Energy Storage Materials 43 (2021) 275–283

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Energy Storage Materials

journal homepage: www.elsevier.com/locate/ensm

Tailoring the redox-active transition metal content to enhance cycling

stability in cation-disordered rock-salt oxides

Ke Zhou

a , b , Yining Li c , Shiyao Zheng

d , Maojie Zhang

a , Chunyang Zhang

e , Corsin Battaglia

b , Haodong Liu

f , Kuan Wang

g , Pengfei Yan

g , Jianjun Liu

c , ∗ , Yong Yang

a , h , ∗

a State Key Laboratory for Physical Chemistry of Solid Surfaces, and Department of Chemistry, College of Chemistry and Chemical Engineering, Xiamen University, Xiamen 361005, China b Empa, Swiss Federal Laboratories for Materials Science and Technology, 8600 Dübendorf, Switzerland c State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China d Department of Mechanical and Nuclear Engineering, The Pennsylvania State University, University Park, PA 16802, USA e Laboratory of Photonics and Interfaces, Department of Chemistry and Chemical Engineering, École Polytechnique Fédérale de Lausanne, Lausanne 1015, Switzerland f Center for Memory and Recording Research, University of California San Diego, La Jolla 92093, USA g Beijing Key Laboratory of Microstructure and Properties of Solids, Faculty of Materials and Manufacturing, Beijing University of Technology, Beijing 100124, China h School of Energy, Xiamen University, Xiamen 361005, China

a r t i c l e i n f o

Keywords: Cation-disordered rock-salt oxides Redox-active transition metal Redox mechanism

Cycling stability Electronic structure

a b s t r a c t

Lithium-excess cation-disordered rock-salt oxides (DRXs) are investigated intensively as cathode materials for future lithium-ion batteries combining cationic and anionic redox reactions. However, the lattice oxygen redox can cause severe oxygen release resulting in rapid capacity fading. Here, we investigate a series of xLi 2 TiO 3 - (1 - x)LiMnO 2 (0 ≤ x ≤ 1) materials and find that only Li 1.2 Mn 0.4 Ti 0.4 O 2 (x = 0.4) and Li 1.1 Mn 0.7 Ti 0.2 O 2 (x = 0.2) can form phase-pure DRXs, which both deliver high capacity ( > 250 mAh g − 1 ). The newly discovered Li 1.1 Mn 0.7 Ti 0.2 O 2 DRX exhibits remarkably high capacity retention of 84.4% after 20 cycles compared to only 60.8% for Li 1.2 Mn 0.4 Ti 0.4 O 2 . Our result indicates that the irreversible oxygen loss is reduced by raising the Mn content. Theoretical calculations further reveal that increasing the redox-active Mn content from Li 1.2 Mn 0.4 Ti 0.4 O 2

to Li 1.1 Mn 0.7 Ti 0.2 O 2 causes the orbitals near the Fermi level to change from O 2p non-bonding (Li –O

–Li unhy- bridized orbitals) to (Mn –O) ∗ antibonding bands, exhibiting a high O

–O aggregation barrier, preventing O 2 release and resulting in sustained capacity retention. Hence, these new findings demonstrate that regulating oxygen re- dox by tailoring the redox-active transition metal content is an effective strategy to enhance the cycling stability of DRXs.

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. Introduction

The discovery that lattice oxygen (O

2 − ) can participate in the chargeompensation of transition metal (TM) oxide cathodes provides a greatpportunity for a new generation of high-capacity cathode materialsor lithium-ion batteries [ 1 , 2 ]. The oxygen redox reaction was con-rmed in several studies on Lithium excess cation-disordered rock-alt oxide (DRX) cathode materials [ 3 ], such as Li 1.3 Nb 0.3 Mn 0.4 O 2 [ 4 ],i 1.2 Ti 0.4 Mn 0.4 O 2 [5-7] and Li 1.25 Nb 0.25 Fe 0.5 O 2 , etc [ 8 ]. These mate-ials can deliver a high reversible capacity of 250 mAh g − 1 , muchigher than that of lithium-stoichiometric layered TM oxides includ-ng LiCoO 2 , LiNi 0.8 Mn 0.1 Co 0.1 O 2, or Li 2 MnO 3 [ 3 , 9 , 10 ]. Nevertheless,ctivating the oxygen redox generally leads to significant irreversibil-ty in the first cycle, and capacity fading and voltage decay in subse-

∗ Corresponding author: E-mail addresses: jliu@mail.sic.ac.cn (J. Liu), yyang@xmu.edu.cn (Y. Yang).

F

ttps://doi.org/10.1016/j.ensm.2021.08.045 eceived 15 June 2021; Received in revised form 26 August 2021; Accepted 29 Auguvailable online 8 September 2021 405-8297/© 2021 Elsevier B.V. All rights reserved.

uent cycles. Normally, in lithiated layered TM oxides, such as LiMnO 2 nd LiMn 0.5 Ni 0.5 O 2 , TM cation migration and structural reconstructioneads to a transition from the layered to a cation-disordered rock-salt oration-disordered spinel structure and negatively affecting their perfor-ance [11–14] In contrast, DRX materials do not undergo a structural

ransition during cycling, as cations are already disordered in the rock-alt structure [15–17] . Previous studies suggest that the higher energyf labile O 2p states in DRX materials renders O

2 − more prone to ox-dation triggering O 2 loss, responsible for capacity fading and voltageecay .[ 18 , 19 ].

Tremendous efforts have been devoted to understand and mitigate 2 loss in DRX materials [ 18 , 20 , 21 ]. Several groups have substi-

uted a fraction of O

2 − by F − to improve the reversibility of the oxy-en redox and the cycling stability [ 5 , 22–25 ]. Lee et al . reported that

− ions could be incorporated into the bulk of disordered Li–Ni–Ti-

st 2021

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

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o oxides, and that consequently, Li 1.15 Ni 0.45 Ti 0.3 Mo 0.1 O 1.85 F 0.15 con-ains more redox-active Ni 2 + resulting in less oxygen loss compared toi 1.15 Ni 0.375 Ti 0.375 Mo 0.1 O 2 [ 24 ]. Zheng et al . reported a series of newi-rich Li 1.2 Mn 0.4

3 + Mn x 4 + Ti 0.4 − x O 2 (0 ≤ x ≤ 0.4, LMM x TO) by replac-ng a fraction of Ti 4 + by Mn 4 + in Li 1.2 Mn 0.4 Ti 0.4 O 0.2 [ 26 ]. Interestingly,his strategy may manipulate the extent of O redox and enabling highapacity retention by adjusting the Li/TM content in the oxide [ 18 , 27 ].ecent experimental data in several studies also support this view [ 8 ,8 ]. Among them, capacity fading in Li 1.4 Nb 0.4 Mn 0.2 O 2 is much moreronounced than Li 1.2 Nb 0.2 Mn 0.6 O 2 . Research thus proposed that in-reasing the ratio of redox-active TMs might maximize the TM-redoxapacity and reduce the O-redox activity [ 18 ]. All these studies indicatehat regulating the redox-active TM ratio plays a vital role in improvingycling stability. However, no definite conclusion has yet been derivedn how to stabilize oxygen redox and capacity retention by tailoring theedox-active TM content in DRXs.

In this study, we synthesized a series of xLi 2 TiO 3 -(1 - x)LiMnO 2 (0 x ≤ 1) lithium-excess cathode materials. Our comparative study in-icates that Li 1.1 Mn 0.7 Ti 0.2 O 2 with a higher redox-active Mn contentxhibits higher cycling stability with a capacity retention of 84.4%fter 20 cycles. Our density functional theory (DFT) calculations re-eal that tailoring the redox-active transition metal content intrinsi-ally regulates frontier orbitals close to the Fermi level E F , whichovern the redox behavior, [ 29 - 32 ] i.e., O 2p (Li –O

–Li unhybridizedrbitals) in Li 1.2 Mn 0.4 Ti 0.4 O 2 and (Mn –O) ∗ antibonding orbitals ini 1.1 Mn 0.7 Ti 0.2 O 2 . The former results in high charge transfer to anionic and a low O

–O aggregation barrier, resulting in low cycling stabilitynd O 2 release. Whereas the latter generates low charge-transfer fromnionic O and a high O

–O aggregation barrier, promoting reversiblexygen redox and sustained capacity retention.

. Experimental

.1. Materials preparation

Cathode materials with nominal composition Li 1.05 Mn 0.85 Ti 0.1 O 2 ,i 1.1 Mn 0.7 Ti 0.2 O 2 , Li 1.2 Mn 0.4 Ti 0.4 O 2 , and Li 1.3 Mn 0.1 Ti 0.6 O 2 were pre-ared using solid-state calcination. A stoichiometric mixture of Li 2 CO 3 10 wt. % excess, Sinopharm Chemical Reagent, ≥ 98.0%), Mn 2 O 3, andiO 2 (Alfa Aesar, 99.0%) was mixed by ball milling at 300 rpm for 6 in acetone. The mixture was dried and pressed into pellets, followedy calcination at 950°C for 16 h in an inert argon atmosphere. Samplesere subsequently ground into fine powders mixed with acetylene blackt a mass ratio of 9:1, followed by ball milling at 500 rpm for 16 h with-ut acetone to generate a carbon coating. Carbon-coated materials areenoted as, e.g., Li 1.1 Mn 0.7 Ti 0.2 O 2 @C.

.2. Electrochemical measurements

Carbon coated active cathode material, acetylene black, and poly(-inylidene fluoride) binder were ball milled at a weight ratio of 80:10:10n N-methyl-2-pyrrolidone solvent at 300 rpm for 3 h. The slurry washen cast on an aluminum foil at a mass loading of 2.0-2.5 mg/cm

2 tobtain the cathode electrode. 1 M LiPF 6 in ethylene carbonate and ethylethyl carbonate (volume ratio 3:7) (Shenzhen CAPCHEM) was used as

he electrolyte. Coin-type cells were assembled in an argon-filled gloveox, using lithium metal as the negative electrode. Cells were cycled on LAND CT-2001A potentiostat at room temperature.

.3. Materials characterizations

Powder X-ray diffraction (XRD) was performed on the Rigaku Ul-ima IV diffractometer in Cu K 𝛼 radiation ( 𝜆 = 1.5406 Å), respectively.ietveld refinement of the XRD patterns was performed using the Gen-ral Structure Analysis System (GSAS) program [ 33 ]. Sample morphol-gy was characterized on a Zeiss Sigma scanning electron microscope

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SEM) with an energy-dispersive X-ray spectroscopy (EDS) detector usedor elemental mapping. The elemental composition of all samples wasonfirmed by inductively coupled plasma optical emission spectroscopyICP-OES) on a Perkin Elmer Optima 2000-DV. 7 Li magic-angle-spinningolid-state nuclear magnetic resonance (ss-NMR) experiments were per-ormed on a Bruker AVANCE III 400 MHz spectrometer using a 1.3 mmouble-resonance MAX probe at a magic angle spinning frequency of 50Hz with a Hahn-echo pulse sequence (90°pulse –𝜏– 180°pulse –𝜏). 1 LiCl aqueous solution was employed to reference chemical shifts (0

pm). Samples for scanning transmission electron microscopy (STEM)ere prepared by focused ion beam (FIB) on an FEI Helios NanoLab 600perated at 2 – 30 kV. An FEI Titan G2 60-300 transmission electronicroscope equipped with a probe spherical aberration corrector op-

rated at 300 kV was used for high-resolution high-angle annular dark-eld STEM (HAADF-STEM) imaging. Operando differential electrochem-

cal mass spectrometry (DEMS) measurements were performed with awagelok-type cell on a Hiden analytical mass spectrometer [ 7 , 34 ].

igh-purity Ar at a flow rate of 0.6 mL min − 1 was used as the carrieras upon cycling. Ex-situ Mn K-edge X-ray absorption spectrum (XAS)ata were collected in electron yield mode at beamline 4W1B of theeijing Synchrotron Radiation Facility (BSRF). Processing of the X-raybsorption near edge structure (XANES) data was performed using thethena software [ 35 ].

To construct initial DRX configurations for first-principles caclua-ions, the special quasi-random structures (SQS) module based on aonte Carlo algorithm in the ATAT software package was employed 21 , 36-39 ] For Li 1.1 Mn 0.7 Ti 0.2 O 2 and Li 1.2 Mn 0.4 Ti 0.4 O 2 , a 5 × 4 × 2upercell based on the primitive unit cell of the rock-salt structure wasmployed containing 80 atoms. First-principles calculations were per-ormed within the formalism of DFT and the generalized gradient ap-roximation (GGA) of the exchange-correlation function as formulatedy Perdew, Burke, and Ernzerhof (PBE) [ 40 ]. The valence electron-ionnteraction was modeled by the projector augmented wave (PAW) po-ential implemented in the Vienna ab initio simulation package (VASP) 41 , 42 ]. The values of the U parameters for Mn 3d, Nb 3d and Tid states were set to 3.9, 1.5 and 4.6 eV, respectively [ 43 ]. Conver-ence criteria of 0.01 meV per for the total energy and 0.02 eV/Å fortomic forces were applied. The Bader charge determined by the atoms-n-molecule (AIM) method provides a reliable estimate of the redoxrocess during charging [ 44 , 45 ]. Lithium-vacancy configurations inelithiated structures were sampled with the supercell program [ 46 ] ,nd the lowest energy structures were calculated by DFT. The crystalrbital Hamiltonian population (COHP) analysis is performed by usinghe Lobster program [ 47 , 48 ]. The transition state structures and en-rgy of O

–O bonding at each delithiation concentration were calculatedy the nudged elastic band (NEB) method [ 49 ].

. Results and discussion

.1. Material characterizations

A group of xLi 2 TiO 3 -(1 - x)LiMnO 2 (0 ≤ x ≤ 1) (Li 1.05 Mn 0.85 Ti 0.1 O 2 ,i 1.1 Mn 0.7 Ti 0.2 O 2 , Li 1.2 Mn 0.4 Ti 0.4 O 2 and Li 1.3 Mn 0.1 Ti 0.6 O 2 ) Li-excessathode materials was prepared by a solid-state method [ 7 ]. Fromhe analysis of the XRD pattern of Li 1.1 Mn 0.7 Ti 0.2 O 2 (x = 0.2) andi 1.2 Mn 0.4 Ti 0.4 O 2 (x = 0.4) shown in Fig. 1 a, we conclude that bothompositions crystallize into a phase-pure DRX structure with spaceroup Fm -3 m (PDF#16–0223) ( Fig. S1 ). Li 1.05 Mn 0.85 Ti 0.1 O 2 (x = 0.1)nd Li 1.3 Mn 0.1 Ti 0.6 O 2 (x = 0.6) are closer to the Pm nm (PDF#35–0749)nd Fd -3 m space group (PDF#35–0782), respectively, which cannot bessigned to the DRX phase due to the presence of a weak 003 reflection. Li ss-NMR was employed to study the local coordination of lithiumn the synthesized materials. As shown in Fig. 1 b, Li 1.05 Mn 0.85 Ti 0.1 O 2 nd Li 1.3 Mn 0.1 Ti 0.6 O 2 show only one narrow resonance, similar to theuclear magnetic peak of layered material and thus an indication for cation-ordered structure [ 50 , 51 ]. On the contrary, a much broader

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

Fig. 1. Characterization of the crystal structure of the DRXs. (a) XRD patterns ( 𝜆 = 1.5406 Å) and ss-NMR spectra (b) of xLi 2 TiO 3 - (1-x)LiMnO 2 (0 ≤ x ≤ 1). (c) Schematics of the atomic structure of cation-disordered rock-salt materials. Rietveld refinement of XRD patterns for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C (d) and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C (e). STEM-HAADF images for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C (f) and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C (g).

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esonance emerges for Li 1.1 Mn 0.7 Ti 0.2 O 2 and Li 1.2 Mn 0.4 Ti 0.4 O 2 , indicat-ng a distribution of paramagnetic shifts resulting from the different lo-al chemical environments of lithium affected by the paramagnetic ofearby Mn ( Fig. 1 c). In addition, a large number of diamagnetic lithiumnvironments can also be detected at 0 ppm (Li 2 CO 3 , Li –O

–Li) [ 52 ].

herefore, DRX structures only stabilize at a suitable content of Li-Mn 3 + -i 4 + in xLi 2 TiO 3 -(1 - x)LiMnO 2 .

In the following discussions, we focus exclusively on the two DRXaterials, Li 1.1 Mn 0.7 Ti 0.2 O 2 and Li 1.2 Mn 0.4 Ti 0.4 O 2 . The particle mor-hology and corresponding elemental maps are shown in Fig. S 2a-f.he particle size is non-uniform, ranging from 5 to 20 microns, and theistribution of Mn and Ti is homogeneous for both samples. For com-leteness, Fig. S2 g, h show the morphology and particle sizes (2 𝜇m)f the materials after carbon coating. The Rietveld refinement resultsf two carbon-coated materials are presented in Fig. 1 d, e, and Table

1 , which are highly consistent with the DRX structure [ 21 ]. Fig. 1 f, ghow HAADF-STEM images. The spacing between layers extracted fromhe HAADF-STEM images is 0.23, and 0.21 nm for two carbon-coatedaterials, corresponding to (111) and (200) reflections, and both sam-les display good crystallinity with Fm -3 m space group.

.2. Electrochemical performance

Galvanostatic voltage profiles of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C andi 1.1 Mn 0.7 Ti 0.2 O 2 @C during the first 20 cycles are shown in Fig. 2 a, b.he specific discharge capacity of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C decreases from40 mAh g − 1 of its initial cycle to 223 (2 nd ), 208 (5 th ), 184 (10 th ),nd 146 (20 th ) mAh g − 1 , with a low capacity retention of 60.8% after0 cycles. Li 1.1 Mn 0.7 Ti 0.2 O 2 @C shows a discharge capacity of 2571 st ), 257 (2 nd ), 256 (5 th ), 246 (10 th ), and 244 (20 th ) mAh g − 1 , with much higher capacity retention of 84.4% after 20 cycles and muchower first-cycle irreversible capacity of 45 mAh g − 1 . The theoreticalapacities based on Mn 3 + /Mn 4 + redox couple are 145 mAh g − 1 and

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31 mAh g − 1 for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C,espectively. In comparison, the specific capacity of both materials ex-eeds its theoretical capacity, indicating an extensive O

2 − /O 2 n − redox

ouples are active in two cathodes, especially in Li 1.2 Mn 0.4 Ti 0.4 O 2 @C.ssuming that the oxidation of the electrolyte is not considered, theapacity contribution from the O oxidation correspond to 105 mAh

− 1 and 26 mAh g − 1 for two materials. Previous studies showed thatnionic redox leads to severe oxygen loss and thereafter results in aarge polarization [ 53 ]. Therefore, it can indeed be observed thati 1.2 Mn 0.4 Ti 0.4 O 2 @C shows a high irreversible capacity of 102 mAh

− 1 in the first cycle. After 50 cycles, Li 1.1 Mn 0.7 Ti 0.2 O 2 @C shows aetter capacity retention than that of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C ( Fig. S3 a).ig. S3 b shows the average discharge voltage of the two materials. Fori 1.1 Mn 0.7 Ti 0.2 O 2 @C, the average voltage fades only by 0.41 V duringhe first 50 cycles, compared to 0.62 V for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C.

To understand the redox behavior, dQ/dV curves of the two materi-ls are shown in Fig. 2 c, d. A broad peak centered at around 3.5 V associ-ted with the Mn 3+ to Mn 4+ oxidation (labeled Ox1, and correspondingo the sloped region of the first cycle in galvanostatic voltage profiles). sharp peak at around 4.5 V associated with the O

2 − to (O 2 ) n − oxi-ation (labeled Ox2, corresponding to the long plateau were observedor both samples, consistent with the two-step redox process reported inrevious studies) [ 7 , 17 ]. Compared to Li 1.2 Mn 0.4 Ti 0.4 O 2 @C ( Fig. S 4),he Mn 3 + to Mn 4 + oxidation reaction (Ox1) increases in intensity fori 1.1 Mn 0.7 Ti 0.2 O 2 @C, while the intensity of O

2 − to (O 2 ) n − oxidation re-ction is reduced. Furthermore, it is clear that the potential of oxygen ox-dation is lowered at the initial charge process in Li 1.1 Mn 0.7 Ti 0.2 O 2 @C.lready after the second cycle, only a weak Ox2 peak can be observed

n Li 1.2 Mn 0.4 Ti 0.4 O 2 @C in Fig. 2 c, which suggests a largely irreversibleeaction of lattice oxygen [ 5 ]. In contrast, for Li 1.1 Mn 0.7 Ti 0.2 O 2 @C,his peak is clearly visible after the second cycle and is still observedfter 20 cycles. At the same time, the reduction peak has been contin-

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

Fig. 2. Electrochemical properties of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C in Li cells. The voltage profiles (a, b) and corresponding dQ/dV curves (c, d) of the Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C between 1.5 and 4.8 V at a current density of 30 mA g − 1 at room temperature.

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ously shifting to a lower voltage for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C, while fori 1.1 Mn 0.7 Ti 0.2 O 2 @C the shift is much smaller ( Fig. 2 d). This indicatesmproved reversibility of the oxygen redox in Li 1.1 Mn 0.7 Ti 0.2 O 2 @C,hich is beneficial for capacity retention and prevents rapid fading of

he discharge voltage. The electrochemical results suggest an overallower contribution of the oxygen redox to the total discharge capacitynd a more reversible oxygen redox for Li 1.1 Mn 0.7 Ti 0.2 O 2 @C.

.3. Cationic and anionic redox reaction

Ex-situ Mn K-edge XANES spectra of the two materials at differenttates of (dis)charge are given in Fig. 3 a, b to investigate the cationicnd anionic redox. A previous study on Li 1.2 Mn 0.4 Ti 0.4 O 2 @C showedhat Ti does not participate in redox upon cycling [ 7 ]. The Mn K-edgebsorption edges of both materials are also compared to Mn 2 O 3 (Mn 3 +

s reference). For Li 1.1 Mn 0.7 Ti 0.2 O 2 @C, the oxidation state of Mn show slight decrease. Combined with the result of ICP-OES ( Table S2 ), thetoichiometry of Li 1.1 Mn 0.7 Ti 0.2 O 2 does not change apparently from theesigned stoichiometric ratio, resulting in the valence state of Mn willlightly deviate from the designed trivalent (Mn 3 + ) to ensure the chargealance of the compound. After charging to 4.3 V, most Mn is oxidizedo Mn 4 + for both materials, with the absorption edges falling very closeo MnO 2 . When charging from 4.3 V to 4.8 V, only a slight additionalbsorption edge shift is observed for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C, which corre-ponds to the long charge plateau dominated by oxygen oxidation. Thisesult is also in line with our previous X-ray absorption spectroscopyXAS) and mapping resonant inelastic X-ray scattering (mRIXS) study 5 ]. In contrast, for Li 1.1 Mn 0.7 Ti 0.2 O 2 @C, the absorption edge keepshifting, indicating further oxidation of Mn between 4.3 and 4.8 V. Upon

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ischarge, the absorption edge of Mn gradually approaches Mn 2 O 3 andnO.

At the fully discharged state (1.5 V), the Mn valence ini 1.1 Mn 0.7 Ti 0.2 O 2 @C is close to Mn 3 + , while in LMTO it is close ton 2 + . This is mainly because the Mn content in Li 1.2 Mn 0.4 Ti 0.4 O 2 @C

s 20% less than 35 % in Li 1.1 Mn 0.7 Ti 0.2 O 2 and then the Mn ini 1.2 Mn 0.4 Ti 0.4 O 2 @C is further reduced to compensate the charge (O

2 −

O

n − ) during the lithiation process [ 17 ]. During the second charge, thebsorption edge for both shifts to higher energies compared to the firstharge. The Mn oxidation state at the end of discharge (1.5 V) after 5 and0 cycles is lower than that after the first cycle for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C.t explains why the broad Mn reduction peak in the dQ/dV curve ofi 1.2 Mn 0.4 Ti 0.4 O 2 @C in Fig. 2 a gradually decreases in intensity and in-icates that the Mn redox is more stable in Li 1.1 Mn 0.7 Ti 0.2 O 2 @C.

Galvanostatic intermittent titration technique (GITT) curves inig. 3 c reveal severe polarization for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C whenharged from 4.3 V to 4.8 V. This is because the oxidation of latticexygen at high voltages and the formation of solid-electrolyte inter-hase (SEI) [ 17 ]. It also shows a considerably lower voltage relaxationmplitude for Li 1.1 Mn 0.7 Ti 0.2 O 2 @C than that of Li 1.2 Mn 0.4 Ti 0.4 O 2 @Csee also insert in Fig. S5 ). DEMS was conducted on the two materi-ls to monitor the irreversible release of lattice oxygen in the form of 2 gas. As shown in Fig. 3 d, Li 1.2 Mn 0.4 Ti 0.4 O 2 @C generates a more

ignificant amount of O 2 (0.052 𝜇mol g − 1 ) than Li 1.1 Mn 0.7 Ti 0.2 O 2 @C0.024 𝜇mol g − 1 ) at the end of the full charge. It indicates a more pro-ounced loss of lattice oxygen (and possibly surface side reactions) fori 1.2 Mn 0.4 Ti 0.4 O 2 @C.

Fig. 3 e shows Nyquist plots of the electrochemical impedancepectroscopy (EIS) data of the two materials at different states of

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

Fig. 3. The redox mechanism for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C during charge/discharge. Ex-situ XANES spectra at the Mn K-edge of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C (a) and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C (b) at different states of (dis)charge. “1st C 4.3 V / DC 3.0 V ” represents being charged / discharged to 4.3 V

/ 3.0 V in the first charge. GITT curves (c) and operando DEMS of the O 2 evolution (d) of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C. The “standing ” empha- sizes that a voltage hold step at 4.8 V is conducted to collect all the gas generated after charging. Nyquist plots of the EIS data (e) of Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 @C at different (dis)charge states and the corresponding values for the charge transfer resistance R ct extracted from the equivalent circuit fit of the EIS data (f).

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harge before and after cycling. The charge transfer resistance asso-iated with the electrode/electrolyte/lithium metal interface(s) (R ct )f Li 1.2 Mn 0.4 Ti 0.4 O 2 @C, increases from 42 Ω cm

2 to 270 Ω cm

2 afterharging to 4.8 V owing to parasitic side reactions between the elec-rolyte and the electrode and lattice oxygen ( Fig. 3 f). In contrast, R ct or Li 1.1 Mn 0.7 Ti 0.2 O 2 @C increases from 35 Ω cm

2 to only 156 Ω cm

2 .lso, after the 20th cycle, R ct of Li 1.1 Mn 0.7 Ti 0.2 O 2 @C is significantly

ower than that for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C, indicating less acute interfaceeactivity of Li 1.1 Mn 0.7 Ti 0.2 O 2 @C at high voltage (4.8 V) [ 54 ].

.4. Local structure distribution and electronic structure analysis

We performed DFT calculations to explore the impact of the con-ent of active-redox TM on the electronic structure and the chemicalature of lattice redox. A statistical ensemble of DRX supercells withtoichiometry of Li 1.2 Mn 0.4 Ti 0.4 O 2 and Li 1.1 Mn 0.7 Ti 0.2 O 2 was generatedsing Monte Carlo techniques [ 46 ]. For each compound, supercells withelatively low energy were further optimized within DFT to determinehe lowest-energy configuration ( Fig. S6 ). The calculated XRD patternserived from the lowest-energy DRX configurations ( Fig. S7 a, b) areonsistent with the experimental XRD patterns ( Fig. S7 c, d). The de-omination OLi n TM m

(n + m = 6) is used to specify a particular localctahedral-coordinated configuration [ 55 ]. Fig. 4 a, b presents the oc-upation of the six nearest-neighbor sites surrounding each lattice oxy-en site. Obviously, the distribution of the possible local coordinationf lattice oxygen sites is different in the two samples ([OLi 6 ], [OLi 5 Mn],OLi 5 Ti], [OLi 4 Mn 2 ], [OLi 4 MnTi] and [OLi 4 Ti 2 ], etc.). The [OLi 4 MnTi]17.6%) and [OLi 3 Mn 2 Ti 1 ] (18.7%) structures occupy the highest ration all local configurations for Li 1.1 Mn 0.7 Ti 0.2 O 2 and Li 1.2 Mn 0.4 Ti 0.4 O 2 .he different numbers of nearest Li/Mn/Ti around the O, electronictructures for every local structure in octahedral [OLi n TM m

] (n + m = 6)re distinctive and will be reported later.

279

Recent studies have shown that lithium-excess materials create morei –O

–Li configurations in which the O 2p orbitals do not hybridize with TM orbital and facilitate the transfer of the oxygen electron, resultingn oxygen redox reactions [ 31 , 56 , 57 ]. In comparison, the [OLi 4 MnTi]as the highest occurrence in Li 1.2 Mn 0.4 Ti 0.4 O 2 with more Li –O

–Li con-guration (41.5%), whereas [OLi 3 Mn 2 Ti] has the highest occurrence

n Li 1.1 Mn 0.7 Ti 0.2 O 2 with less Li –O

–Li configuration (35.1%). In ad-ition, the tendency towards oxygen oxidation is generally quantifiedy analyzing distinct distributions of frontier orbitals near E F [ 1 , 58 ].

ote that E F for the cathode oxides lies in the e g ∗ band (formed byhe overlap between 3/4d x2-y2 , 3/4d z2 and O 2p orbitals) or t 2g

∗ bandformed by the overlap between 3/4d xy , 3/4d yz , 3/4d xz and O 2p or-itals), and oxidation proceeds by removing an electron from thesetates [ 59 , 60 ]. Hence, Fig. 4 c, d and Fig. S8 plot the density oftates (DOS) and the projected density of states (pDOS) of the O 2prbitals and Mn 3d orbitals for the two materials, respectively. Com-ared to pristine states, a much larger DOS originates from the O 2pith higher energy than Mn 3d orbitals in Li 1.2 Mn 0.4 Ti 0.4 O 2 between 0nd − 1.0 eV. For Li 1.1 Mn 0.7 Ti 0.2 O 2 , the (Mn –O) ∗ bands are located nearhe E F . Besides, according to the COHP analysis ( Fig. S7 e, f) [ 48 ], theMn –O) ∗ bands of Li 1.1 Mn 0.7 Ti 0.2 O 2 are located at higher energy anddjacent to E F . During delithiation, only the (Mn –O) ∗ states are emp-ied and move above E F in Li 1.1 Mn 0.7 Ti 0.2 O 2 , while in Li 1.2 Mn 0.4 Ti 0.4 O 2, oth the (Mn –O) ∗ and unhybridized O bands move successively abovehe E F . The evolution of the Bader charges of the O atoms in dif-erent local structures are shown as a function of state of charge fori 1.2-x Mn 0.4 Ti 0.4 O 2 and Li 1.1-x Mn 0.7 Ti 0.2 O 2 (x = 0.0, 0.2, 0.4, 0.6, and.8) in Fig. 4 e, f, respectively [ 61 ]. The local structure containing the Tiuch as [OLi5Ti] and [OLi4MnTi] exhibit low redox activity because de-reased electronic states of unhybridized O 2p orbital with the presencef Ti 4 + . The oxygen atoms in [OLi 6 ] and [OLi 5 Mn] in Li 1.2 Mn 0.4 Ti 0.4 O 2 xhibit a rapid charge decrease and high anionic redox activity after

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

Fig. 4. Local structure distribution characteristics and electronic structure analysis. Occupation statistics of the six Li/TM sites, which are the nearest neighbor to the lattice O sites, in Li 1.2 Mn 0.4 Ti 0.4 O 2 (a) and Li 1.1 Mn 0.7 Ti 0.2 O 2 (b). DOS of Mn and O electronic structures in delithiation process for Li 1.2 Mn 0.4 Ti 0.4 O 2 (c) and Li 1.1 Mn 0.7 Ti 0.2 O 2 (d). Change in Bader charge of oxygen ions in [OLi 6 ], [OLi 5 Mn], [OLi 5 Ti], [OLi 4 Mn 2 ], [OLi 4 MnTi], [OLi 3 Mn 3 ] and [OLi 3 Mn 2 Ti] configurations in the initial lithiated state for Li 1.2 Mn 0.4 Ti 0.4 O 2 (e) and Li 1.1 Mn 0.7 Ti 0.2 O 2 (f).

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i 1.2-x Mn 0.4 Ti 0.4 O 2 is reaching x = 0.4 (1.6 charges lost) and x = 0.80.7 charges lost). In contrast, most charge reduction is attributed ton atoms in Li 1.1-x Mn 0.7 Ti 0.2 O 2 , and only O atoms in [OLi5Mn] struc-

ure contribute a small fraction ( Fig. 4 e, f, and Fig. S9 ). To summa-ize, oxygen oxidation preferentially occurs in Li 1.2 Mn 0.4 Ti 0.4 O 2 thann Li 1.1 Mn 0.7 Ti 0.2 O 2 , which is consistent with the electrochemical per-ormance.

.5. Two different oxygen redox mechanisms

By systematically analyzing the electrochemical performance andhe electronic band structure evolution of both materials, we found the 2p orbitals that are adjacent to E F of Li 1.2 Mn 0.4 Ti 0.4 O 2 cause removalf the O lone-paired electron and highly active O

n − species redox. Thiseads to large voltage hysteresis, oxygen loss, and capacity fading, suchs in Li 2 MnO 3 and in some Li-rich oxides [ 62 , 63 ]. By contrast, theMn –O) ∗ antibonding orbitals of Li 1.1 Mn 0.7 Ti 0.2 O 2 are located at theop of the valence band and the bottom of the conduction band. It im-lies that redox reactions are mainly dominated by the antibonding or-itals of [MnO 6 ], as reported for Li 2 RuO 3 , and Li 2 IrO 3 , which maintain

280

delicate balance between reversible oxygen redox, O 2 loss, and sus-ainable cycling [ 64 , 65 ]. The two ways of electron removal indicateshat Li 1.1 Mn 0.7 Ti 0.2 O 2 and Li 1.2 Mn 0.4 Ti 0.4 O 2 have different redox mech-nisms: 1) the oxygen redox in Li 1.2 Mn 0.4 Ti 0.4 O 2 is rooted in O lone-pairrbital, 2) while in Li 1.1 Mn 0.7 Ti 0.2 O 2 it mainly comes from Mn and par-ial charge compensation from O.

To further exhibit O

–O formation kinetics influenced by differ-nt oxygen redox mechanisms, we performed transition state cal-ulations for the O

–O formation processes of Li 1.1 Mn 0.7 Ti 0.2 O 2 andi 1.2 Mn 0.4 Ti 0.4 O 2 ( Fig. S10 ) [ 30 ]. The calculated O

–O bond forma-ion barrier is shown as a function of the state of charge in Fig. 5 a.or Li 1.2 Mn 0.4 Ti 0.4 O 2 , the O

–O bond formation barriers gradually de-rease because the charge on the O atoms are reduced upon delithia-ion. The barriers for Li 1.1 Mn 0.7 Ti 0.2 O 2 are much higher than the cor-esponding ones for Li 1.2 Mn 0.4 Ti 0.4 O 2 . The highest barrier with about.0 eV is found at x = 0.2 of Li 1.1-x Mn 0.7 Ti 0.2 O 2 , indicating a low prob-bility for O

–O bond formation. The distinct behavior with respect tohe O

–O bond formation barrier provides an intuitive explanation forhe enhanced O 2 gas evolution and for the lower cycling stability ofi 1.2 Mn 0.4 Ti 0.4 O 2 [ 29 ]. As mentioned above, such higher barriers in

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

Fig. 5. (a) O

–O bond formation barrier for Li 1.2 Mn 0.4 Ti 0.4 O 2 @C and Li 1.1 Mn 0.7 Ti 0.2 O 2 as a function of the state of charge. (b) Charge changes of O

–O no bonding and bonding in Li 1.2-x Mn 0.4 Ti 0.4 O 2

and Li 1.1-x Mn 0.7 Ti 0.2 O 2 . (c) Changes of local charge density around the O

–O bonding process in LiMn 0.4 Ti 0.4 O 2 and Li 0.9 Mn 0.7 Ti 0.2 O 2 . (d) The relationship between Mn/O ratio, energy difference of (Mn –O) ∗ antibonding and O 2p orbitals, and the capacity retention of 20 cycles according this work and references. (e) Schematic electronic band diagram of the two oxygen redox scenarios with different positions of (Mn –O) ∗ antibonding and O 2p orbitals.

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i 1.1 Mn 0.7 Ti 0.2 O 2 should be attributed to more pronounced charge re-uction in O atoms than those in Li 1.2 Mn 0.4 Ti 0.4 O 2 ( Fig. 5 b). To exhibithe charge transfer difference of O atoms between LiMn 0.4 Ti 0.4 O 2 andi 0.9 Mn 0.7 Ti 0.2 O 2 , we plotted their local charge density distributions inhe energy range from -0.5 eV to 1 eV during O

–O bond formation ishown in Fig. 5 c. It can be seen that the charge density around O doesot change much with O

–O bonding in LiMn 0.4 Ti 0.4 O 2 , so it does noteed to lose too many electrons to overcome the lower potential barrieror forming O

–O bonds. On the contrary, the charge density around Ohanges more in Li 0.9 Mn 0.7 Ti 0.2 O 2 , and more electrons are lost, whichause a larger O

–O bond formation potential barrier. So far, we have concluded that the redox-active Mn content is re-

ated to the respective positions of the (Mn –O) ∗ antibonding and O 2prbitals, resulting in different degrees of oxygen reactions and cyclingerformance. Moreover, we find that oxides with high content of activeM show relatively excellent capacity retention in many previously re-orted studies, as shown in Table S3 . Therefore, we also carried out a

281

OS analysis for Mn-based DRXs with different active-Mn content andummarized the results in Fig. S11 . There is no plot for other active-M oxides (Ni, Fe) because of the significant difference in TM bandositions. Combined with electrochemical performance and energy dif-erence between (Mn –O) ∗ antibonding orbitals and O 2p orbitals in thisork (Li –Mn –Ti –O), and a series of Li –Mn –Nb –O and Li –Mn –Ti/Nb –O

–FRXs from literatures [ 18 , 61 , 66 ] , we can plot the relationship be-

ween the capacity retention and the respective positions of the two or-itals in Mn-based DRXs with different Mn content in Fig. 5 d. First, wend that with increasing Mn content the (Mn –O) ∗ antibonding orbitalshift towards E F leading to a larger energetic difference between tworbitals. Second, we find that the capacity retention is proportional tohe energy of the (Mn –O) ∗ band, which is proportional to the Mn con-ent. Overall, Fig. 5 e illustrates the two oxygen redox scenarios withifferent positions of (Mn –O) ∗ antibonding and O 2p orbitals. The O 2prbitals are adjacent to the E F, which shows irreversible oxygen redox,eading to O loss and capacity decay. In contrast, electron removal from

2

K. Zhou, Y. Li, S. Zheng et al. Energy Storage Materials 43 (2021) 275–283

t

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he (Mn –O) ∗ antibonding orbital triggers reversible oxygen redox, de-ermining sustained capacity retention. Consequently, we can adjust thelectronic structure to stabilize the oxygen redox and capacity retentiony tailoring the redox-active TM content in DRXs.

. Conclusions

In summary, a series of xLi 2 TiO 3 -(1 - x)LiMnO 2 (0 ≤ x ≤ 1) lithium-xcess oxides were synthesized and investigated. We show that a phase-ure rock-salt structure is only formed at precise ratios of Li, Mn and Ti.he oxides (Li 1.1 Mn 0.7 Ti 0.2 O 2 ) with a high content of the redox-activen shows less O 2 release and much-improved cycling stability. We pro-

ose that (Mn –O) ∗ antibonding orbitals close to the EF generates lowharge transfer from O and a high O

–O aggregation barrier, promot-ng reversibility of the oxygen redox and sustainable cycling stability.onversely, when the O 2p orbitals is close to the EF, it exhibits highharge transfer from O and a low O

–O aggregation barrier, resultingn low cycling stability and O 2 release. Thus, our study provides somemportant guidelines for understanding cationic-anionic activity and de-eloping cathodes with high capacity and cycling stability by tailoringhe redox-active transition metal content in DRXs.

eclaration of Competing Interest

There are no conflicts to declare.

RediT authorship contribution statement

Ke Zhou: Conceptualization, Investigation, Methodology, Data cu-ation, Writing – original draft. Yining Li: Methodology, Formalnalysis, Writing – review & editing. Shiyao Zheng: Data curation.aojie Zhang: Data curation. Chunyang Zhang: Data curation. Corsin

attaglia: Supervision, Methodology, Funding acquisition, Writing – re-iew & editing. Haodong Liu: Writing – review & editing. Kuan Wang:

oftware. Pengfei Yan: Software. Jianjun Liu: Conceptualization, For-al analysis, Writing – review & editing. Yong Yang: Supervision,roject administration, Funding acquisition, Writing – review & editing.

cknowledgements

This work is financially supported by the National Key Researchnd Development Program of China (grant no. 2018YFB0905400,016YFB0901502), National Natural Science Foundation of Chinagrant no. 21935009, 21761132030, 21973107), Shanghai Academicesearch Leader (grant no. 20XD1404100). The authors sincerely thank

he assistance of the Beijing Synchrotron Radiation Facility and the Na-ional Synchrotron Radiation Laboratory for the beam time and kindupport for the measurements of hard X-ray absorption spectroscopy.e Zhou would like to thank the Chinese Scholarship Council for theirnancial support to her in-split Ph.D. study in Switzerland. Corsinattaglia acknowledges financial support by the Ernst Göhner Founda-ion.

upplementary materials

Supplementary material associated with this article can be found, inhe online version, at doi: 10.1016/j.ensm.2021.08.045 .

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