# Novel Insights into Enhanced Stability of Li-Rich Layered and High-Voltage Olivine Phosphate Cathodes for Advanced Batteries through Surface Modification and Electron Structure Design

Zhili Liang, Abdulaziz Baubaid, Mariusz Radtke, Maximilian Mellin, C.

Maheu, Sandipan Maiti, Hadar Sclar, Igor Píš, Silvia Nappini, Elena

Magnano, et al.

► To cite this version:

Zhili Liang, Abdulaziz Baubaid, Mariusz Radtke, Maximilian Mellin, C. Maheu, et al. Novel Insights into Enhanced Stability of Li-Rich Layered and High-Voltage Olivine Phosphate Cathodes for Advanced Batteries through Surface Modification and Electron Structure Design. Advanced Science, 2024, pp.2413054. 10.1002/advs.202413054. hal-04879868

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# Novel Insights into Enhanced Stability of Li-Rich Layered and High-Voltage Olivine Phosphate Cathodes for Advanced Batteries through Surface Modification and Electron Structure Design

Zhili Liang, Abdulaziz Baubaid, Mariusz Radtke, Maximilian Mellin, Clément Maheu, Sandipan Maiti, Hadar Sclar, Igor Píš, Silvia Nappini, Elena Magnano, Federica Bondino, Robert Winkler, René Hausbrand, Christian Hess, Lambert Alff, Boris Markovsky, Doron Aurbach, Wolfram Jaegermann, and Gennady Cherkashinin*

The design of cathode/electrolyte interfaces in high-energy density Li-ion batteries is critical to protect the surface against undesirable oxygen release from the cathodes when batteries are charged to high voltage. However, the involvement of the engineered interface in the cationic and anionic redox reactions associated with (de-)lithiation is often ignored, mostly due to the difficulty to separate these processes from chemical/catalytic reactions at the cathode/electrolyte interface. Here, a new electron energy band diagrams concept is developed that includes the examination of the electrochemical- and ionization- potentials evolution upon batteries cycling. The approach enables to forecast the intrinsic stability of the cathodes and discriminate the reaction pathways associated with interfacial electronic charge-transfer mechanisms. Specifically, light is shed on the evolution of cationic and anionic redox in high-energy density lithium-rich 0.33Li₂MnO₃·0.67LiNi₀.₄ Co₀.₂Mn₀.₄O₂ (HE-NCM) cathodes, particularly those that undergo surface modification through SO₂ and NH₃ double-gas treatment to suppress the structural degradation. The chemical composition and energy distribution of the occupied and unoccupied electronic states at the different charging/discharging states are quantitatively estimated by using advanced spectroscopy techniques, including operando Raman spectroscopy. The concept is successfully demonstrated in designing artificial interfaces for high-voltage olivine structure cathodes enabling stable battery operation up to 5.1 V versus Li⁺/Li.

## 1. Introduction

The limited energy density provided by cationic redox in the classical cathode materials for Li-ion batteries has targeted extensive research on triggering reversible anionic redox to achieve extra capacity.[1,2] In conventional lithium transition metal oxides (LMO) of layered structure, the O 2p states hybridized with antibonding M 3d states are situated further from the Fermi level ($E_F$) than M 3d states that requests a higher voltage (>4.2 V vs Li⁺/Li) to trigger the anion redox. An electron release from the O 2p states of the lattice is generally an irreversible process for the inverse $O^− → O^{2−}$ chemical reaction. In parallel to that, the partially oxidized $O^{2−}$ ions of the lattice oxygen can leave the solid in gaseous form, thereby determining the intrinsic voltage limit of the first-generation cathode materials.[3] The situation is different for Li-rich cathodes (LRC) where the structurally modified electronic configuration facilitates anion redox chemistry, which is believed to enable reversible oxygen redox.[4–6] A most known approach for reversible anionic redox is the

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/advs.202413054

[+]Present address: Nantes Université, CNRS, Institut des Matériaux de Nantes Jean Rouxel, IMN, F-44000 Nantes, France

© 2024 The Author(s). Advanced Science published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/advs.202413054

Z. Liang, A. Baubaid, M. Mellin, C. Maheu[+], R. Winkler, R. Hausbrand, L. Alff, W. Jaegermann, G. Cherkashinin Institute of Materials Science

Technische Universität Darmstadt
Peter-Grünberg-Str. 2, D-64287 Darmstadt, Germany

E-mail: gennady.cherkashinin@tu-darmstadt.de

M. Radtke, C. Hess

Department of Chemistry

Eduard-Zintl-Institut für Anorganische und Physikalische Chemie

Technische Universität Darmstadt

Peter-Grünberg-Str. 8, D-64287 Darmstadt, Germany
chemical synthesis of LRC via the integration of Li₂MnO₃ and LiMO₂ (M = transition metal) phases,[⁷] where an excess of lithium composition is thermodynamically favorable to substitute the transition metal in the MO₂ layers by Li-ions.[⁵,⁶] The oxygen redox is activated by delithiation of LRC at ≥ 4.5 V when the Fermi level crosses the top electronic states of Li₂MnO₃ which are mostly contributed from oxygen, whereas Mn 3d is situated at higher energy.[⁸] The accompanied electron removal from the O 2p level leads to the formation of Li₂O with its partial release from Li₂MnO₃ without the structural destabilization of the lattice inherent for conventional layered oxides.[⁹] According to density functional theory (DFT) calculations, a lithium excess in LRC leads to the formation of a non-bonding, non-hybridized O 2p orbital, which is structurally positioned within the Li-O-Li bond and electronically between the antibonding M 3d state and the O 2p bonding state of the lattice oxide.[⁶] Energetically, the non-hybridized O 2p orbital is more prone to donate an electron due to a closer energy to the E<sub>F</sub> in comparison with the O 2p orbital bonded to M 3d states. Moreover, triggering the anionic redox associated with the non-hybridized O 2p orbital does not lead to the structural destabilization of the lattice due to the localized character of this oxygen orbital.

Nevertheless, along with anomalous high capacity, LRCs are characterized by various structural changes such as layered-to-spinel transformation upon cycling followed by the formation of an electrochemically inactive phase with a rock-salt type disordered structure responsible for degradation, notable voltage hysteresis and long-term voltage fade.[¹⁰] Since the given phase transitions begin typically from the surface accompanied by oxygen release,[¹¹,¹²] several approaches have been developed to design the surface conditions suppressing structural degradation.[¹³] These include chemical treatment of LRC with various solutions and gases,[¹⁴⁻¹⁶] doping of the surface layers,[¹⁷] incorporation of an anion counterpart into oxygen site,[¹⁸] etc. Conceptually, the role of the modified surfaces is exclusively considered as a protective layer against the involvement of the cathodes into chemical reactions with the electrolytes and the oxygen release, or as Li-ion conducting layer facilitating Li⁺ transport across the electrode-electrolyte interfaces. However, the redox activity of the modified surfaces is often neglected thereby greatly underestimating the

importance of the bulk-surface-interface to electrolyte electronic charge transfer mechanisms affecting both an extra capacity and the internal stability of LRC.

Among the high-energy density cathode materials, Li-rich transition metal oxides with the chemical formula of xLi₂MnO₃(1-x)Li(M)O₂ (M = Mn, Ni, Co; x < 0.5) has drawn attention due to their high discharge capacity (>200 mAh g⁻¹) and environmental friendliness. 0.33Li₂MnO₃•0.67LiNi₀.₄Co₀.₂Mn₀.₄O₂ (named as HE-NCM) provides a discharge capacity >250 mAh g⁻¹. A higher capacity retention and rate capability, low voltage hysteresis and low impedance can be further achieved via the thermal surface treatment of HE-NCM with SO₂ and NH₃ gases studied in our previous work.[¹⁹] Such solid-gas reactions modify the surface of the cathode to a depth of several nanometers,[²⁰] mostly impacting the surface electronic configuration and chemical composition,[²¹] that might lead to more complex charge transfer mechanisms upon charging the HE-NCM in comparison with the untreated counterpart or with the bulk related redox processes.

In this contribution, we propose a novel concept based on the electron energy band diagrams approach to shed light on the cationic and anionic redox within the modified surface. We study the evolution of the electronic structure and interfacial chemical composition coupled to the examination of the electrochemical potential shift across the M 3d- and O 2p- occupied states caused by delithiation/lithiation of a series of HE-NCMs upon their cycling to ≥ 4.8 V (versus Li⁺/Li). For this purpose, we apply *quasi-in-situ* surface sensitive electron spectroscopy techniques, combined with theoretical calculations of X-ray absorption spectra and operando Raman spectroscopic experiments to probe the structure at the surface and in the bulk of the cathodes. Such a complementary techniques approach enables to explain the relevant electrochemical reactions, charge transfer mechanisms, and the nature of distortions at the oxygen site occurred at the modified surfaces and interfaces upon electrochemical cycling of various high-voltage cathodes for Li ion batteries.

## 2. Results and Discussion

HE-NCMs studied here are: a) *untreated HE-NCM*, b) *treated HE-NCM*, c) *untreated HE-NCM composite* = *untreated* cathode, d) *treated HE-NCM composite* = *treated* cathode (see the Experimental section for the details).

### 2.1. Oxidation States of the Cations and Anion at the Surface of As-Prepared Cathode Materials

Soft X-ray photoelectron spectroscopy (SPES) and X-ray absorption spectroscopy (XAS) probe depth of a few nanometres giving a unique opportunity to study the surface evolution and its involvement in the cationic- and anionic- redox processes. The O K-edge X-ray absorption near-edge structure (XANES) occurs due to an electron transition from the O 1s core-level to unoccupied O 2p states hybridized with empty M 3d electronic states in the 529–535 eV photon energy range.[²²] The wide band above 535 eV is associated with the O 1s to M 4s- and 4p- unoccupied states electron transition (Figure 1a). In the *untreated*

S. Maiti, H. Sclar
Department of Chemistry
Institute for Nanotechnology and Advanced Materials (BINA)
Bar-Ilan University
5290002 Ramat-Gan, Israel

I. Pfš, S. Nappini, E. Magnano, F. Bondino
CNR – Istituto Officina dei Materiali (IOM)
Strada Statale 14, km. 163,5 in Area Science Park Basovizza, 34149
Trieste, Italy

E. Magnano
Nanotechnology Research Laboratory
Faculty of Engineering
University of Sydney
2006 Camperdown, Australia

B. Markovsky, D. Aurbach
Department of Chemistry
BINA – BIU Institute of Nanotechnology and Advanced Materials
INIES – Israel National Institute of Energy Storage (supported by Israel
Ministry of Energy and Infrastructures)
Bar-Ilan University
5290002 Ramat-Gan, Israel
<figure><img src="image_3.png" /><figcaption>**Figure 1.** a) O K-edge XANES of HE-NCM prepared under different conditions: *untreated composite* (i) and *treated composite* (ii) obtained in the total electron yield (TEY) mode. The insert shows the zoomed area of O K XANES in the 528 – 537 eV range after the background subtraction and fitting the spectrum; ΔCFS is crystal field splitting. b–d) Mn L₃ (b), Co L₃ (c), and Ni L₃ (d) XANES of the *untreated composite* (i), *treated composite* (ii). e) The calculated Mn L₃ XANES of Mn²⁺ state in Oₕ symmetry (i), Mn³⁺ state in D₄ₕ symmetry (ii), Mn⁴⁺ state in Oₕ symmetry (iii); linear combination of Mn³⁺ (13%), Mn²⁺ (20%) and Mn⁴⁺ (67%) (iv). f) Calculated Co L₃ XANES with Co²⁺ (high-spin, HS) state (blue) and Co³⁺ (HS) (red) in Oₕ sites (i); Co²⁺ (HS) in the tetrahedral (T₉) (red) and Co³⁺ (low-spin, LS) in Oₕ (blue) symmetries (ii); linear combination of Co²⁺ HS in T₉ (44%) and Co³⁺ LS in Oₕ (56%) (iii). g) The calculated Ni L₃ XANES for Ni²⁺ in Oₕ symmetry. h, i) C 1s photoelectron spectra (hv = 1486.7 eV) (h) and Raman spectra (hv = 632 nm) (l) of as-prepared: *untreated HE-NCM composite* (i) and *treated HE-NCM composite* (ii). j) The most intense features are ascribed to C–C (A = 284.5 ± 0.1 eV), C–H (B = 285.0 ± 0.1 eV), C–O (C = 286.4 ± 0.2 eV), and C–F (from PVDF, D = 290.9 ± 0.2 eV). l) The vertical arrows show the expected band positions of LiOH, Li₂O₂ and Li₂CO₃. Note however, that these species are not detected in our Raman spectroscopic experiments. The vertical lines indicate vibrations associated with Li₂MnO₃ (A ≈ 488 cm⁻¹) and LiMO₂ (B = 491–497 cm⁻¹ and C = 518–614 cm⁻¹).</figcaption></figure>

HE-NCM, the oxidation and spin states of transition metals are Mn⁴⁺ (high-spin, HS), Co³⁺ (low-spin, LS) and Ni²⁺ (see the M 3d electronic configuration in details in Figure S1a-i,v,ix, Supporting Information).[21] Accordingly, the t₂g (lowest energy) and e₉ (highest energy) electronic levels of M 3d (labeled as A and B in the O K-edge, Figure 1a-i) are bound to the O 2p states in an octahedral (Oₕ) symmetry.[21] For the ground state electronic configuration, the t₂g orbitals of the Co³⁺ and Ni²⁺ states are fully occupied. Therefore, the A peak (529.6 eV) can be associated only with the Mn⁴⁺ (t₂g) empty state (Figure S1a-i, Supporting Information). The empty e₉ states of Mn⁴⁺, Co³⁺, and Ni²⁺ (partially) (Figure S1a-i,v,ix, Supporting Information) mostly contribute to the B peak at 531.8 eV. The energy difference between the A (t₂g) and B (e₉) peaks defined by crystal field splitting (ΔCFS), gives ΔCFS = 2.2 eV in a good agreement with ΔCFS ≈ 2.4 eV inherent for tetravalent metal ions bound to oxygen p state in Oₕ symmetry.[23] The relevant ground state electronic configuration (with almost negligible contribution from Mn³⁺ (Figure 1b-i) is also evidenced by M L₃,₂- edges XANES, which occurs as dipole allowed electron transitions from the M 2p core level to the M 3d

empty states (Figure 1b-i–d-i). The conclusion on oxidation states of M is in agreement with theoretical calculations of the Mn L₃, Co L₃ (Co³⁺, LS) and Ni L₃,₂ XANES (Figure 1e-iii,f-ii,g) and the experimental data obtained on a model NCM thin-film cathode (Figure S1b-ii–d-ii, Supporting Information).

The SO₂ and NH₃ double gas treatment modifies the electronic configuration at the surface of HE-NCM. The purpose of such a chemical treatment was discussed in our previous works.[15,21] A brief discussion on the surface modification with either SO₂ and NH₃ alone can be found in the Discussion Section S1 (Supporting Information). The double gas treatment of HE-NCM induces numerous modifications observed in the O K XANES (Figure 1a-ii), a decrease in the A and B peaks intensity while the peak C emerges and two minor features (D⁺ and B⁺) also appear. The A and B decrease evidences an increasing the M 3d occupation due to a reduction of transition metal ions that would have a trend in lowering the M 3d–O 2p hybridization and a smaller crystal field splitting for Oₕ symmetry.[23] However, the estimated ΔCFS = 2.8 eV is higher than expected value for both the Oₕ and tetrahedral (T₉) symmetries (it should be noted that ΔCFS for T₉ is
typically lower than that of an O<sub>h</sub> environment). A more reason-
able assumption is an additional empty state contributing to the
peak B. This “hidden” state is resolved via spectral decomposition
of the O K XANES followed by a weighted least-squares fitting of
the spectral region (Figure 1a-ii, insert). Accordingly, the empty
state B (532.0 eV) is assigned to the M 3d (e<sub>g</sub>) state giving Δ<sub>CFS</sub>
= 2.3 eV. The B<sup>+</sup> feature at higher energy (532.7 eV) is most prob-
ably associated with vacancies in the nitric oxide (NO) species
formed in a small amount due to chemical reactions of NH<sub>3</sub> with
the HE-NCM surface.[15] The empty O 2p state of N–O is typi-
cally manifested at 532.7 eV in O K XANES.[23] In spite of a weak
adsorption of NH<sub>3</sub> onto the HE-NCM surface,[21] our previous
research showed that ammonia modifies the electronic structure
of the cathodes significantly, acting as a reducing agent,[15] which
in turns leads to the formation of NO<sub>x</sub> (Figure S2-i, Supporting
Information). The shoulder D<sup>*</sup> in the O K XANES in the 537.5
– 538.5 eV range (Figure 1a-ii is ascribed to SO<sub>4</sub><sup>2-</sup> groups,[24,25]
in accordance with the O 1s (≈ 532 eV) and S 2p (169.4 eV) pho-
toelectron emissions of the *treated* HE- NCM composite (Figures
S3-iii and S4, Supporting Information).[21] A reduction of the oxi-
dation states of transition metals caused by SO<sub>2</sub> and NH<sub>3</sub> double
gas treatment is also detected by the Mn L- and Co L- XANES
with the clear visible superposition of Mn<sup>4+</sup> with Mn<sup>3+</sup>/Mn<sup>2+</sup>
and Co<sup>3+</sup>/Co<sup>2+</sup> oxidation states, whereas Ni remains mostly in 2+
state (Figure 1b-ii,c-ii,d-ii). The coordination of transition met-
als with oxygen at the surface of the *treated* HE-NCM was ex-
plored by theoretical calculations of the Mn L<sub>3</sub>- and Co L<sub>3</sub>- edges
XANES in various environments of the M 3d states with the O
2p level. These calculations show that the experimental Mn L<sub>3</sub>
and Co L<sub>3</sub> XANES are well described by a linear combination
of the O<sub>h</sub>- and D<sub>4h</sub>- symmetries and of the O<sub>h</sub>- and T<sub>d</sub>- symme-
tries, respectively (Figure 1e-iv,f-iii). Note, that T<sub>d</sub> symmetry is
often assigned to spinel structure, whereas D<sub>4h</sub> symmetry is a
reduction of O<sub>h</sub> symmetry due to Jahn-Teller effect. More inter-
esting is the peak C (533.7 eV) which emerges in O K XANES
*only* after the double gas treatment of HE-NCM (Figure 1a-ii).
Lithium carbonate, which has a similar energy, can be ruled
out, because the chemical composition of the carbon species is
very similar for both the *untreated*- and *treated*- HE-NCM compos-
ites (Figure 1h). The absence of Li<sub>2</sub>CO<sub>3</sub> in the bulk is also con-
firmed by Raman spectroscopy (Figure 1l–ii). Thus, O K XANES
at 533.7 eV is assigned to the surface Li<sub>2</sub>O. A more discussion
on other oxide species appears in the Discussion Section S2,
Supporting Information). In the as-prepared *untreated* HE- NCM
composite, Li<sub>2</sub>O associated with a Li–O–Li bond is not detected
neither by the surface sensitive techniques nor by operando Ra-
man spectroscopy which probes the subsurface/bulk of the sam-
ples (Figure 11a–i,1–i; Figure S5, Supporting Information).

## 2.2. Distribution of Electronic States Near the Fermi Level as a Predictor of Redox Mechanisms

Quantitative energy band diagrams are crucial to understand the
underlying mechanisms occurring during the double SO₂ and
NH₃ gas treatment. Moreover, understanding of the distribu-
tion of electronic states near E<sub>F</sub>, coupled to an ability to monitor
changes in the electrochemical potential of the cathodes (μ<sub>C</sub><sup>e</sup> =
E<sub>F</sub>) upon electrochemical cycling of batteries, is important to pre-

dict the involvement of the cations and the anion into redox reac-
tions. In our work, the energy distribution of the occupied elec-
tronic states near E<sub>F</sub> is derived from the results of resonant pho-
toemission spectroscopy (RPES). By collecting the valence band
photoelectron spectra while scanning the photon energies across
the M L- and O K- edges, it is possible to observe the resonances
of the M 3d- and O 2p- valence states. These resonances arise
from a quantum mechanical interference between the direct pho-
toemission process and the recombination of an excited state.[26]
Our RPES experiments evidence three states closest to E<sub>F</sub>. One
is associated with Co<sup>2+</sup> in the *treated* cathode (A<sup>1</sup><sub>T</sub> = 1.7 eV ±
0.1 eV), whereas Co<sup>3+</sup> at A<sup>2</sup><sub>T</sub> = 2.0 ± 0.1 eV and A<sub>U</sub> = 1.9 ±
0.1 eV is revealed for the *treated*- and *untreated*- cathodes, respec-
tively (Figure S6, Supporting Information). The Ni<sup>2+</sup> 3d level is
situated at higher energies, E<sub>bin</sub> = 2.3 – 2.5 eV for both cathodes
(Figure 2a,b). The presence of the Mn<sup>4+</sup>/Mn<sup>3+</sup>/Mn<sup>2+</sup> oxidation
states in the *treated* cathode leads to the several resonance fea-
tures of Mn 3d states (Figure 2c,d). These states are more distant
from E<sub>F</sub> as compared to the Ni 3d and Co 3d (Figure 2a,b; Figure
S6, Supporting Information) and have a higher binding energy
than Mn<sup>4+</sup> of the *untreated* cathode (Figure 2c,d). The more pro-
nounced resonances of the M 3d occupied states of the *untreated*
HE-NCM might indicate that the electrons of the transition met-
als are more localized in this cathode as compared to the *treated*
counterpart (the mapping of photon energies versus binding en-
ergies is shown in Figure S7, Supporting Information).

The O 2p resonance shows the general characteristics for both
cathodes in the range of 9 – 11 eV (Figure 2e,f). A slight exten-
sion of the O 2p state to higher binding energies in the case of
*treated HE-NCM composite* is associated with the O 2p – M 3d
bonding hybrid states. The main difference in the valence band
structure associated with the O 2p states of the two cathodes is the
resonance for the *treated* cathode (absent in the *untreated* cathode,
Figure 2f) with an onset at E<sub>bin</sub> ≈ 4.5 eV and the maximum at E<sub>bin</sub>
= 5.5 eV at photon energy of ≈ 534 eV (Figure 2e; Figure S8a,c,
Supporting Information). This occupied O 2p state is mostly as-
signed to Li<sub>2</sub>O (see above our discussion on the nature of the C
peak in the O K XANES) in agreement with the previously re-
ported E<sub>bin</sub> values for lithium oxide lying in the ≈ 4.5 – 10 eV
range.[27] The lack of the relevant O 2p resonance in the *untreated*
HE-NCM composite (Figure 2f; Figure S8b, Supporting Informa-
tion) makes the presence of Li<sub>2</sub>O unlikely in the absence of the
SO<sub>2</sub> and NH<sub>3</sub> surface treatment. LiOH is expected at a slightly
higher binding energies (7–10 eV,[28]) and is not detected neither
for the *treated*- or *untreated*- cathodes (Figure 2e,f).

Thus, our RPES results show that the binding energy sequence
of the M 3d (t<sub>2g</sub>) occupied states hybridized with the O 2p (latt-
tice) are in the following order: Mn > Ni > Co. Also, the O 2p
occupied state of Li<sub>2</sub>O of the *treated* HE-NCM is situated at lower
binding energy from the O 2p state of the lattice oxygen. The en-
ergy distribution of the unoccupied electronic states near E<sub>F</sub>, ob-
tained from the O K-, Co L-, Ni L-, and Mn L- edges XANES, is
plotted on a unified energy scale (Figure S9a,b, Supporting Infor-
mation). It is worth noting that the L-edges do not directly repre-
sent the density of states (DOS) due to a complex impact of the
final-state effects on the shape of L-edge.[29] Nevertheless, the rep-
resentation of M L XANES with respect to E<sub>F</sub> provides valuable
insights into the relative energy positions of unoccupied M 3d
states.[30,31]
<figure><img src="image_4.png" /><figcaption>**Figure 2.** Valence band (VB) photoemission spectra of HE-NCM *treated* - (a,c,e) and *untreated* - (b,d,f) composites collected at photon energies of the 849 eV – 856 eV (a, b), 637 eV – 646 eV (c, d), 527.9 eV (off-resonance of O K-edge XANES) (e, f) and the 533.4 – 534.9 eV (on-resonance for the C peak in O K-edge XANES of the *treated* cathode). The solid circles in the inserts show excitation energies. The VB spectra collected at the Ni L₃-edge (a, b) of the highest intensity (in red) show the Ni²⁺ 3d resonance at a) A<sub>T</sub> = 2.5 eV and b) A<sub>U</sub> = 2.35 eV for the *treated*- and *untreated*- HE-NCM, respectively. The VB spectra collected at the Mn L₃-edge of the *treated* HE-NCM (c) show the two resonance structures (highlighted in red) at ≈ 640 and ≈ 642.7 eV photon energies, where A = 3.6 ± 0.1 eV, B = 4.1 ± 0.05 eV, C = 5.3 ± 0.1 eV and D = 6.2 ± 0.2 eV are associated with the Mn 3d levels of Mn²⁺, Mn³⁺ and Mn⁴⁺ oxidation states. d) The resonances at A = 2.5 ± 0.2 eV, B = 3.4 ± 0.4 eV, C = 3.7 ± 0.2 eV and D = 4.4 ± 0.2 eV in the *untreated* HE-NCM are mostly with Mn⁴⁺ state. e) The resonance in the Δ = 4.5 – 7.0 eV binding energy range (labelled by vertical lines and horizontal double arrow) of the *treated* HE-NCM is mostly assigned to a Li₂O state. (e, f) The O (KVV) Auger transition contributed at E<sub>bin</sub> > 5 eV is shown by the dashed line.</figcaption></figure>

By neglecting the band bending of the electronic states, which
can impact the relative energy positions of the DOS,[³¹] one can
conclude that the occupied electronic states of the *treated* cath-
ode are shifted to higher binding energies by ≈ 0.2 eV as com-
pared to the *untreated* cathode, evidenced by the valence band
maxima (VBM) at 1.1 ± 0.1 and 0.85 ± 0.1 eV versus E<sub>F</sub> for the
*treated*- and *untreated*- HE-NCM composites, respectively (Figure
S10a-i,b-i, Supporting Information). The observed VBM shift in
the *treated* cathode reflects lifting up of E<sub>F</sub> toward the unoccu-
pied states and is not attributed to the chemical shift. The SO<sub>2</sub>
and NH<sub>3</sub> double gas treatment causes the reduction of Co<sup>3+</sup> →
Co<sup>2+</sup> and Mn<sup>4+</sup> → Mn<sup>2+/3+</sup>. In turn, a partial occupation of the
Mn 3d (e<sub>g</sub>)- and Co 3d (e<sub>g</sub>)- empty states (see the electronic con-
figuration in Figure S1a-i–vi, Supporting Information) leads to a
shift of E<sub>F</sub> away from the VBM, i.e., the photoelectron spectra are
shifted to higher binding energies as compared to the *untreated*

counterpart. For the *untreated* HE-NCM *composites*, the relevant
3d (e<sub>g</sub>) levels are empty (Figure S8a-i,v, Supporting Information)
that explains the *E*<sub>F</sub> position closer to the VBM.

The electron energy band diagrams of the occupied- and unoccupied- electronic states for the as-prepared *treated*- and *untreated*- HE-NCM composites derived from the current photoemission experiments are shown (Figure 3a,b).

2.3. Evolution of the Electrochemical Potential, Cation-, and Anion- Redox During Electrochemical Cycling of HE-NCM Cathodes

A driving force for the battery operation, that defines the open
circuit voltage of a battery, is the difference in the chemical poten-
tials of lithium (Δμ<sub>Li</sub>) in the cathode (μ<sub>C</sub><sup>Li</sup>) and in the anode (μ<sub>A</sub><sup>Li</sup>).
<figure><img src="image_5.png" /><figcaption>**Figure 3.** The energy band diagrams of as-prepared *treated*- (a) and *untreated*- (b) HE-NCM cathodes. The electronic states of Ni, Co, and Mn are presented in dark blue, red, and green, respectively. Li-O-Li bond with a vacancy at oxygen (*V*<sub>O</sub>) is also shown (a). The main difference in the energy diagram of the *treated* cathode is oxygen vacancies in Li<sub>2</sub>O and the presence of N-O bond. Δ<sub>Co3d</sub> = 1.85 eV ± 0.15 eV for *treated* cathode is an average value of the binding energy of the Co<sup>2+</sup> 3d (*E*<sub>bin</sub> = 1.7 eV) and Co<sup>3+</sup> 3d (*E*<sub>bin</sub> = 2.0 eV) states with respect to *E*<sub>F</sub> (see RPES in Figure S6a, Supporting Information for details). VBM<sup>T</sup>, VBM<sup>U</sup>, φ<sub>e</sub><sup>T</sup>, and φ<sub>e</sub><sup>U</sup>, are the valence band maximum and the work function of the *treated* and *untreated* cathodes, respectively. The VBM and work function were measured at hv = 1486.7 eV and hv = 47 eV, respectively. The estimated values of the ionization potential at these photon energies are IP<sup>T</sup> = 5.8 ± 0.2 eV and IP<sup>U</sup> = 5.3 ± 0.2 eV for the *treated*- and *untreated*- cathodes, respectively. The increase of work function and ionization potential after the SO<sub>2</sub> and NH<sub>3</sub> treatment is assigned to the formation of Li<sub>2</sub>O. The downward band bending (i.e., away from the *E*<sub>F</sub> in the bulk) of γ ≈ 0.4 eV was observed by probing the electronic structure with higher photon energies (Table 1). For the *treated* cathode, a more colored area near the empty states (as compared to the *untreated* cathode) reflects the *E*<sub>F</sub> shift toward the unoccupied states due to their partial occupation due to the Mn<sup>2+</sup>/3+ and Co<sup>2+</sup> states. The work function is different also due to the presence of Li<sub>2</sub>O and N-O species at the surface. c) Evolution of the work function, φ<sub>e</sub><sup>T</sup>, and valence band maximum (VBM<sup>T</sup>) upon charging the *treated* cathode. φ<sub>e</sub><sup>T</sup> and VBM<sup>T</sup> are labeled by red- and blue-color balls, respectively. The dashed line shows a linear extrapolation of φ<sub>e</sub><sup>T</sup> and VBM<sup>T</sup> to charging state of 5.1 V.</figcaption></figure>

Previous model considerations supported by our recent experi-
mental studies have demonstrated that the electronic part (μ<sub>C</sub><sup>e</sup>)
of the chemical potential,[32] plays the dominating role in the bat-
tery voltage as compared to the ionic part (μ<sub>C</sub><sup>Li+</sup>).[33–35] In terms of
the evolution of the electronic structure upon electrochemical cy-
cle, charging (discharging) will lead to removal (accommodation)
of a Li<sup>+</sup> ion from (in) the cathode that will immediately lead to re-
moval (accepting) of a valence electron from HE-NCM to satisfy
the charge-neutrality principle. Accordingly, the E<sub>F</sub> of the cathode
is increased upon charging and decreased upon discharging.[3]
In other words, the internal work function, φ<sub>e</sub>, is increased by
delithiation of the cathodes, whereas binding energy of the VBM
is decreased versus E<sub>F</sub>. The opposite direction takes place upon
lithiation of the cathodes. Note, this simplified model does not
take into account the surface effects, such as a double layer poten-
tial drop (i.e., dipole potential, δ<sub>dip</sub>) at the surface of the cathode
material occurred, for example, due to the cathode/electrolyte in-
terfacial (CEI) layer formation. Nevertheless, the E<sub>F</sub> energy shift,
coupled to the changes in occupation of the M 3d- and O 2p-
electronic states versus cycling, allows to determine the voltage
profiles upon cycling various intercalation battery systems.[34,36]
Moreover, by tracking the current E<sub>F</sub> position and knowing redox
levels of the electrolyte components (and impurities, for instance
water), the nature of various chemical reactions occurred at the
electrode/electrolyte interface, and which are not associated with
delithiation/lithiation of the electrodes, can be explained.[31,37] In
spite of the importance of band structure analysis for advanced
batteries, still only few activities are devoted to its practical ap-
plication for the battery field.[31,35,38] Here, we demonstrate the

experimental validity of the electrochemical potential evolution
upon delithiation (lithiation) of the HE-NCM cathodes (Figure 3c
and Table 1; Figure S10, Supporting Information).

2.3.1. Cationic Redox

Based on the energy band diagrams derived from our SPES,
XANES and RPES experiments (Figure 3a,b), as well as in agree-
ment with theoretical expectations, E<sub>F</sub> enters the occupied Co 3d-
and Ni 3d- states at first in the beginning of charging the HE-
NCM cathodes. In turn, this leads to the oxidation of Co and
Ni as supported by the Co L<sub>3</sub>- and Ni L<sub>3</sub>-XANES, whose spec-
tral features change in intensity (Figure 4a,b,d,e; Figures S13
and S14, Supporting Information), as well as by the Co K- and
Ni K- XANES which shift to higher energies by charging the
cathodes ≤ 4.5 V (Figures S15 and S16, Supporting Informa-
tion). Mn ions do not change their oxidation state by charging
the *treated- and untreated- HE-NCM* cathodes below 4.5 V as ev-
idenced by the Mn L XANES and Mn 3s photoelectron spectra
(Figures 4c,f and 5a-i-iii and Table 2; Figure S17, Supporting In-
formation). Such finding is in good agreement with the DOS dis-
tribution of the as-prepared HE-NCMs where the Mn 3d states
have a higher binding energy with respect to the Co 3d- and
Ni 3d- states (Figure 2a–d; Figure S6, Supporting Information).
No significant changes in the O K-edge XANES up to 3.7 V evi-
dence that lattice oxygen is not involved in the anionic redox pro-
cess (Figure S18, Supporting Information). The minor changes
in the O K XANES of the *treated HE-NCM* cathode are ascribed
to evolution of the oxygen-related species induced by the NH<sub>3</sub>
<table><caption>**Table 2.** The binding energy evolution of the Mn 3s photoemission and exchange satellite ($S_{exc}$) versus charging/discharging of the *treated* HE-NCM cathode.</caption><thead><tr><th></th><th>Mn 3s [eV]</th><th>S<sub>exc</sub> [eV]</th><th>ΔE<sub>Mn3s-Sexc</sub> [eV]</th><th>Valence</th></tr></thead><tbody><tr><td>Pristine</td><td>84.25 ± 0.05</td><td>89.15 ± 0.05</td><td>4.9 ± 0.1</td><td>≈ 3.5+</td></tr><tr><td>Soaking</td><td>83.85 ± 0.05</td><td>89.2 ± 0.1</td><td>5.35 ± 0.15</td><td>≈ 3+</td></tr><tr><td>3.7 V charged</td><td>84.0</td><td>89.0</td><td>5.0 ± 0.1</td><td>≈ 3.4</td></tr><tr><td>4.5 V charged</td><td>83.8</td><td>88.9 ± 0.1</td><td>5.1 ± 0.1</td><td>≈ 3.3+</td></tr><tr><td>4.6 V charged</td><td>83.7</td><td>89.5 ± 0.1</td><td>5.8 ± 0.1</td><td>≈ 2.3+</td></tr><tr><td>4.8 V charged</td><td>84.0</td><td>89.7 ± 0.2</td><td>5.7 ± 0.2</td><td>≈ 2.4+</td></tr><tr><td>4.4 V disch.</td><td>84.0 ± 0.1</td><td>88.8 ± 0.1</td><td>4.8 ± 0.2</td><td>≈ 3.6+</td></tr><tr><td>3.75 V disch.</td><td>84.05 ± 0.05</td><td>89.0 ± 0.1</td><td>4.95 ± 0.15</td><td>≈ 3.4+</td></tr><tr><td>After 1<sup>st</sup> cycle</td><td>84.05 ± 0.05</td><td>89.25 ± 0.05</td><td>5.2 ± 0.1</td><td>≈ 3.1</td></tr></tbody></table>

surface gas treatment (see discussion in more details in the Discussion Section S3, Supporting Information). Importantly, $E_F$ does not cross the O 2p state of Li₂O at a charging state below 4.5 V that is evidenced as the absence of the marked changes in the C peak intensity of the O K XANES (Figure S18, Supporting Information). In the voltage profiles of the HE-NCM cathodes, the plateau region labelled by 1 corresponds to oxidation of the relevant transition metals (Figure 6). A more detailed discussion on the electrochemical properties of the HE-NCM cathodes is provided in the Discussion Section S4 and Figure S11 (Supporting Information).

### 2.3.2. Activation of Anionic Redox

The $E_F$ approach to the O 2p state of the lattice oxygen hybridized with the M 3d states leads to distinct differences in the O K XANES (Figures S18a(iv),b(ii), and S19, Supporting Information). At the 4.5 V charging state, the $E_F$ position equals to 5.25 ± 0.1 eV (versus the vacuum level, $E_{vac}$) for both the

*untreated- and treated- HE-NCM* cathodes (Table 1). For the *treated* HE-NCM, this means that $E_F$ is equal to the energy of the non-bonding O 2p orbital of Li₂O, which lies in the ≈ 4.5 – ≈ 5.5 eV binding energy range (Figure 2a). As the result, the C peak (the empty states of a Li–O–Li bond) disappears irreversibly in the O K XANES (Figure 7a-iii), thereby evidencing oxidation of the Li₂O related non-bonding states and removal of lithium oxide from HE-NCM. In addition, the A* shoulder (528.8 eV) occurs in O K XANES associated with the lattice oxygen in both the *treated- and untreated- HE-NCM* cathodes (Figure 7a-iii,b-ii). Whilst there are different models of oxidation mechanisms in LRC, namely O²⁻/O⁻ redox, or O²⁻/(O₂)ⁿ⁻ (peroxo-like dimer species) or isolated O⁻ anions,[40] we assign the A* feature with the hole formation on oxygen (lattice) in agreement with the previous work.[31] Note that peroxide formation is not detected also by Raman spectroscopy (Figure S20, Supporting Information). Thus, the movement of $E_F$ into the O 2p state of the lattice oxygen leads to the O²⁻ oxidation, or in other words, to hole transfer to the O 2p state. In the voltage profiles, the plateau region associated with oxygen redox is labelled by 2 (Figure 6). An enhanced holes formation (A* = 528.8 eV) in the *untreated* HE-NCM evidences a stronger distortion of the oxygen site as compared to the *treated* HE-NCM (Figure 7a-iii,b-ii). A higher amount of oxygen holes in *untreated* HE-NCM is accompanied with Li₂O activation (irreversible) manifested by the C peak at 534.0 eV (Figure 7b-ii), which evidences to triggering oxidation of the lattice oxygen. In this regard, *untreated* HE-NCM behaves as a classical LRC, where charging the cathodes above 4.4 V leads to activation of Li₂MnO₃ via the formation of Li₂O (supported also by operando Raman spectroscopy, Figure S20, Supporting Information), followed by its decomposition to evolve oxygen gas.[9] In parallel to the involvement of the lattice oxygen in charge compensation, the transition metals are more oxidized in both cathodes. This is evidenced by changes in the Co L-, Co K-, Ni L-, and Ni K- XANES (Figure 4a-iii,b-iii,d-ii,e-iii; Figures S13–S16, Supporting Information), as well as in a relative intensity increase of the O K XANES peaks associated

<table><caption>**Table 1.** Evolution of the work function ($\phi_e$.T and $\phi_e$.U), valence band maximum (VBM) and ionization potential (IP.T and IP.U) of the *treated- and untreated- HE-NCM* cathodes.</caption><thead><tr><th></th><th>Φ<sub>e</sub>.T [eV](hv = 47 eV)</th><th>VBM [eV](hv = 250 eV)</th><th>IP.T [eV](VBM + Φ<sub>e</sub>.T)</th><th>VBM [eV](hv = 1486.7 eV)</th></tr></thead><tbody><tr><td><strong>Treated HE-NCM</strong></td><td></td><td></td><td></td><td></td></tr><tr><td>As-prepared</td><td>4.70 ± 0.1</td><td>0.75 ± 0.1</td><td>5.45 ± 0.2</td><td>1.1 ± 0.1</td></tr><tr><td>3.7 V charge</td><td>5.05 ± 0.1</td><td>0.60 eV ± 0.1</td><td>5.65 ± 0.2</td><td>0.8 ± 0.1</td></tr><tr><td>4.5 V charge</td><td>5.25 ± 0.1</td><td>0.50 ± 0.1</td><td>5.75 ± 0.2</td><td>0.4 ± 0.1</td></tr><tr><td>4.6 V charge</td><td>5.26 ± 0.1</td><td>0.22 ± 0.1</td><td>5.48 ± 0.2</td><td>0.3 ± 0.1</td></tr><tr><td>4.8 V charge</td><td>5.37 ± 0.1</td><td>0.18 ± 0.1</td><td>5.55 eV ± 0.2</td><td>0.4 ± 0.1</td></tr><tr><td>5.1 V charge</td><td>≈ 5.43<sup>a)</sup></td><td>≈ 0.15<sup>a)</sup></td><td>5.58</td><td>—</td></tr><tr><td><strong>Untreated HE-NCM</strong></td><td>Φ<sub>e</sub>.U (hv = 90 eV)</td><td>VBM (hv = 450 eV)</td><td>IP.U eV(VBM + Φ<sub>e</sub>.U)</td><td>VBM<sup>c)</sup> (hv = 1486.7 eV)</td></tr><tr><td>As-prepared</td><td>4.45 ± 0.1 eV<sup>b)</sup></td><td>0.72 ± 0.1 eV</td><td>5.17 ± 0.2 eV</td><td>0.85 ± 0.1 eV</td></tr><tr><td>4.5 V charge</td><td>5.20 ± 0.1 eV</td><td>0.55 ± 0.1 eV</td><td>5.75 ± 0.2 eV</td><td>0.50 ± 0.1 eV</td></tr><tr><td>4.6 V charge</td><td>5.25 ± 0.1 eV</td><td>0.35 ± 0.1 eV</td><td>5.60 ± 0.2 eV</td><td>—</td></tr><tr><td>4.8 V charge</td><td>5.77 ± 0.1</td><td>0.15 ± 0.1 eV</td><td>5.90 ± 0.1 eV</td><td>0.15 ± 0.1 eV</td></tr></tbody></table>

<footnote>a) The values obtained by a linear extrapolation of $\phi_e$.T and VBM to 5.1 V; b) $\phi_e$.U was measured at hv = 47 eV; c) VBM was measured on the *untreated* HE-NCM (without conductive carbon and PVDF). The work function is defined as $E_F$ with respect to the vacuum level, $E_{vac}$. The ionization potential, IP = $\phi_e$. + ($E_F$ – VBM) + $\Delta\delta_{dip}$, is defined as the energy between the highest occupied state (i.e., VBM) and $E_{vac}$. Some discrepancies in the equal change of the work function, $\phi_e$, and VBM are caused by a surface dipole layer, $\Delta\delta_{dip}$, which depends on the interfacial conditions (e.g., cathode electrolyte interface (CEI) layer formation). The downward band bending in the as-prepared *treated- and untreated-* cathodes is evidenced by a shift of VBM to higher binding energy (i.e., away from $E_F$) at higher photon energies, which probe the electronic structure deeper into the bulk.</footnote>
<figure><img src="image_6.png" /><figcaption>**Figure 4.** Evolution of Co L$_{3}$, Ni L$_{3,2}$ and Mn L$_{3}$ XANES of the *treated*- (a–c) and *untreated*- (d–f) HE-NCM cathodes versus charging (a–f) and discharging (a–c) during the first electrochemical cycle with the upper cut-off of 4.8 V. As-prepared are labelled by “as is” (i), the charging states are colored in red. The discharging states are colored in violet [a–c(vi–vii)], cycled [a–c(viii)]. a) Co L$_{3}$ XANES: a steadily intensity decrease of the low energy A* shoulder by charging the *treated* HE-NCM up to 4.5 V [a(i–iii)] evidences oxidation of Co$^{2+}$; the A*, E and D* features inherent to Co$^{2+}$ state increase in intensity upon charging the *treated* HE-NCM to 4.6 V [a(iv)]; a shift of the B to higher energies upon charging the cathodes [a(i–v), d] evidences an increase in the oxidation state of Co. The opposite B shift occurs upon discharging the cathodes [a(vi–viii)]. b, e) Ni L$_{3,2}$ XANES: the intensity increase of the B shoulder and its shift to higher energies are due to an increase in the oxidation state of Ni. Upon discharging, the B decreased in the intensity and shifted in the opposite direction, thereby confirming reduction to the Ni$^{2+}$ state in the *treated* HE-NCM. c,f) Mn L$_{3}$ XANES: the increase in intensity of the A (≈ 640.2 eV), B (640.8 eV), and D (641.8 eV) shoulders of the *treated* HE-NCM cathode [c(iv)] is due to the contribution of Mn$^{2+}$ and Mn$^{3+}$ (see calculations in Figure 1e); Mn L$_{3}$ XANES is not changed markedly by charging the *untreated* HE-NCM cathode (f).</figcaption></figure>

with the t$_{2}$$_{g}$ and e$_{g}$ empty states mostly of Co and Ni (Figure S19-iii, Supporting Information). Minor changes in the shape of the Mn L XANES, as well as the Mn 3s photoelectron spectra (Figure S17a-iii,b-ii, Supporting Information; Figure 5a-iv) support inactivity of Mn 3d in charge compensation at such voltage. Interestingly to note that a higher oxidation of the Ni ions in the *untreated* HE-NCM, as compared to the treated counterpart (see Ni L$_{3,2}$ XANES in Figure 4b-iii,e-ii), correlates with a higher distortion of the lattice oxygen in the *untreated* HE-NCM

charged to the same potential (4.5 V) (Figure 7a-iii,b-ii). As such, we can assume that E$_{F}$ penetrates deeper into the Ni 3d- and O 2p- occupied states of the *untreated* HE-NCM. This suggestion agrees with the relative energies of the electronic states of as-prepared cathodes, where E$_{F}$ is situated closer to the VBM of the *untreated* HE-NCM as illustrated by the energy diagrams (Figure 3a,b).

By charging the HE-NCM cathodes to a higher potential (4.6 V), the E$_{F}$ shifts deeper down (Figure 3c and Table 1).
<figure><img src="image_7.png" /><figcaption>**Figure 5.** a) Evolution of the Mn 3s photoelectron spectra (hν = 1486.7 eV) of the *treated* HE-NCM cathode versus charging and discharging during the first electrochemical cycle with the upper cut-off of 4.8 V. As-prepared sample (i), the charged states of 3.7 V (ii), 4.5 V (iii), 4.6 V (iv), 4.8 V (v), discharged states of 4.4 V (vi) and 3.75 V (vii), cycled (viii). ΔE is the exchange splitting defined as the energy difference between the Mn 3s peak and exchange satellite, S<sub>exc</sub>. A higher number of electrons on the Mn 3d level (i.e., higher reduction of Mn) leads to a stronger ΔE. The increase in the S<sub>exc</sub> intensity with respect to Mn 3s is assigned to the impurities (the binding energy area is labelled by the blue box).[21] b) The dependence of ΔE versus number of 3d electrons for various manganese oxides.[39,21]</figcaption></figure>

Therefore, more Co and Ni ions are expected to be involved in
the charge compensation, thereby getting a higher oxidation to
the Co⁴⁺ and Ni³+/Ni⁴⁺ states, as generally accepted in the con-
ventional NCM layered oxides. Instead, Co and Ni are not further
oxidized in the *treated* HE-NCM composite (Figure 4d-iii,e-iii;
Figures S14b-iii, S15c-iii, and S16c-iii,d-iii, Supporting Informa-
tion). The trend in oxidation of transition metals agrees with the
ligand-to-metal charge transfer (LMCT) mechanism inherent for
anionic redox LCR.[41] We suggest that the charge extracted from
M 3d is balanced back by charge transfer from O 2p state accord-
ing to Co⁴⁺ – O²⁻ → Co³⁺ + O⁻ and Ni³+/⁴⁺ – O²⁻ → Ni²⁺ + O⁻
(0 ≤ σ < 1; 1 < n < 2). In the *treated* HE-NCM composite, such
electronic charge transfer is accompanied with a strong distor-
tion at the oxygen site and the molecular oxygen (O₂) formation
(A*⁺ ≈ 530.9 eV,[23]) in addition to the oxygen holes (A* = 528.7 eV)
(Figure 7b-iii). It was recently pointed out that generation of O₂
triggers the formation of closed voids/cracks near the particles
of the active material.[42] Such voids trap molecular oxygen dur-

<figure><img src="image_8.png" /><figcaption>**Figure 6.** Voltage profiles of the first charge-discharge cycle of the *untreated* (black) and the double gas (SO₂ and NH₃) *treated* (blue) HE-NCM cathodes at C/15 rate in 1 M LiPF₆ /EC-EMC solution. Voltage hysteresis is shown by arrows. The charging/discharging potentials, upon which the electronic properties of the cathodes were probed by XPS, SPES, XANES, are labeled by solid filled circles. The plateau regions labeled by 1 and 2 correspond to oxidation of the cations and anion, respectively.</figcaption></figure>

ing cycling that limits the involvement of O₂ in electrochemical activity,[42] and results in an increasing voltage fade.

Whilst our finding reveals that the M ions are no further oxidized in the *treated* cathode being charged to 4.6 V, they are obviously reduced in the *treated* cathode in comparison with charging state of 4.5 V (see the Co L- and Ni L- XANES in Figure 4a-iv,b-iv). A more pronounced reduction occurs at the Mn- site with the distinct contribution of Mn²⁺/Mn³⁺ in Mn L XANES (see Figure 4c-iv; Figure S17a-iv, Supporting Information, in comparison with theoretical calculations, Figure 1e). A nominal oxidation state of Mn is 2.3+ at 4.6 V charging state, as determined from Mn 3s photoelectron emission (Figure 5a-v and Table 2). Reduction of transition metals at the cathode/electrolyte interface is also evidenced by a significant increase of occupation of the O 2p (lattice) vacancies (Figure S18v, Supporting Information). Such a reduction is often assigned with M-dissolution, side reactions, or electrolyte-to-surface electronic charge transfer, i.e., electrolyte oxidation.[31,43] However, our results show that the electrochemical potential of both cathodes has very similar values, whereas the ionization potential of the *treated* cathode is even lower as compared to its *untreated* counterpart (Table 1). Such a finding rules out the reduction of the transition metals at the surface due to an electron transfer from the electronic levels of electrolyte to the M 3d levels of the *treated* cathode, taking into account that the same electrolyte is used for both HE-NCM cathodes. A stronger reduction of the M-ions in the *treated* cathode at 4.6 V cannot be also explained by a deeper penetration of E<sub>F</sub> into the M 3d and O 2p (lattice) levels. This would lead to a stronger involvement of the lattice oxygen in the charge compensation, therefore, to the formation of greater number of holes at the oxygen site. In contrast to that, we observe an essential decrease of the number of O<sup>n-</sup> holes (A<sup>*</sup> = 528.8 eV) in the *treated*
<figure><img src="image_9.png" /><figcaption>**Figure 7.** The evolution of O K XANES of the *treated*- (a) and *untreated*- (b) HE-NCM cathodes versus charging (a, b), and discharging (a) during the first electrochemical cycle with the upper cut-off of 4.8 V. The charging states are colored in red, the discharging states are colored in blue [a(vi-vii)], a, b) as-prepared (i), charged to: 3.7 V [a(ii)], 4.5 V [a(iii), b(ii)], 4.6 V [a(iv), b(iii)], 4.8 V [a(v), b(vi)]; discharged to: 4.4 V [a(vi)], 3.7 V [a(vii)], cycled [a(viii)]. The electron removal from the lattice oxygen leads to the O⁻ hole formation (denoted as A⁺) [a(iii)]. A decrease of the empty states in the lattice oxygen (vanishing of the A⁺ intensity) correlates with an increase of the shoulder B⁺ due to a hole transfer to a non-bonding O 2p orbital [a(iv)]. The B and B⁺ states are assigned to M 3d (e_g) and non-bonding O 2p orbital, respectively. b) A⁺, A⁺⁺ and C are empty states due to a hole formation on the lattice oxygen, molecular oxygen (O₂) and Li₂O, respectively.</figcaption></figure>

cathode (Figure 7a-iv). Importantly, the oxidation of O²⁻ ions (lattice oxygen) does not form O₂ by charging the *treated* HE-NCM composite to 4.6 V that agrees with a decrease in voltage hysteresis (Figure 6).

Mitigation of distortion of the oxygen site correlates with the occurrence of the shoulder B⁺ ≈ 532.8 eV in the O K XANES, which is missing in the *untreated* HE-NCM cathode (Figure 7a-iv,b-iii). The revealed A-B⁺ difference of 3.1 V (Figure S19-iv, Supporting Information) is higher than the crystal field splitting for both the O_h and T_d symmetries. This makes assigning of the B⁺ shoulder to an empty 3d state unlikely, at least for the considered environments with oxygen (see more discussion Section S5, Supporting Information).

Therefore, it is reasonable to assume that the presence of the O 2p empty states (assumed non-bonding) in addition to the O 2p (lattice) and overlapping these oxygen states with the M 3d states, might cause a complex charge transfer between the O 2p- and M 3d- states upon charging the *treated* HE-NCM cathode. As an example, the charge extracted from M 3d orbital is balanced back from one of O 2p states to the e_g or t_2g orbitals, depending on the electronic configuration of M-ion in a charged state of HE-NCM (see Figure S1a, Supporting Information). The relevant occupation of the O 2p (lattice) vacancies in the *treated* cathode, which were formed by charging to 4.5 V, mitigates the distortion of the lattice oxygen and the oxygen loss in comparison with the *untreated* cathode. As for lithium peroxide (Li₂O₂), we cannot support its existence in both *treated-* and *untreated-* cathodes due to the absence of the Raman peak at ≈ 790 cm⁻¹ (Figure S20, Supporting Information) and an inability to resolve peroxide in O K XANES due to its energy vicinity to O₂ (the former is commonly observed at ≈ 530.5 eV of the O K-edge,[44]).

### 2.3.3. *A Deep Stage of Delithiation*

Upon an increasing the charging potential to 4.8 V, E<sub>F</sub> moves deeper into the M 3d hybridized with O 2p (lattice) states manifested by a higher values of the work function (Table 1). In the *treated* HE-NCM cathode, the number of empty M 3d states increases again evidencing further oxidation of the transition metals. This conclusion is supported by an intensity increase of the A and B peaks in the O K XANES (Figure S19, Supporting Information) and the change in the shape of the M L-edges XANES (Figure 4c-v; Figures S13-v and S14a-v, Supporting Information). In the *untreated* HE-NCM cathode, the M ions are no further oxidized (Figure 4d-iv,f-iv), whereas a higher amount of O<sup>n⁻</sup> holes (lattice) and O₂ species are formed in comparison to the status at 4.5 V (Figure 7b-iv). Importantly, the highest oxidation states, such as Co<sup>4+</sup> and Ni<sup>4+</sup>, are not attained in either cathode, thereby ruling out a complete oxidation of the transition metals at 4.8 V (see more discussion on oxidation states of the M ions in Discussion Section S6, Supporting Information). Thus, this partial oxidation of the transition metals in HE-NCM cathodes agrees well with the LMTC model,[41] according to which the Co<sup>4+</sup> and Ni<sup>4+</sup> formal oxidation states are compensated by an electron transfer from O<sup>2-</sup> 2p states. Note that even for conventional layered structure oxides, a strong hybridization of O 2p- and M 3d- states leads to a competition between cationic and anionic redox processes at a deep stage of delithiation.[45] Importantly, the O K XANES of the *treated* HE-NCM cathode shows only a minor distortion of the lattice oxygen at 4.8 V, very similar to that occurring at 4.5 V (Figure 7a-iii,v). The difference is a slightly visible shoulder at A⁺⁺ = 530.9 eV occurred at the higher charging potential that evidences molecular oxygen (O₂) formation at 4.8 V
<figure><img src="image_10.png" /><figcaption>**Figure 8.** Evolution of the electrochemical potential and oxidation states of the cations and the anion upon delithiation during the first charge of the *treated*- (a) and *untreated*- (b) HE-NCM cathodes. a(i), b(i) The *E*<sub>F</sub> position in as-prepared cathodes; a(ii) oxidation of the M cations (δ < 1 and δ < 2 for Co and Ni, respectively) below 4.5 V; (iii) an illustration of the possible electron transfer to vacancy in the Li–O–Li bond via oxidation of the lattice oxygen (O<sup>2-</sup> → O<sup>n-</sup>: 1 < n < 2) at 4.5 V; (iv) oxidation of Co and Ni due to the *E*<sub>F</sub> shift and a partial reduction (labeled by σ) of Co, Ni, and Mn via an electron transfer from the O 2p (lattice) to the M 3d levels at 4.6 V; (v) oxidation of M 3d states and an electron removal from the lattice oxygen at 4.8 V. The O 2p state of N–O bond is not shown. b) (ii), (iii) oxidation of Co and Ni due to the *E*<sub>F</sub> shift, and a partial reduction of Co, Ni, and Mn via an electron transfer from the O 2p (lattice) to the M 3d levels in the 4.5 – 4.6 V range. The O 2p state of Li<sub>2</sub>O activated at 4.5 V is also shown (ii); (iv) oxidation of M 3d states and an electron removal from the lattice oxygen at 4.8 V. The empty- and occupied- states are shown by line and area plots, respectively. Φ<sub>e.</sub><sup>T</sup> (Φ<sub>e.</sub><sup>U</sup>), IP<sup>T</sup> (IP<sup>U</sup>), *E*<sub>F</sub>, and *E*<sub>v</sub> are the work function, ionization potential, Fermi level, and vacuum level, respectively. The ionization potentials were estimated at the VBM and Φ<sub>e.</sub><sup>T</sup> (Φ<sub>e.</sub><sup>U</sup>) measured at *h*ν = 250 eV and *h*ν = 47 eV, respectively (Table 1). The decrease of IP<sup>T</sup> and IP<sup>U</sup>, whereas Φ<sub>e.</sub><sup>T</sup> and Φ<sub>e.</sub><sup>U</sup> remained almost unchanged [a(iii, iv), b(ii, iii)] evidences a double layer potential drop, Δδ<sub>dip</sub>, formed at the cathode/electrolyte interface. Δδδ<sub>dip</sub> in fact lowers the work function.</figcaption></figure>

(Figure 7a-v). This finding indicates a higher stability threshold for the lattice oxygen of the *treated* HE-NCM cathode. The difference in the electronic charge transfer mechanisms upon delithiation of the *treated*- and *untreated*- HE-NCM cathodes is illustrated (Figure 8).

**2.4. Effect of the Electronic Structure Design and Surface Coating on Stability of High-Voltage Cathode Materials at Charging State Above 4.8 V**

Getting higher energy density via a complete delithiation of the cathode materials is attractive, although mitigating the O₂ lattice

release coupled to the structural instability upon a deep charging (≥ 4.8 V vs Li⁺/Li) is a challenge. Considering the intrinsic stability of the high-voltage cathodes, the key question is how far the O 2p states (lattice) can be engineered below the Fermi level to prevent their involvement in charge compensation. A good example is high-voltage LiMPO₄ (M = Co, Ni) olivine structure cathodes with the more ionic character of the chemical bonding between M 3d and O 2p states provided by the inductive effect.[46] A more electronegative P⁵⁺ countercation pulls the electronic density away from the M 3d state, thereby resulting in enlarging the M 3d–O 2p bond distance and reducing the M–O covalency that leads to M²⁺/³⁺ redox couple above 4.8 V. As the result, the M 3d–O 2p states are sharp and well separated from each other, as
<figure><img src="image_11.png" /><figcaption>**Figure 9.** a) The electronic configuration of LCP olivine- structure thin film cathode (without both conductive carbon and binder) near the $E_F$ = 0 derived from photoemission experiments: on the left, the valence band (VB) structure (hv = 1486.7 eV); the right, O K-edge XANES (TEY). The spectra are plotted on a unified energy scale. A = 3.7 eV is associated with the $t_{2g}$ and $e_g$ states of Co, B = 6.0 eV is ascribed to the Co 3d hybridized with O 2p bonding and antibonding states, the C = 10.2 eV and D = 12.6 eV are related to the oxygen and phosphorous strongly bound states, respectively. $\Delta$ = 9.3 eV is the energy gap between the Co 3d and $PO_4^{3-}$. In the O K-edge XANES the most intense peak situated at $\approx$ 5 eV from the $E_F$ is occurred due to the electron transition from the O 1s core-level to the empty O 2p band. The broadness of the O K- edge of phosphates, as compared to layered structure oxides, reflects a strong covalent oxygen-phosphor bonding in the $PO_4^{3-}$ unit.[35] b) The evolution of O K-edge XANES of the *treated* HE-NCM cathode (i) and LCP thin film cathode (ii) by charging to 5.1 V (in the middle) and the first electrochemical cycle with the upper cut-off of 5.1 V (on the right). The assignment of the $A_{HE}$ (i), A (ii) and B (ii) is described (Figures S21a and S24, Supporting Information). c-e) Co L$_3$ (c), Ni L$_3$ (d), Mn L$_3$ (e) XANES of the *treated* HE-NCM cathode charged to 4.8 V (i) and 5.1 V (ii) during the first cycle. f) Cyclic voltamograms of the LCP thin-film cathode coated with LiPON charged to 5.1 V. The CV after the first electrochemical cycle is shown in red. All potentials are referenced versus Li$^+/Li$.</figcaption></figure>

probed by photoelectron spectroscopy (Figure 9a; see more discussion on the electronic structure of LiCoPO$_4$ in Discussion Section S7, Supporting Information). Due to a large energy separation between the M 3d and O2p levels, the $E_F$ moves down upon charging the olivine-type cathodes without pinning the O2p state even at high charging potentials.[35]

The O K XANES of the *treated* HE-NCM cathode and LiCoPO$_4$ thin-film cathode cycled to 5.1 V are shown in Figure 9b. The several chemical states (denoted as $A_{HE}$) arise in the *treated* HE-NCM cathode at the highest charging potential (see O K XANES in more details in Figure S21a, Supporting Information). The shoulder at 528.8 eV and the next peak (529.7 eV) are ascribed to the $O^-$ species and M 3d–O 2p hybridized states, respectively, in agreement with our finding for HE-NCM charged to 4.8 V (Figure 7a-v). The shoulder at 531.1 eV is mostly ascribed to molecular oxygen ($O_2$) formation. In this regard, the spectral features in the O K XANES are very similar to those observed for the *untreated* HE-NCM cathode charged at 4.8 V (Figure 7b-iv). The main difference is an unexpectedly high intensity of the M 3d ($e_g$) state (532.0 eV) and the occurrence of shoulder at $\approx$ 533 eV. The revealed changes in O K XANES are mostly associated with

the formation of carbonyl- (C = O, 531.2 eV), carboxyl- (O–C = O, 532 eV) and carboxylic- (533 eV) groups,[47] as well as $Li_2CO_2$ (533 eV,[48]) due to oxidation of the electrolyte. This conclusion is supported by examination of $E_F$ and IP of HE-NCMs (Table 1). The oxidation potentials of a mixture of EC-EMC carbonate solvent are 4.8–4.9 V (vs Li$^+/Li$),[49] whereas the $E_F$ and the IP equal to $\approx$ 5.4 and $\approx$ 5.6 eV, respectively, when the *treated* HE-NCM cathode is charged to 5.1 V (Figure 3c and Table 1). This means the $E_F$ is shifted below the electronic levels of the carbonate solvents that leads to a hole transfer from the cathode to the electrolyte (see the relevant discussion.[31]). However, such a charge transfer is not accompanied by reduction of the transition metals at the surface in comparison with 4.8 V (Figure 9c,d). Moreover, Ni is even slightly oxidized, although oxidation state of Ni remains lower than 4+ (compare Figure 9d with Figure S22, Supporting Information). The reduction takes place probably at the lattice oxygen, although drawing an explicit conclusion is difficult due to a complexity of the LMCT mechanism at the cathode/electrolyte interface. Nevertheless, the *treated* HE-NCM cathode demonstrates a good reversibility of the electronic structure after the first electrochemical cycle (Figure 9b-i,c,d). This finding
suggests that the practical capacity of the HE-NCM cathodes can be further increased by selecting an appropriating electrolyte stable at 5.1 V, although a long-time cycling stability of the lattice oxygen at such a high potential remains uncertain.

The O K XANES of the LCP thin film cathode shows the lattice oxygen is not involved in redox process at 5.1 V (Figure 9b-ii, middle). The changes in O K XANES caused by delithiation (Figure S24, Supporting Information) are mostly associated with carboxyl groups due to the electrolyte oxidation (A ≈ 532 eV), as well as alteration in the Co 3d–O 2p hybridization and an increasing number of unoccupied Co 3d–O 2p hybridized states (B ≈ 529.9 eV).[35] Thus, enough distance of the O 2p band from E<sub>F</sub> in LCP excludes the involvement of the lattice oxygen in charge compensation at 5.1 V. Nevertheless, remaining of the B peak after cycling (Figure 9b-ii), is a sign of an irreversible occupation of the Co 3d empty states. The issue with an instability of the electrolyte at the surface of the cathode materials in a deep charging state can be solved via design of the interfacial electronic properties. Besides the SO<sub>2</sub> and NH<sub>3</sub> gas treatment reported here for HE-NCM, coating of the redox active oxides also changes their electrochemical potential with respect to the electronic levels of electrolyte, as shown in our recent research.[50] If so, the material for an efficient coating can be selected based on the energy diagrams and electrochemical potential of the redox active oxides obtained from the experiments.[50] Based on such a forecast, coating the LCP thin-film cathodes with lithium phosphorus oxynitride (LiPON) results in a significant improvement in cycling stability (Figure 9f; Figure S25, Supporting Information).[51] Further discussion on the selection criteria of the artificial interfaces is provided in the Discussion Section S8 (Supporting Information) (see also ref. [50]). We note that severe issues associated with dissolution of transition metals at the electrode/electrolyte interface of high-voltage layered structure cathodes, as well as gas evolution upon charging to a high potential can be overcome by choosing an appropriate coating and by selecting an electrolyte with a high electrochemical stability window, respectively, ensuring that the cathode and electrolyte are chemically compatible. In this regard, the modified surface of the *treated* HE-NCM cathode plays also a role for the protective layer, namely, the treatment modifies the E<sub>F</sub> position favorable to avoid the involvement of lattice oxygen into anionic redox at a high charging potential, thereby improving cyclic stability. The above-described approaches can be used to design artificial interfaces for various electrochemical systems, in particularly by means of the electrode material's gas treatment, as it was previously reported.[20,52,53]

### 3. Conclusion

Here, the energy band diagrams concept is introduced to explore the intrinsic voltage limit of high-voltage cathodes and to explain the evolution of the cationic and anionic redox upon electrochemical cycling up to ≤ 5.1 V versus Li⁺/Li. This approach, coupled to examination of the changes in the electrochemical potential and ionization potential, enables to elucidate the mechanisms of chemical reactions occurred at the cathode/electrolyte interface, which differ from redox processes associated with (delithiation reactions. The general remarks on the separation of lithiation/delithiation reactions from interfacial chemical reac-

tions can be found in the Discussion Section S9 (Supporting Information). Specifically, we exam the effect of SO₂ and NH₃ double gas treatment of xLi₂MnO₃·(1-x)Li(M)O₂ (HE-NCM) layered structure cathodes using advanced electron spectroscopy techniques combined with theoretical calculations and operando Raman experiments. The SO₂ and NH₃ treatment alter the surface chemistry and electronic structure via a partial reduction of Co and Mn ions and the formation of Li₂O with the O 2p orbital situated between the hybridized M 3d and O 2p (lattice) levels. The surface modification lifts up the electrochemical potential of the *treated* cathode toward the unoccupied M 3d and O 2p (lattice) electronic states. This shift in the E<sub>F</sub> position is beneficial for the stability of the cathode, because upon delithiation, the Fermi level is shifted downwards into the occupied O 2p states to a less extent, as compared to the untreated counterpart charged to the same potential (Figure 8a(v), 8b(iv)). As the result, the *treated* HE-NCM cathode exhibits reduced oxygen loss, voltage hysteresis and interfacial impedance. In the as-prepared *untreated* HE-NCM cathode, the existence of Li₂O is not confirmed neither at the surface nor in the bulk. This finding means that the relevant Li–O–Li bond (if any) is structurally integrated into xLi₂MnO₃·(1–x)Li(M)O₂ and, therefore, cannot be resolved as a separated phase by using the independent experimental techniques. However, Li₂O is activated by charging the *untreated* HE-NCM cathode to 4.5 V, in agreement with a generally accepted model of anionic redox in xLi₂MnO₃·(1–x)Li(M)O₂.[19] By comparing the E<sub>F</sub> and IP positions with respect to oxidation levels of the electrolyte components, we are able to separate the lattice distortion of HE-NCM from the changes in the interfacial composition induced by electrolyte oxidation which occurs by charging the cathodes to 5.1 V. The apparently stable redox activity of the surface layer even at high potentials implies a high stability which in consequence protect the HE-NCM against oxygen lost. The *treated* HE-NCM cathodes with available SO₄²⁻ bonds at the surface might induce a similar effect in the improvement of the battery performance that incorporation of XO₄ (X = P, S, Si) units into the bulk of the oxide lattice recently discovered for the Li-excess integrated rocksalt-polyanion cathodes.[12] The bulk incorporation of polyanion groups mitigates oxygen loss from the lattice, thereby enhancing the cyclic stability.[54] The energy band diagram approach, coupled to the prediction of E<sub>F</sub> position upon charging the cathodes to a high voltage, is used to design a protective interface for the LiCoPO₄ (LCP) olivine structure cathode which commonly degrades fast upon electrochemical cycling. The LCP coated with lithium phosphorous oxynitride (LiPON) shows significant improvements in cycling stability. To sum up, the proposed concept allows us to predict the intrinsic voltage limit of various cathodes,[31,55] the interfacial stability upon charging batteries to a high voltages,[35] explain the reason of chemical compatibility (or incompatibility) between artificial layers and the electrodes,[50] and an efficiency of the redox active oxides with respect to the oxidation of water.[56] Conceptually, our research opens new horizons in the interface engineering of high voltage cathodes, thereby enabling to increase the energy density of lithium-ion batteries. We believe that the proposed approach is reliable and applicable for various cathode materials, such as spinel structure oxides, disordered rocksalt (DRX) cathodes, high entropy layered oxides, etc.
## 4. Experimental Section

**Materials:** Details of chemical synthesis of 0.33Li₂MnO₃·0.67 LiNi₀.₄Co₀.₂Mn₀.₄O₂ (HE-NCM) cathode materials were reported elsewhere.[19] In this work, the following terminology was used: a) the *untreated* HE-NCM (pristine, without conductive carbon and PVDF additives), *untreated composite* HE-NCM = *untreated* cathode (includes conductive carbon and PVDF additives), *treated* HE-NCM (double-gas treated with SO₂ and NH₃, without conductive carbon and PVDF additives) and *treated composite* = *treated* cathode (double-gas treated with SO₂ and NH₃, includes conductive carbon and PVDF additives). Polyvinylidene difluoride (PVDF) was a binder. The *untreated-* and *treated*-HE-NCM powders (without conducting carbon and PVDF) were pressed into Al-mesh of 7 mm diameter. The LiCoO₂ (LCO), LiNi₀.₂Co₀.₇Mn₀.₁O₂ (NCM), and LiCoPO₄ (LCP) thin film cathode materials were prepared in DAISY-BAT laboratory by radio-frequency (RF) magnetron sputtering technique.[57] The thin-film deposition processes were described in more details.[31,50,55] The conditions for coating of the LCP thin films with phosphorous oxynitride (LIPON) were described elsewhere.[50] The LCP thin films and LiPON layer of ≈ 150 and 50 nm thick, respectively, were studied in the current work. After preparation, HE-NCM materials were stored in a glove box under argon atmosphere (H₂O and O₂ contents were <0.5 ppm).[21]

**Electrochemical Characterization:** For the quasi-in-situ electron spectroscopy experiments, the Swagelok-type battery cells were assembled in a glovebox. Metallic lithium anode and Celgard2500 as the separator were used. 1 M LiPF₆ ethylene carbonate (EC): diethyl carbonate (DEC) (50:50) electrolyte (Sigma Aldrich, Germany), and 1 M LiPF₆ in 4:1 w/w dimethyl carbonate (DMC): fluorinated ethylene carbonate (FEC) + 0.2 wt.% trimethylboroxine (TMB) electrolyte (Solvionic, France) were used for the HE-NCM- and LCP thin-film cathodes, respectively. The electrochemical measurements were performed with the help of a BioLogic (VMP2 Princeton Applied Research) potentiostat. The LCP thin-film cathodes were charged or discharged in cyclic voltammetry mode in the voltage range of 3.0–5.1 V with a 0.1 mV s⁻¹ scan rate at room temperature (RT). The HE-NCM cathode materials were charged or discharged to a defined potential in galvanostatic mode at C/15 rate in the 2.8–5.1 V range at RT. After charging the cathodes to a high potential (≥ 4.6 V), they were kept under charging for several hours to reach thermodynamic equilibrium, until the current drops typically in the range of 0.1–0.5 µA and remains stable. After the electrochemical performance, the cells were disassembled in a glovebox, the cathodes were rinsed in DMC and first dried in a glovebox. For synchrotron measurements at Elettra (Trieste, Italy) and BESSY II (Berlin, Germany), the samples were transferred under UHV conditions by using the UHV suitcases which hold pressure at ≈ 10⁻⁹ mbar. For synchrotron measurements at Diamond Light Source (Oxford, UK), the HE-NCM samples were transported under Argon atmosphere followed by their drying under UHV conditions. For the Raman spectroscopic and electrochemical impedance (EIS) measurements, the HE-NCM materials were pressed on Al₂O₃ mesh under the pressure of 3 tons followed by assembling of electrochemical cell (EL-Cell GmbH, Hamburg, Germany) with a borosilicate glass (Alfa Aesar, U.K) and the Whatman 10 µm filter paper as the separator. Thick lithium foil (0.76 mm) (Alfa Aesar, U.K.) was used as the anode. The BioLogic VSP potentiostat was equipped with an EIS impedance card operating in the frequency range of 1 MHz–5 Hz, with 6 points in decade and a 50 mV amplitude in a SPEIS-experiment (staircase potentio-electrochemical impedance spectroscopy starting at 3.2 V after the 1st full charge-discharge cycle to minimize side-reactions and, therefore, ensure accuracy in the measurements. All potentials were referenced against Li⁺/Li (-3.022 V vs SHE).

**XPS, SPES, RPES, and XANES Analysis:** The HE-NCM cathodes, as prepared and after electrochemical charging (discharging) to a defined-potential, were moved from the glovebox to the UHV systems for photoelectron spectroscopy studies by using various facilities (laboratory and synchrotron sources). All experiments were performed in quasi-in situ regime, i.e., in vacuo sample transfer without contact to air. XPS measurements were performed in the Darmstadt Integrated System for Fundamental Research (DAISY-FUN) using a PHOIBOS 150 spectrometer

(SPECS Surface Nano Analysis GmbH) and a monochromatic Al Kα (hv = 1486.7 eV) source. The base pressure in the analysis chamber was pₓpxs < 10⁻¹⁰ mbar. The photoelectron spectra were collected at an electron escape angle, θ = 90° with respect to the surface. The binding energies were referred to the Fermi level of Ag (or Au) foil in electrical contact with the sample. The probing depth of XPS is d≈ 3 × λ × sin(θ), where λ(Eₖₗ) was the electron inelastic mean free path, which depends on photoelectron kinetic energy (Eₖₗ) and material properties. The probing depth for XPS using Al Kα radiation lies in the range of 5 – 100 Å. The areas under the spectra and energy positions of the photoelectron peaks were obtained by a weighted least-squares fitting of model curves of 70% Gaussian and 30% Lorentzian character to the experimentally measured spectra. CasaXPS and XPSPeak41 software packages were used for the fitting procedure. The background was subtracted using a Shirley-type function. The LCO-, NCM-, and LCP- thin film cathodes after preparation were transferred under UHV conditions (pₐ₉₅ < 10⁻⁸ mbar) for the XPS measurements in DAISY-BAT. The X-ray photoelectron spectra of the HE-NCM powder materials (which do not contain conductive carbon and PVDF) were also measured. In the case of the composite materials, the Ni 2p photoelectron spectra overlap with the F (KLL) Auger peaks contributed from PVDF. This complicates the interpretation of oxidation state of Ni ions when using XPS.

Synchrotron photoemission electron spectroscopy (SPES), soft X-ray absorption near edge structure spectroscopy (XANES) and resonant photoemission spectroscopy (RPES) experiments were performed at the BACH beamline of CNR at the Elettra synchrotron facility (Trieste, Italy). The treated HE-NCM composite cathode materials were delivered to the BACH beamline endstation under ultra-high vacuum (UHV) conditions to avoid exposure to air. The quasi-in-situ Co L₃,₂, Ni L₃,₂, and Mn L₃,₂ XANES were measured in total electron yield (TEY) mode by monitoring the drain current through the sample using a Keithley 428 current amplifier. The energy resolution was better than 250 meV, and the intensities were normalized to the photon flux derived from the total photoelectric current recorded at the last mirror of the beamline. In case of O K-edge XANES measurements, the incoming beam flux was monitored using a gold mesh placed before the sample. The photon energy scale was calibrated by measuring the kinetic energy of Au 4f photoelectrons from a gold reference sample. RPES data were collected at normal emission using a Scienta R3000 electron energy analyzer. The total energy resolution was set to 300 meV. Binding energies were calibrated to the Au 4f signal from a gold reference in electrical contact with the samples. In RPES spectroscopy, valence band photoemission spectra were collected by varying the photon energy across an absorption edge. This can result in interference between direct photoemission and autoionization channels, leading to a resonant enhancement of the valence band electronic states for the element corresponding to the probed edge.

SPES and soft XANES in partial electron yield (PEY) mode were performed at BESSY II (Berlin, Germany). The PM3 undulator beamline equipped with a soft X-ray plane-grating monochromator optimized for elliptical dipole radiation from modern sources enabling the photon energy range of 20–2000 eV.[58] The SoLiAs integrated UHV system with a SPECS PHOIBOS 150 MCD-9 electron analyzer was connected to the beamline.[59] The binding energies of the spectra had been referenced to the Fermi level of a clean Ag-polycrystalline foil. Kinetic energies (Eₖₗ) of photoelectrons in the given XANES (PEY) experiments were varied from ≈ 200 to 850 eV enabling non-destructive depth profiling. The photon energies for the SPES experiments were selected in such a manner to collect the photoelectrons from the same depth.

In addition, Diamond Light Source (Oxford, UK) facilities were used to measure the samples by soft- and hard- XANES X ray monochromators, enabling to vary photon energies in the range of 100–2100 eV and 2.1–20 keV. These experiments were performed at the 109 beamline for Surface and Interface Structural Analysis (SISA). The samples were delivered to the beamline in an UHV suitcase filled by Argon. The XANES experiments were performed in TEY mode.

**Raman Spectroscopy:** All Raman spectra were acquired on a Jobin-Yvon Horiba LabRam HRS 800 (600 lines per mm) under the inert argon atmosphere (spectroelectrochemical cell enclosure), using 632.8 nm excitation from a He-Ne laser and a ×100 objective (Olympus, NA = 0.95). The
excitation power was set to 3 mW. The operando spectra were recorded by using a long-working distance objective (Olympus × 50 LWD, NA 0.55, MPLN).

**Supporting Information**

Supporting Information is available from the Wiley Online Library or from the author.

**Acknowledgements**

This work was supported by the German Science Foundation (DFG project numbers: HA 6128/6-1, HE 4515/12-1, and CH 566/4-1). The authors acknowledge Elettra Sincrotrone Trieste, CNR-IOM (Trieste, Italy) and Diamond Light Source (Oxford, UK) for providing access to synchrotron radiation facilities and for financial support (the project No. SI31579). Helmholz Zentrum Berlin (HZB) is acknowledged for granting beam time at BESSY II under application 222-11382. The authors thank Ilargi Napal for the assistance in synchrotron measurements at the BACH beamline, Pardeep Kumar and Tien-Lin Lee for their support during the synchrotron measurements at the I09 beamline. I.P., S.N., E.M., and F.B. acknowledge funding from the EUROFEL project (RoadMap Esfrí).

Open Access funding enabled and organized by Projekt DEAL.

**Conflict of Interest**

The authors declare no conflict of interest.

**Data Availability Statement**

The data that support the findings of this study are available from the corresponding author upon reasonable request.

**Keywords**

cathode/electrolyte interface, cationic and anionic redox, electrochemical potential and electronic structure, in-situ and operando, LCP, Lithium-rich NCM, XANES and RPES and XPS and Raman

Received: October 30, 2024

Revised: December 6, 2024

Published online:

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