Reversible Dual Anionic-Redox Chemistry in NaCrSSe 1 with Fast Charging Capability 2

: Utilizing the anionic redox reaction opens new approaches for the development of new battery cathode materials with extra capacities. Although, it suffers from several obstacles such as voltage hysteresis and sluggish kinetics. In this paper, a 35 new layered chalcogenide-based on dual anionic-redox reaction is reported. The newly designed layered NaCrSSe exhibits the capacity of almost all Na + intercalation/deintercalation (137 mAh g -1 at 50 mA g -1 ), and a unique charge/discharge feature with a small polarization of 0.15 V and high energy efficiencies of ~92% in 39 initial cycles. Furthermore, a superior high-rate charge capacity of 115.5mAh g -1 (83.7% 40 retention) was achieved at 27.8 C (4000 mA g -1 ), which is impressive in all bulk materials for sodium-ion batteries. Systematic characterization studies on evolution and DFT calculation show the charge compensation of S and Se anions during cycling. These results will enrich the anion redox chemistry and provide valuable information for developing new anion redox based cathode materials with high capacity and fast kinetics. couples by tuning the hybridization of dual anionic Se 4p and S 3p bands. The high columbic efficiency (99.7%) and energy efficiency ( ＞ 90%) in the initial cycles and 2 excellent rate capability (the high charge capacities of 115.5 mAh g -1 at the rate of 27.8 3 C) could be attributed to the positive effect with dual anionic in NaCrSSe. In addition, 4 NaCrSSe are believed to reduce the chances of releasing S 2 owing to its minimizing the 5 probability of the chemical formation of S 2 as compared to NaCrS 2 . Our results fill the 6 void of multi anionic-redox chemistry in layered compounds for batteries, enrich the 7 intercalation charge compensation chemistry, and guide the design of layered 8 intercalation compounds by new approach. This work, along with the former reports of 9 anionic-redox based layered sulfides, explores a new pathway of layered chalcogenides 10 to solve the practical problem of anionic-redox for high-rate batteries. We expect a 11 combination of multi anionic-redox and traditional multi cationic-redox to design high- performance practical materials in the future. dimethyl (1:1 with 5 wt% (FEC) as the electrolyte. Galvanostatic charge-discharge experiments and 3 GITT (Galvanostatic Intermittent Titration Technique) were carried out on a Land 4 CT2001A battery tester at room temperature. Cyclic voltammograms (CV) and AC 5 impedance test were performed on an electrochemical workstation CHI 660B. X-ray absorption spectroscopy: XAS of Cr and Se- K edge was measured at beamline 7 8-ID and 7-BM of National Synchrotron Light Source II (NSLSI-II) in Brookhaven 8 National Laboratory (BNL). The data were collected in the transmission mode. The 9 reference spectrum of Cr metal foil or Se were simultaneously collect during each 10 measurement and used for energy calibration. Ex-situ S K-edge XAS spectra were 11 obtained in the fluorescence mode at beamline 4B7A in Beijing Synchrotron Radiation 12 Facility and beamline 9-BM of Advanced Photon Source (APS) at Argonne National 13 Laboratory (ANL). The X-ray absorption near-edge structure (XANES) and Extended 14 X-ray absorption fine structure (EXAFS) spectra were processed using the Athena 15 software package. 45

The intercalation chemistry of layered compounds exhibits a wide variety of 2 charge compensation mechanisms by the oxidation and reduction of transition metal 3 cations, anions or simultaneous cations and anions during the reversible 4 intercalation/deintercalation of the guest ions in the host framework 1-4 . The capacity of 5 electrode materials is closely related to the redox couples during the electrochemical 6 process such as cationic redox (M n+ /M (n+1)+ , where M is a transition metal), or anionic 7 redox ((2O 2-/(O2) n-) or (2S 2-/(S2) n-)). Therefore, high energy density can be achieved 8 through the proper selection of a cation group, or a combination of the cation-anion 9 group. As the traditional multi-cationic-redox-based materials gradually reach the limit 10 of their theoretical capacities, it's becoming more and more necessary to utilize the 11 anionic redox reaction to further enhanced the capacity of reversible ion batteries to 12 satisfy the demand for high-energy-density. For example, the Li-rich material of 13 Li[Li0.2Ni0.13Co0.13Mn0.54]O2 contains three cation redox couples of Ni 2+ /Ni 3+ , Ni 3+ /Ni 4+ , 14 and Co 3+ /Co 4+ , together with one anionic redox couples of 2O 2-/(O2) n-5,6 , which nearly 15 double the capacity of today's LiCoO2 and LiNi1/3Mn1/3Co1/3O2 with the extra 16 contribution from anionic redox reaction. The P2-type compound of 17 Na2/3[Mg0.28Mn0.72]O2 with Mn 3+ /Mn 4+ redox and 2O 2-/(O2) nredox was reported to 18 obtain a capacity of ~230 mAh g -1 in Na ion batteries. 7,8 Also, P2-Na0.72[Li0.24Mn0.76]O2 19 was reported with anionic redox reaction of O anions and Mn 3+ /Mn 4+ redox, delivers 20 the highest energy density (700 Wh kg -1 , 270 mAh g -1 ) among all Na cathode materials. 21 9 However, the widely-studied redox of O anions is likely difficult to break through the 22 bottleneck for practicability, which are mainly voltage hysteresis (low energy efficiency) 23 and sluggish kinetics (low rate capability). 4,10 While, layered chalcogenides may be the 24 substitution of oxides. Single anionic redox chemistry in layered NaCrS2 was reported 25 not long ago, which undergoes S2 n-/2S 2redox, while the contribution of Cr 4+ /Cr 3+ is 26 negligible during the deintercalation/intercalation of Na ions. 11 Furthermore, a layered 27 NaCr2/3Ti1/3S2 was successfully developed to fully utilize the synergy effect of the 28 cation (Ti 3+ /Ti 4+ ) and anion ((S 2-/S n-), (S n-/(S2) n-)) redox. As expected, the capacity 29 from ~100 mAh g -1 in NaCrS2 increased to ~190 mAh g -1 . 12 Very recently, a Li-rich 30 sulfide Li1.33 - 2y/3Ti0.67 - y/3FeyS2 with cationic (Fe 2+ / 3+ ) and anionic (S 2-/S n-, n < 2) redox 31 processes was reported, which obtained a high capacity of ~245 mAh g -1 and undergo 32 comparable low voltage hysteresis and fast kinetics. 13 These studies advanced the 33 development of new layered compounds with high-energy-density. 34 It is worthwhile to point out: although the utilization of multi-cationic-redox or 35 cation-redox coupled with single-anionic-redox has been widely reported, no similar 36 work on multi-anionic-redox-based layered compounds has been reported for 37 rechargeable ion batteries. Besides, it is well known that cation doping was generally 38 adopted to tune the band structure and property of layered transition metal oxides 3,14,15 39 and sulfides 12,16,17 , the choice of anions is always monotonous, however. As a new 40 approach, in this work, a strategy of anion doping is applied. The reversible dual-41 anionic-based redox chemistry in NaCrSxSe1-x is reported for the first time. By 42 introducing the dual anion-redox, the designed NaCrSSe electrode shows a reversible 43 specific capacity of 137 mAh g -1 (96% of the theoretical capacity) with a small 44 polarization of 0.15 V (no voltage hysteresis), high energy efficiency of ~92% and 1 unapproachable high-rate charge capability (83.7% retention at 27.8C), which is 2 impressive in all bulk materials for sodium-ion batteries, to the best of our knowledge. 3 These superior properties have never been reported in other layered materials based on 4 anionic redox chemistry. In addition, systematic characterization studies and DFT 5 calculation was utilized to study the structure evolution in dual anionic-redox of S and 6 Se, as well as to explain the high-rate charge capability.   Figure 1c gives the corresponding annular bright-field (ABF) image, note 18 that the ABF images show a Z 1/3 relationship, therefore, both light and heavy atom 19 columns are visible simultaneously. 18,19 In ABF image, Na columns are also 20 distinguished clearly (see also the blue line in Figure 1b). The configuration of atom 21 columns displays the O'3-type structure, which is a slightly distorted O3-type structure 22 ( Figure S1). The distance of S/Se-Cr-S/Se slabs is measured to be 6.4 Å. It should be 23 noted that a small amount of impurity phase is found ( Figure S2).

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X-ray diffraction (XRD) spectra of as-prepared NaCrSxSe2-x powders were 25 measured to further confirm the structure and purity of samples. As shown in Figure   26 S3, XRD patterns of all the prepared NaCrSxSe2-x (x=0.5,1,1.5) samples show the O3-27 type structure. It can be seen that, with the increase of Se ratio, peaks shift to lower 28 degree, and parameters a and c are both elongated (simply derived from distinctive (001) 29 and (104) peaks). As we know, atom size of Se is larger than that of S and the calculated 30 parameters a and c values are 3.632Å and 19.351Å for NaCrS2 and3.826Å and 22.447 31 Å for NaCrSe2, respectively, so it's reasonable that lattice constants of a and c increased 32 as S substituted by Se, leading to the decrease of diffraction 2θ value. The structure and 33 lattice parameters of NaCrSSe is further confirmed by the Rietveld refined results of 34 synchrotron XRD data of NaCrSSe powders as presented in Figure 1d. A distorted O'3-35 type layered structure (P1) for NaCrSSe and an OP6 model for impurity phase Cr(S/Se)x 36 are adopted in the Rietveld refinement and the structure parameters are listed in Table   37 S1. The refined lattice parameters of NaCrSSe are a =3.7683(5) Å, b =3.6103(6) Å, c =  (Table S2) for NaCrSSe, which is almost consistent with the NaCrSSe 10 stoichiometry, while the values of NaCrS1.5Se0.5 and NaCrS0.5Se1.5 also accord with 11 stoichiometry as shown in Table S3 and S4. The pristine NaCrSSe mainly consist of 12 high-density flaky structure with thickness of 2~3 micrometer ( Figure S6).  indicates that extra capacity in Se substituted samples may be attributed to the redox of 42 Se in addition to the S. 43 densities from 50 mA g -1 (0.35C) to 4 A g -1 (27.8 C). At the rate of 3A g -1 (20.8 C) and 1 4 A g -1 (27.8 C), high charge capacities of 122.4 mAh g -1 (0.85 Na + /CrSSe, 88.7% of 2 that at 50 mA/g) and 115.5 mAh g -1 (0.80 Na + /CrSSe, 83.7% of that at 50 mA/g) are 3 achieved, respectively. While the discharge capacity of 64.4 mAh g -1 (47.0% of that at 4 50 mA g -1 ) is obtained at 4 A g -1 . This result indicates that NaCrSSe cathode has super 5 first-class charge rate performance and is able to store charge more than 85% in a 2~3 6 minutes, which is an impressive rate performance among in all sodium storage layered 7 bulk materials reported previously ( Figure S9). Apparently, NaCrSSe does not exhibit 8 the feature of sluggish kinetics for anion redox based materials 4 . Such a fast rate ability 9 for a bulk NaCrSSe cathode material without nanometer-sized, porous structures,  Table S6.  Figure S12. These results indicate a non-crystallization process and reduction of 42 crystallinity of NaCrSSe during the electrochemical process. To obtain clear insight to 43 the local structure evolution of NaCrSSe after electrochemical process, x-ray pair distribution function analysis (PDF) was utilized. Compared with the XRD technique 1 which uses Bragg scattering only and provides long-range average structural 2 information, the PDF method utilizes both Bragg diffraction and diffuse scattering. 3 Therefore, PDF method can provide local structural information especially the atomic 4 pair distribution, which is related to chemical, structural and morphological information 5 of materials. 22-28 PDF of NaCrSSe collected at pristine, charged and discharged state 6 within the long-r range of 0-85 Å are shown in Figure S13. Stronger G(r) peaks are 7 observed clearly from PDF of the pristine sample even at the large atomic pair distance 8 (60-70 Å), indicating the good crystallinity of the pristine sample. However, the 9 intensities of G(r) peaks after charge-discharge cycling reduce significantly, especially 3b) also reveals that O'3-NaxCrSSe (0 < x < 1) is not thermodynamically favorable with 28 the positive formation energies of 10~20 meV/f.u., implying the phase separation of 29 O'3-NaCrSSe and O1-CrSSe. This is in good agreement with XRD patterns in Figure   30 3a.  Table S7. It should be noted that impurity Cr(S/Se)x phase can be  is fully reversible. In addition, the broad peak B at ~12670 eV (blue area in Figure 4b) 19 caused by multiple scattering in the coordination sphere 31 shifts to high energy and its 20 intensity reduces upon charging. It shifts back to the pristine state after fully discharge.

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All these changes demonstrate that the charge compensation is partly achieved by the 22 redox of Se anions.

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S K-edge XAS: XAS spectra at S K-edge at different charge/discharge stages is 24 presented in Figure 4c. The relative intensity of peak A at ~2470eV (red area in Figure   25 4c) varies greatly during the cycling ( Figure S15), and the evolution of peak B at 26 ~2480eV (blue area in Figure 4c) has the similar moving trend as that for Se K-edge.

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The peak A could be attributed to the transition of 1s electron to the unoccupied Cr 3d-28 S 3p hybrid orbitals. During the charge process, the emerging and growing of the 29 shoulder peak located at ~2468 eV and addition peak formation at 2470.7 eV are clearly  The corresponding FT-EXAFS spectra of Cr, Se and S are also shown in Figure 4d, 4e 36 and 4f respectively. Note that the FT-EXAFS spectra have not been phase-corrected so 37 that the actual bond lengths could be ~0.4 Å longer in our case. 11,34 As shown in Figure   38 4d, the peak at 2.1 Å of the pristine sample can be attributed to the closest S/Se ions to 39 the core Cr ions, and it shifts to 2.0 Å upon charging. The peak at 2.2 Å in Figure 4e at 40 the pristine state is corresponding to the closest Cr ions to the core Se, which shift to 41 2.1 Å after fully charging. The peak at 2.1 Å for the pristine state in Figure 4f is 42 corresponding to the closest Cr ions to the core S, which shift to 2.0 Å after fully 43 charging. During charging (desodiation), Cr-S/Se bonds shrink, also shown as the decrease in lattice parameters "a" and "b" obtained by XRD, caused by the contraction 1 of the S and Se anion radii at higher oxidation states. These structural changes are in 2 line with the mechanism of 'normal unit cell breathing'. 35 During discharge, the 3 changes of Cr-S/Se distances reverse, indicating the reversible evolution of the local 4 structure after sodiation. It should be mentioned that, a small shoulder peak at 1.7 Å in 5 Se FT-EXAFS (Figure 4e) and at 1.6 Å in S FT-EXAFS (non-symmetry of the main 6 peak in Figure 4f) arises after charging, which indicates of the formation of (S/Se)2 m-7 species. 12 8 X-ray pair distribution function (PDF): The structural evolution of NaCrSSe 9 electrode material during the charge and discharge process is also investigated by using 10 X-ray pair distribution function (PDF). As shown in Figure 4g, the first peak at 2.52 Å 11 attributes to the Cr-S/Se bond and the third peak at 3.60 Å attributes to Cr-Cr and S/Se-   Table   26 S8. The shift of major doublet at low binding energy to higher energy upon charging 27 and shift back to lower energy upon discharge indicates the reversible formation of the 28 holes of Se nand S n-. The binding energy of minor doublet peak evolved from S or Se 29 XPS spectrum at (163.6 eV, 164.8 eV) or (54.6 eV, 55.8 eV) could represent the 30 formation of (S/Se)2 mdimers according to these previous XPS studies of S2 2and 31 Se2 2-. 36,37 As well known, XPS is a surface analytical measurement. It should be 32 assumed that the composition from the surface is identical to that from bulk combining 33 with XAS data. Based on the ratio of their areas, it can be estimated that ~1/7 of Se and 34 S dimerize, while the rest ~6/7 anion generate Se nand S nholes upon charging.

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Formation of (S/Se)2 mdimers is also further reveled by the peak formed at 450   As the schematic (Figure 6c) shown, the Se 4p states lay above S 3p states, 18 substituting S partly with Se in layered NaCrS2 will raise the total valance bands of 19 S/Se anions, and make anions easier to take part in charge compensation. Apparently, 20 the dual anionic redox of S and Se is quite different from those oxygen-redox materials compensation to a great extent, but also activate S to participate more in charge 32 compensation.

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The diffusion barrier of Na is calculated and shown in Figure S20. We assumed 34 three intermediate phases -Na0.92CrSSe, Na0.5CrSSe and Na0.08CrSSe between pristine 35 NaCrSSe and full charged CrSSe. It can be concluded that with the decrease of local 36 Na content from 0.92 to 0.08, the diffusion barrier of Na declines from 0.34 eV to 0.19 37 eV, owing to less Na-Na electrostatic repulsion, which also is contributed to the fast 38 kinetics of NaCrSSe. anions such as S, Se and Te, the anionic redox process will be enhanced due to the 8 increased overlap of transition-metal(d)-anion(p) orbitals 43 . In NaCrSSe, Se has less 9 electronegativity compared with S to provide additional capacity, as well as high 10 electronic conductivity. In addition, introducing Se would enlarge lattice parameters of 11 comparing with sulfides, and may increase the mobility of Na ions. EXAFS. This dimerization reaction should be triggered by the formation of non-29 coordinated S3p and Se4p orbitals due to the Cr migration to the Na sites in layered 30 structure, which could be investigated by STEM and DFT results. This dual anionic-31 redox charge compensation mechanism explains the high capacity corresponding to one 32 Na per NaCrSSe in initial charge/discharge process.

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The lattice parameter of c axis in O'3-NaCrSSe and O3-NaCrS2 are 20.0742 Å 34 and 19.3506(1) Å, respectively. Apparently, the slab distance of S-Cr-S slabs increases 35 by substituting S partly with Se in O3-NaCrS2 structure. This undoubtedly provides a 36 more favorable fast channel for ion transport. Na + diffusion coefficients from CV curves 37 is measured to be 2.67×10 -9 cm 2 /s for charge process, which is almost superior to all 38 sodium storage bulk materials reported previously (Table S6). The diffusion barrier of 39 Na is as low as 0.19 eV, which is very impressive among the O3 type sodium cathode 40 materials (Table S6). Since the top of Se 4p states lie above S 3p bands, introducing of Co., Ltd, purity:99.95%), S(Alfa Aesar, purity: 99.5%) and Se(aladdin,purity:99.9%).

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Note that Na2S, S and Se are excess of 5%. The starting powders were carefully 19 grounded using a mortar and pestle and then pressed into a pellet. The pellet was heated 20 at 900℃ for 6 hours in Ar phenomena. After a cooling process, the resulting product 21 was transformed into Ar-filled glove box immediately to avoid contacting with air.  Impedance spectra of a NaCrSSe pellet (mass: 0.4810 g, diameter: 12 mm, thickness: 10 1.302 mm) at 25 ℃, the equivalent circuits used for fitting the impedance spectra is 11 inserted in the figure.