRO shifts from 22140.0 to 22141.5 eV (Δ = 1.5 eV corresponding to 1.0 e − per f.u. of Ru), with the corresponding WL intensity decrease and back peak intensity increase, taking the typical shape of the Li 1 Ru 5+ O 3 plot 31 .However, as the cell potential remains beyond 4 V, where oxygen dimerization evolves, the pre-edge peak protrudes due to structural distortion, and the back peak weakens due to the loss of the O neighbor to scatter.The results consistently imply that the HV-redox of LRO involves progressive structural degradation with uctuating lattice oxygen.In contrast, LNRO shows different redox behavior, as con rmed by the differential crossline intensity (ΔI: showing no de nite propensity related to the structural distortion and oxygen loss in the vicinity of 22120 eV.Remarkably, the Ru WL energy in LNRO increases from 22139 eV, which is lower than that in LRO (22140 eV), indicating that Ru 4+ in LNRO holds a higher charge density than Ru 4+ in LRO.After charging, a peak shift of 2 eV (corresponding to 1.3 e − per f.u. of Ru) occurs in LNRO.In the meantime, a shift in the WL energy of Ni K-edge spectra from 8350 to 8353.6 eV proceeds until charging up to 4.6 V (Fig. 3c,f).
Considering that the Ni K-edge XANES spectra of LNRO are shifted to larger energy with respect to the Ni 2+ O reference 25 , Ni 2+ in pristine LNRO seems to have a lower charge density than that in NiO.As mentioned earlier, the results show that MMCT occurs from electropositive Ni to electronegative Ru via the O-bridge.The phenomena were also identi ed through computed average Bader charge, indicating that the more Ni is located near Ru, the higher charge value Ru has (Supplementary Fig. 10).Notably, Ni is oxidized up to approximately 3+ during charging as the redox potential of Ni increases with decreasing charge density [32][33][34] , in contrast to the perception that Ni can be oxidized to 4+ on typical NCM cathode materials 6,30 .The same is identi ed through rst-principles calculations in Supplementary Table 2, indicating that full oxidation (Ni 2+/4+ ) of partial Ni is thermodynamically hindered depending on the TM sites.Thus, we emphasize that the Ru 4+/5+ and Ru 5+/6+ redox reactions via MMCT are the main chargecompensating factors, whereas Ni species play a role in auxiliary electron donors for charge redistribution in LNRO.The following section investigates how such electrochemistry is coupled with structural changes.In the charged state, the FT intensity of the O peak in the LRO-Ru spectra, which is an indicator of atomic disordering (Debye-Waller factor, σ) and coordination number 31 , decreases abruptly with no complete recovery at the end of the formation cycle, which is direct evidence for irreversible oxygen evolution.In contrast, the FT intensity of the O peak in the LNRO-Ru spectra remains constant during the 1 st cycle, providing convincing proof that Ni aids the local structural stability of Ru, which is the main redox species.At the same time, Ni spectra exhibit highly uctuating but reversible spectral changes.Empirically, we determined that Ni ceaselessly migrates along its thermodynamically preferential routes, as veri ed through operando EXAFS results and calculations of relative site energies.The 2D contour map tracking migrated Ni species relative to xed Ru species shows that the T-peak at ~3.4 Å splits into two peaks, ~3.2 Å peak for the Li layer, and ~3.4 Å peak for the TM layer, as shown in Fig. 4a, right panel and Supplementary Fig. 13.Considering that Ru is xed at its original sites, such a peak split represents Ni migration toward its thermodynamically preferential neighboring site depending on the stoichiometry of delithiated LNRO.Additionally, the FT intensity of the O peak in the Ni spectra initially decreases, followed by a reciprocal tendency at the end of the charging process (marked by the white arrow in Fig. 4a), indicating that Ni undergoes out-of-plane migration from Oh TM to Td Li sites and further to Oh Li sites, simultaneously changing its coordination number (6 à 4 à 6) (Supplementary Fig. 14).The energy landscapes also provide consistent proof of whether Ni is the most stable at the Td Li sites or Oh Li sites, depending on the state of charge (SOC) (Supplementary Fig. 16).This is in agreement with a previous report that TM tends to trap the Li layer with a thermodynamic preference 8 .These results highlight the importance of the research addressing such an irreversible cationic migration, which will be discussed in the next section.

High
We compared the O peak in the EXAFS spectra depending on the cell voltage (Fig. 4b,c) to specify the oxygen stability during HV-redox.The O peaks in the LRO-Ru spectra can be distinguished into three regions corresponding to the short TM-O peak (~1.14 Å, denoted as S), main TM-O peak (~1.61 Å, denoted as M), and long TM-O peak (~2.05 Å, denoted as L).In particular, the interatomic distance of the S peak is similar to that of double bonds 15 .Two remarkable changes are observed in the EXAFS spectra of LRO upon charging: 1) as the cell voltage increases, the M peak generally contracts with L peak splitting, which is attributed to the distorted Ru octahedron due to the increased interaction between Ru and O originating from the higher oxidation states 31 .The extended TM-O bond weakens and ultimately gives rise to decoordinated oxygen 20 .
2) The decrease in S peak intensity and peak broadening also represent a disordered phase, which correlates with an unstable oxygen state upon charging.On the other hand, the O peaks of LNRO-Ru remain unchanged during delithiation, except for the S peak.The S peak gradually evolves while the cell is charged, indicating Ru-O bonds approaching double bonds through additional π-bonds.This terminal oxo species, which is double bound to Ru, can be stably maintained compared with single bound oxyl species 35,36 , nally leading to a highly reversible HV redox.
We performed rst-principles calculations to unveil the energetic effect of Ni on the redox system in the Ru-based Li-excess layered oxides (Fig. 4d,e).First, we visualized the electronic structures calculated on the basis of the Heyd-Scuseria-Ernzerhof (HSE06) hybrid functional 37  that in LRO (Supplementary Fig. 18).As expected from the XANES analysis, the MMCT process induced by Ni substitution enriches the charge density around Ru, making the redox system more reversible even at a high voltage, where the σ-bonding between p orbitals on neighboring oxygen ligands is inhibited by the competitive π-type interaction between Ru and O 38 .
Our suggestions regarding oxygen states coupled with the coordinated environment are summarized in Fig. 4f,g.According to crystal eld theory 39   and LNRO (89.2%) over 50 cycles at 0.5 C between 4.6 and 2.0 V versus Li 0 /Li + at 25 °C (Fig. 5b).To better understand the structural stability, we checked the average potential on each cycle.LRO with irreversible O-dominant HV-redox shows a remarkable voltage decay from 3.29 V for the second cycle to 2.96 V for the 50 th cycle (Δ = 0.33 V).LNRO shows a relatively stable average voltage drop of 3.48 V to 3.28 V (Δ = 0.20 V), manifesting less severe Ni-induced structural degradation caused by oxygen evolution.As expected, the LANRO exhibits superior stability with a slight voltage drop from 3.40 to 3.32 V (Δ = 0.08 V), proving that the harmony between Ni and Al brings outstanding redox stability in Ru-based Li-excess materials.These advantages are also observed through ex-situ synchrotron XRD of pristine, 4.6 V charged, 2.0 V discharged materials, and those after 50 cycles (Supplementary Fig. 22).All XRD patterns exhibit well-de ned representative peaks of the honeycomb-phase superstructure (shaded region).In particular, these peaks are still clearly seen even at 4.6 V charging and after 50 cycles, which means a well-maintained cationic arrangement with almost no migration within TM layers during HVredox and subsequent cycling.
Interestingly, HR-TEM data of each Li-excess oxide after cycling reveal how severely the cells are aggravated during cycling and how much Al doping enhances the structural integrity at the atomic level.(Fig. 5c and Supplementary Fig. 25) After prolonged cycles under the 0.5 C-rate, severe vacancies and microcracking were observed in LRO, as shown in Fig. 5c (left), which originate from severe latticeoxygen-evolution.In addition, atomic arrangements are not recognized on high magni cation (high-MAG) HR-images with random cationic disordering and numerous defects and vacancies.In contrast, low magni cation (low-MAG) HR-images of LNRO do not reveal microcracking, but numerous voids are still clearly seen, especially near the surface (Fig. 5c, middle).We attributed these voids to Ni dissolution into electrolytes, eventually toward the opposite side, which catalyzes electrolytic decomposition on the anode side and increases the cell resistance 43,44 .This phenomenon is clearly identi ed through TOF-SIMS mapping on anode surfaces after cycling, as shown in Supplementary Fig. 26.For the sake of convenience, we conducted full-cell tests with a graphite anode under identical conditions.Heavier Ni dissolution, compared with Ru dissolution, is observed in LNRO, where Ru dissolution might be attributed to the destructive impact of lattice-oxygen-evolution 45 and Ni dissolution might be attributed to its intrinsically unstable repetitive out-of-plane migration 46 .In particular, these migrated Ni atoms are presumably trapped into tetrahedral sites in the Li layer (see high-MAG TEM images in Fig. 5c, middle), in agreement with a previous report indicating that the Ni species are thermodynamically stable on the intermediate Td site of the O3 structure 8 .In stark contrast with both previous samples, the cycled LANRO shows almost awless low-MAG TEM images and well-de ned TM layers with slight cationic uctuations in the high-MAG TEM images in Fig. 5c (right).In addition, TOF-SIMS data also con rm that Al doping is notably effective in suppressing Ni dissolution (see Supplementary Fig. 26).
To reveal local structural changes in LANROs during cycling, we analyzed the migration behavior of each atom with 2D contour maps.In contrast to LRO and by analogy to LNRO (Fig. 4a-c), EXAFS at the Ru Kedge of LANRO demonstrates an entirely static propensity on the formation, the second cycle, and the 50 th cycle (Supplementary Fig. 27).However, a remarkable feature arises at the end of the charging where the Ru-TM 2 peak irreversibly splits into those corresponding to the TM layer (~3.6 Å) and Li layer (~3.2 Å), which is attributed to Al migration.Such irreversible peak splitting does not occur in the next cycle; instead, a new peak reversibly appears at a similar interatomic distance.Therefore, the irreversible out-oflayer migration of Al atoms on the formation step induces a reversible atomic rearrangement within the structural network after the 2 nd cycle, suggesting the positive interplay of Ni and Al in Ru-based materials.

A thermodynamic penalty of Ni migration by Al dopants
A previous report proposed that Al atoms located in the tetrahedral sites in Al-doped LiNiO 2 may mitigate the phase transformation upon increasing the temperature by disrupting Ni out-of-plane migration 42 .This statement is also experimentally veri ed in this work by the change in Ni coordination number during charging from (6→4→6) to (6→4) with Al doping, blocking the nal migration toward octahedral sites (Supplementary Fig. 27).
To clarify how Al thermodynamically hinders Ni migration, we conducted a computational study on the site preferences of Ni and Al (tetrahedral or octahedral sites) in Li slabs.Beforehand, we statistically examined the distribution of Ni ions neighboring Al among the 15 most stable LANRO structures based on Ewald summation followed by GGA+U relaxation (Supplementary Fig. 28).In 9 of 15 structures, all Ni ions (3Ni) are located on TM sites adjacent to Al within the second nearest neighboring Oh sites in the TM layer.Five of 15 structures hold two-thirds of Ni in the identical range.Statistically, Al is surrounded mainly by Ni in approximately 93% of modeled structures, i.e., most Ni migration can be controlled by Al ions within the crystal structure, in agreement with a previous paper reporting the same phenomenon through NMR spectroscopy 47 .The intriguing point is that Al migration to Td sites is thermodynamically and kinetically preferential for Ni migration nearby.The nudged elastic band (NEB) calculation in Fig. 5e shows that the migration barrier for Ni (1.54 eV) is higher than that for Al (1.16 eV), implying sluggish Ni migration around Al dopants.
Furthermore, it is noteworthy that Al Td signi cantly increases the site energy of both Td and Oh sites in Li slabs (Fig. 5f).Fundamentally, all Ni ions tend to migrate to adjacent Td sites on delithiated LNRO.Among them, the relative site energy of the Oh site near Ni (i) is estimated to be -0.92eV, which indicates further migration toward the Oh site.As the rest of the Ni ions also have a minor thermodynamic barrier for nearby Oh sites, their further migration is also feasible.In contrast to delithiated LNRO, however, Al prevents Ni migration toward Td sites, changing the site energy from negative to positive, except for Ni (i).
Although Ni (i) can move toward the Td site (-0.29 eV), Ni (i) migration is entirely suppressed in comparison with LNRO (-0.88 eV), and it is hard to migrate toward the Oh site (+0.10 eV) further.Thus, it is concluded that Al doping can support the structural integrity not only by physically preoccupying the empty sites toward which Ni tends to migrate but also by raising the relative thermodynamic free energy of neighboring sites.

Discussion
Resonant inelastic X-ray scattering (RIXS) has been considered one of the most powerful tools to understand Li-excess cathode materials due to its high resolution to detect oxygen oxidation.In this regard, it has been known that the O-redox of Li 1.2 Ni 0.2 Ru 0.6 O 2 is inactivated so far because the RIXS spectra do not show ngerprinting features of the O NB hole at 530.8 eV excitation and 523.7 eV emission energies, meaning that its delivery capacities might be compensated for both Ni 2+/4+ and Ru 4+/5+ . 27,48ecently, however, it was reported that Na 2−x Mn 3 O 7 does not show a characteristic RIXS signature despite its O-redox activity 49 .Based on the fact that Na Currently, RIXS analyses have been carried out in many pieces of research related to Li-excess materials, regardless of 3d, 4d, and 5d TM.Taking into account that 4d and 5d TM with stronger TM-O covalency induce additional hybridization of t 2g states and O NB states, as mentioned earlier, interpretations of RIXS results for "O-redox" should be differentiated, depending on whether the origin of HV-redox is the oxidation of O NB states or high-valent redox of TM.Considering that thermodynamic energies associated with unstable highly oxidized states cause O-O pairing induced by structural distortion, their new approach for sharp RIXS signatures at 530.8 eV excitation is su ciently acceptable.
We spectroscopically identi ed the coexistence of LMCT (O→Ru) and MLCT (Ni→O) in Fig. 3, even at a pristine state, supported by Bader charge analysis (Supplementary.Fig. 10), which is similar to the inductive effect 50 .This interplay of LMCT and MLCT, that is, MMCT (Ni◊Ru), shifts the energy states of Ru Operando X-ray absorption spectroscopy (XAS) analysis.Operando XAS data were collected on the BL10C beamline (WEXAFS) at the Pohang light source (PLS-II) with top-up mode operation under a ring current of 250 mA at 3.0 GeV.Monochromatic X-ray beams were obtained with a liquid-nitrogen-cooled double-crystal monochromator (Bruker ASC) with available in situ exchange in a vacuum between a Si (111) and Si (311) crystal pair.The Si (111) crystal pair was used for TM K-edge XAS measurements.For operando XAS measurements, pouch-type half-cells were assembled with a counter electrode of Li-metal foil in a dry room in transmittance mode using N 2 gas-lled ionization chambers (IC-SPEC, FMB Oxford) for the incident and transmitted X-ray photons.To eliminate higher-order harmonic contamination, we detuned the incident X-ray intensity by ~30%.With reference to TM foils placed in front of the third ion chamber, energy calibration processes were conducted for each XAS measurement.The XAS spectra for all TMs were obtained with a scanning time of 20-30 minutes for one spectrum.Using 0.2 C-rate of charging and discharging, 25-35 scans of XAS spectra for each TM were conducted to acquire reliable electronic/crystalline structures.Based on the standard XAS analysis procedure, the reduction of spectral data was performed into normalized XANES and Fourier transformed radial distribution functions (RDFs).
AUTOBK and FEFFIT modules in the UWXAFS package were used to obtain the k 3 -weighted TM K-edge EXAFS spectra, k 3 χ(k), with the removal of background and the normalization processes on the edge jump.The k 3 χ(k) spectra were Fourier transformed (FT) in the k range between 2.5 and 12.0 Å −1 .A difference of 0.3-0.5 Å between actual interatomic distances and observation values of radial distribution functions exists because phase correction was not conducted.
Ru and Ir can sustain peroxo-like O-O species via ligand-to-metal charge transfer (LMCT), the so-called reductive coupling mechanism, such structural distortion induces O-O bond shortening from approximately 2.5 Å in peroxo-like O-O species to 1.5 Å in O-O dimers, eventually resulting in the decoordination of dioxygen from TMs.Several studies have demonstrated that Sn-substituted materials (Li 2 Ru 1-y Sn y O 3 21 and Li 2 Ir 1-y Sn y O 3 15 ) serve as reversible redox hosts in the form of M Li -V M antisite-cationvacancies even at a high voltage, inducing stable 1.8 Å TM=O oxo species and short 1.4 Å O-O dimers, but the underlying redox chemistry remains elusive.Here, we propose a metal-to-metal charge transfer (MMCT) process driven by substituting species (TM 3d ), where internal charge transfer is activated from electropositive (TM 3d ) to electronegative species (TM 4d ), as shown in Fig. 1d.Such charge transfer induces a signi cant difference in TM-O covalency within the asymmetric TM 3d -O-TM 4d backbone, rectifying orbital overlap between TM 4d and O NB with a shifted redox potential of b 1 * (Fig. 1e).Independent of its oxidation number, denoted by the integer number according to the octet rule, TM with a large charge density affords metallic characteristics to the b 1 * -voltage redox stability coupled with local structural change To monitor the local structural change caused by HV-redox, we measured the operando extended X-ray absorption ne structure (EXAFS) at the Ru and Ni K-edges.Radial distribution functions of Fouriertransformed (FT) k 3 -weighted EXAFS provide local structure information on speci c TM species and interatomic distances, corresponding to TM-O bonding (~2.05 Å, denoted as O) and TM-TM bonding (~3.55 Å, denoted as T) 25 .As shown in Fig. 4a, EXAFS contour maps demonstrate a drastic contrast between LRO and LNRO, especially in the TM-O peak range corresponding to the rst coordination shell.
, enabling us to predict the partial density of states (pDOS) of the O 2p and Ru 4d states around the Fermi level (E F ) for LRO and LNRO.For both pristine LRO and LNRO, hybridized orbitals consisting of Ru 4d and O 2p located just below the Fermi level indicate the most preferential redox of Ru 4+/5+ at a low voltage (top of Fig. 4d,e).However, after all, Ru converts to the 5+ oxidation state, and the behaviour of the redox-contributing energy states of LRO and LNRO is quite different (bottom of Fig. 4d,e).The charge compensation of LRO at a relatively high voltage stems from the unstable Ru 5+/6+ , which is caused by O-dominant states.In stark contrast, pDOS of LNRO exhibits much more TM-dominant features near E F .The corresponding hybridized orbital states can be identi ed by visualizing the charge density for the energy range between 0 and -1.64 eV.The labile electrons of delithiated LNRO are contributed from TM-dominant b 1 * states, whereas those of delithiated LRO are from O-dominant b 1 * states (inset in the bottom of Fig. 4d,e).Computational analyses such as electron localization function (ELF) and integrated crystal orbital overlap population (ICOOP) also demonstrate a strong covalency between Ru 4d orbitals and O 2p orbitals in LNRO compared with

Fig. 5a compares
Fig. 5a compares the voltage pro les of the rst and 50 th cycles versus the normalized capacity of LRO, LNRO, and LANRO.The voltage pro le of LANRO is similar to that of LNRO, which delivers 286.9 mAh g −1 with Coulombic e ciency (C.E.) of 81.0 % on charging.The low rst-cycle C.E. might be attributed to the synergistic effect of the slightly increased O-dominant HV-redox due to the reduced Ni/Ru atomic ratio and irreversible Al migration.Meanwhile, the capacity retention and voltage decay are signi cantly enhanced starting from the second cycle.LANRO shows a superior stability of 97.5% over LRO (86.5%) 2−x Mn 3 O 7 does not form O-O peroxide (<1.5 Å), they proposed a new interpretation for the origin of RIXS signatures; such a sharp RIXS feature originates from the excitation of 1s → σ* in the short O-O peroxides, not the hole of O NB states.

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Figure 5 The
Figure 5 10,24,25.17Ru 0.61 O 2 (denoted as LNRO) in contrast to that in Li 2 RuO 3 (denoted as LRO) using We reveal that the reactive Ru-O part is stabilized via π back-bonding between O (Lewis base) and Ru (Lewis acid), preferentially over the formation of O-O peroxides, whereas Ni species partially reversibly migrate to tetrahedral sites, and then part of them further migrate toward octahedral sites in the Li layer.It is known that such migration often induces irreversible phase transformation from layered to spinel, which is a primary cause of voltage hysteresis and capacity fading10,24,25.We con rm how additional Al doping thermodynamically suppresses Ni migration through electrochemical tests, 23,22sensitive operando X-ray absorption spectroscopy and rst-principles calculations.The redox chemistry of LRO consists of a stable TM-dominant Ru 4+/5+ state and an unstable O-dominant Ru 5+/6+ state11,22.We demonstrate that Ni substitution induces polarized electrons toward electronegative Ru away from electropositive Ni, such that Ru ions pull the partial electron density away from Ni ions, leading to a stable TM-dominant Ru 5+/6+ redox state.In particular, it is well known that bond breaking is much more likely to occur with weak covalent TM-O bonds under bias voltage23, leaving highly reactive dangling oxygen.
41,42severe distortion, followed by the formation of peroxo-like O-O species.On the other hand, the d states of Ru in LNRO hold higher charge densities by taking the negative charge from the Ni species via the MMCT process, where stronger Ru-O covalency induces larger splitting of b 1 /b 1 * states (black dotted line).This shifted energy state of b 1 * makes both redox couples (Ru 4+/5+ and Ru 5+/6+ ) TM-dominant (blue line in Fig.4g).Additionally, Ru-O bonding can be stabilized by pulling the O ligands close to Ru accompanied by the spontaneous breaking of the Ni-O bond.Here, we conclude that the most critical chemistry inherent in stable HV-redox is the stabilization of high valent redox couples35along with terminal oxo species through double bonding before RuO 6 distortion, which leads to the formation of O-O peroxide.Improving structural reversibility during extended cyclingThus far, we have revealed how Ni species affect the redox chemistry of Ru-based Li-excess layered oxides.However, various Ni-related issues have already been considered the primary degradation mechanism, such as Ni trapping and irreversible phase transitions 8,10,40 .To prevent such irreversible Ni migration, we propose the aluminum doping of stoichiometric Li 1.22 Al 0.05 Ni 0.12 Ru 0.61 O 2 (denoted as LANRO), one of the most effective methods to suppress structural issues in NCM and NCA cathode materials41,42.
, the orbital energy state of Ru splits into upper e g * states and lower t 2g states.In particular, partially lled t 2g states, which are affected by O 2p neighbors via π-type interactions, split into b 1 * antibonding states and b 1 bonding states.As mentioned earlier, the b 1 * states of Ru in LRO in Fig. 4f consist of a highly reversible redox of Ru 4+/5+ (blue line) and unstable redox of Ru 5+/6+ (red line).When the energy states reach unstable b 1 * states of Ru 5+/6+ during HV-redox, the crystal structure suffers 5+/6+ toward higher energy levels, enabling reversible TM-dominant HV-redox.In particular, accompanying out-of-plane Ni migration can be enough to mitigate the energetic penalty associated with local distortion such that short O-O pairing is suppressed, as veri ed through reduced voltage hysteresis in experimental GITT results.That is why Li 1.2 Ni 0.2 Ru 0.6 O 2 does not show RIXS features for conventional O-redox, in contrast to Li 2 RuO 3 and Li 1.2 Ni 0.2 Mn 0.6 O 2 (see Supplementary Fig. 29).Furthermore, we also reveal that large voltage hysteresis is not entirely due to cationic migration but due to local structural distortion of the main redox species because Li 1.22 Ni 0.17 Ru 0.61 O 2 exhibits in nitesimal voltage hysteresis, despite its ceaseless Ni migration during cycling, which is similar to Li 1.17 Ti 0.33 Fe 0.5 O 2 28 .Such behaviour is in accordance with the ultimate design goal for Li-excess materials.Thus, our discovery provides not only new perspectives on a general understanding of the nature of anionic redox but also guidance for the design of cathode materials.Conclusion Our work is based on the following hypothesis that more spatially expanded TM, such as Ru and Ir, can strongly hybridize the neighboring O via π-type interactions, consequently leading to stronger covalence of TM-O bonds followed by additional splitting of t 2g and O NB energy levels into b 1 and b 1 * states.In particular, we believe that a higher charge density near Ru via MMCT induces more signi cant splitting of b 1 /b 1 *, making the b 1 * state more reversible during HV-redox.This proves that electropositive species play critical roles in stabilizing the HV-redox of electronegative species coupled with double-bonded O ligands.Therefore, we propose that strategies to improve the performance of Li-excess cathode materials should be differentiated based on the energy structure of their intrinsic molecular orbitals, namely, whether HV-redox utilizes the O NB or b 1 * energy level.Such observations guide us to nd a structural stabilizer to control cationic migration, such as redox-inactive Al dopants, and provide new insights into the design strategy to tune the relative energy states and control irreversible cationic migration for stable anionic redox.Methods Sample preparation.LRO, LNRO, and LANRO were prepared by using stoichiometric amounts of Li(CH 3 COO), RuO 2 , Ni(CH 3 COO) 2 •4H 2 O, and Al(NO) 3 •9H2O as precursors using a planetary ball mill with a small amount of acetone.A 5 wt% excess of Li(CH 3 COO) was used to compensate for lithium evaporation during high-temperature synthesis.The mixed powder was heat-treated at 650 ℃ for 6 h and at 1000 ℃ for 12 h at a ramp rate of 2 ℃/min in ambient atmosphere, followed by an intermediate grinding process and further heating at 900 ℃ for 12 h.Subsequently, the furnace was cooled to room temperature.Electrochemical characterizations.All electrochemical evaluations in this work were performed by using 2032 coin-type cells under the same conditions on each material.Electrodes were prepared by mixing 80 wt% active material, 10 wt% carbon black, and 10 wt% polytetra uoroethylene binder dissolved in Nmethyl-2-pyrrolidone.The obtained slurry was cast using a doctor blade and then dried under vacuum at 120 ℃ for 2 h.Electrochemical cells were assembled in an argon-lled glove box.The electrolytes were used with a composition of 1 M LiPF6 in ethylene carbonate (EC)/ethyl methyl carbonate (EMC)/dimethyl carbonate (DMC) in a volume ratio of 3:4:3.Electrochemical evaluations were carried out between 2.0 and 4.6 V at 0.1 C-rate during the formation process, followed by cycling at 0.5 C-rate at room temperature and high temperature (60 ℃).At room temperature, constant voltage steps (CV step) at 4.6 V charged state was added to every cycle until the current decreased to 0.05 C-rate.The galvanostatic intermittent titration technique was conducted under constant current pulses (C/10-rate) for 1 h, followed by relaxation for 4 h.