X-ray Absorption Spectroscopy and In-Operando Neutron Diffraction

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C: Energy Conversion and Storage; Energy and Charge Transport

X-ray Absorption Spectroscopy and In-operando Neutron Diffraction Studies on Local Structure Fading Induced Irreversibility in a 18650 Cell with P2-Na2/3Fe1/3Mn2/3O2 Cathode in Long Cycle Test Tsan-Yao Chen, Bin Han, Chih-Wei Hu, Yuh-fan Su, Yong-Xiang Zhou, Hung-Yuan Chen, Ping-I Pan, Chun-Ming Wu, Alice Hu, Ji-Jung Kai, Yung-Der Juang, and Chia-Chin Chang J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b02908 • Publication Date (Web): 04 Jun 2018 Downloaded from http://pubs.acs.org on June 4, 2018

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The Journal of Physical Chemistry

X-ray Absorption Spectroscopy and In-operando Neutron Diffraction Studies on Local Structure Fading Induced Irreversibility in a 18650 Cell with P2-Na2/3Fe1/3Mn2/3O2 Cathode in Long Cycle Test

Authors: Tsan-Yao Chen,a,b*† Bin Han,c Chih-Wei Hu,a Yuh-Fan Su,d Yong-Xiang Zhou,e Hung-Yuan Chen,c Ping-I Pan,d Chun-Ming Wu,f Alice Hu,c Ji-Jung Kai,c Yung-Der Juang,e*† and Chia-Chin Changd*†

Affiliation: a.

Department of Engineering and System Science, National Tsing Hua University,

Hsinchu 30013, Taiwan. E-mail: [email protected]; TEL: +886-3-5715131#34271. b.

Institute of Nuclear Engineering and Science, National Tsing Hua University,

Hsinchu 30013, Taiwan. c.

Department of Mechanical and Biomedical Engineering, City University of Hong

Kong, Hong Kong, China. d.

Department of Greenergy Technology, National University of Tainan, Tainan 7005,

Taiwan. E-mail: [email protected]; TEL: +886-6-268-9274, Fax: +886-6-260-2205. e.

Department of Materials Science, National University of Tainan, Tainan, 70005

Taiwan. E-mail: [email protected] f.

National Synchrotron Radiation Research Center, Taiwan

† Tsan-Yao Chen, Yung-Der Juang, and Chia-Ching Chang are co-correspondence to this article.

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Abstract P2-Na2/3[Fe1/3Mn2/3]O2 (NFMO) crystal with a maximum capacity of ~150 mAh was synthesized by a solid-state annealing method and used as a cathode in a sodium ion battery. By combining focused-ion beam section scanning electron microscopy, ex-situ X-ray absorption spectroscopy, X-ray photoemission depth profiling, and in-operando neutron diffraction, we found that Na ion intercalation and extraction distort the local structure in NFMO crystal, resulting in irreversibility of the sodium ion battery (SIB). This reaction pathway is controlled by the transformation kinetics of the Fe sites from octahedral (Oh) to tetragonal (Td) in the charge and discharge processes. For a SIB operated at 2.0 to 3.8V, steady kinetics between the Na intercalation and chemical state evolution on the Fe sites enable the homogeneous restructuring in both local and global regimes in NFMO crystal. For a SIB operated at 2.0 to 4.5V, substantially higher kinetics in the Fe chemical state evolution induce a dramatic lattice expansion. This expansion cracks the interface between the P2 and Na intercalated regions, thereby causing substantial irreversibility of NFMO in a SIB.

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Introduction Rechargeable metal-ion batteries hold promise for energy storage applications in public transportation platforms (i.e., hybrid, plug-in hybrid, complete electric vehicles) or buffers of green energy conversion systems (wind, solar power plants). Considering the variety and adaptability of energy conversion and its applications, the abundance and cost of metal ions are key factors in mass energy storage systems. In these considerations, the sodium ion battery (SIB) has attracted much attention in the academic and industrial sectors.1 A SIB comprises electrochemical and geometric configurations identical to those of a lithium ion battery (LIB), so such a device could be produced without significant modification of production facilities. However, unlike LIBs, SIBs have higher humidity sensitivity, a larger ion storage energy barrier, a lower capacity, and higher irreversibility of electrode materials, which are fatal issues for the commercialization of SIBs.2-6 P2- and O3-type NaxMO2 (M state 3d transition metals, including Co, V, Ni, Mn, and Cr) demonstrate a capacity ranging from 90 to 160 mAh g-1 and are potential candidates as cathodes for SIBs.7-11 On the basis of elemental abundance, NaxFeO2 (NFO) is a proper candidate for commercial production of SIBs. However, Fe4+ ions are not stabilized in the crystal framework under normal charge–discharge conditions. Proper doping by a 2nd transition metal improves the stability and capacity of NFO. If 50% of the Fe sites are replaced with Mn dopant, Fe1/2Mn1/2O2 shows a reversible capacity of 170–180 mAh g-1 as the cathode in a SIB. Such a capacity is compatible with that of the cathode in a lithium ion battery and thus suggests a potential for the commercialization of SIBs.12-14 In this study, P2-Na2/3[Fe1/3Mn2/3]O2 (NFMO) crystal was prepared by thermal annealing method. An experimental NFMO with an optimum capacity of ~190 mAh g-1 was achieved. To disclose the factors in capacity loss, the local and crystal structures of NFMO at selected voltages in coin cells operated between 2.0 to 3.8 V and 2.0 to 4.5 V and after long-cycle charge-discharge tests in 18650 full cells were cross-referenced. Our results revealed that in the first charge process, local structural distortion appears at the Fe sites and not at the Mn sites of NFMO crystals. After 100 cycles, the structure transforms from octahedral to tetragonal symmetry, indicating the chemical interaction of Fe sites in NFMO crystal. 3



Such a transformation can be attributed to the formation of amorphous NaFeOx structures. In SIBs operating between 2.0 to 4.5 V, Fe-O bonds break upon Na+ intercalation at the Fe sites of the NFMO. Consequently, the dissolution of Fe and Mn atoms, followed by their deposition on the hard carbon counterpart electrode, results in substantial capacity irreversibility in the SIBs. Details of the characterization results and discussions are provided in later sections.

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Experimental P2-Na2/3[Fe1/3Mn2/3]O2 (NFMO) powder was synthesized by a solid-state thermal reaction at 900ºC for 12h. The precursor was a mixture of Na2CO3 : Mn2O3 : Fe2O3 with a molecular ratio of 2.0 : 1.0 : 2.0. Before being annealed, the mixture was dispersed in ethanol, ground in a planetary ball mill for 24 h, and then dried at 70°C in air for 72 h. After annealing, the obtained NFMO powder was cooled to room temperature and ground prior to the preparation of the cathodes in the SIBs. Experimental cathodes were prepared by mixing NFMO powder with Super P (Timcal Ltd.) and VGCF (diameter 150 nm, length 5–10 µm, Showa Denko) conductive additives, and with polyvinylidene difluoride binder (PVDF, W1700, Kureha Battery Materials Japan Co., Ltd) in a weight ratio of 84 : 7 : 3 : 6 in a solvent of n-methyl-2-pyrrolidinone (NMP, International Specialty Products Inc.). The hard carbon electrode was made of 91.9wt% hard carbon (SF grades, Kureha Battery Materials Japan Co., Ltd.), 2wt% Super P, 1 wt% VGCF, and a 5wt% PVDF (W9100, Kureha Battery Materials Japan Co., Ltd) binder dissolved in NMP and 0.1wt% oxalic acid. The resulting slurries were coated by tape-casting on aluminium foil and copper foil to fabricate the NFMO and hard carbon electrodes, respectively, before being dried at 100°C in a vacuum for 1 h to evaporate the NMP. The electrodes were stored in a glove box with oxygen and humidity content maintained below 1 ppm for more than 24 h before electrochemical characterization. The experimental SIBs were assembled in an argon-filled glove box with oxygen and humidity content below 1 ppm. In CR2032 coin cell SIBs, the cathode was an NFMO electrode and the anode was sodium foil (99.99%, Sigma-Aldrich Fine Chemicals). In 18650 full cell SIBs, the anode was a hard carbon electrode and the cathode was an NFMO electrode. In both types of SIB, the electrolyte was a mixture of ethylene carbonate and diethyl carbonate (EC/DEC, 1:1 by weight) containing 1.0 M of NaPF6. To increase the packing density and reduce the internal contact resistance in the coin cells, the NFMO electrode cast on Al foil was roll-pressed prior to cell assembly. This is the reason why the capacity of the NFMO electrode was higher in the coin cells than in the 18650 full cells. For preparation of the NFMO cathode in an 18650 full cell, NFMO thin film was cast on Al foil with a length of 720 5



mm and a width of 70 mm. After the casting, the electrode was cut to a size of 480 x 50 mm to fit into an 18650 cell. In this step, roll press treatment was omitted due to the uniformity concern of pressure loading over such a large area. In the 18650 full cells, the mass ratios were NFMO : SP = 92 : 8 in the cathode and hard carbon : SP = 99 : 1 (without counting binders) in the anode, respectively. The powder of the NFMO sample was screened by a 212 µm mesh (ZSICSA.-212) prior to slurry preparation. In an optimized cell, the mass loadings of active materials (including NFMO and super P carbon = 140 mAhg-1) were 9.41 mg cm-2 at the cathode (hard carbon “HC” 270 mAhg-1) and 4.89 mg cm-2 at the anode, corresponding to an N/P ratio of 1.003. The charge/discharge behaviors of coin cells operated between 1.8 to 3.8 V and 1.8 to 4.3 V were evaluated at 0.1 C. The 18650 cells were formed at 0.1 C. For 18650 full cells, the cycling performance was tested at 0.2 C with cut-off voltages from 2.0 to 3.8 V and 2.0 to 4.5 V. Cycling tests were conducted at 25°C for 100 cycles. The crystal structure of the NFMO cathode was characterized by X-ray diffraction (XRD) with a commercial diffractometer (BRUKER D8 Advance ECO) using Cu Kα radiation (λ = 0.15418 nm). The local atomic and electronic structures of the NFMO cathodes were determined by X-ray absorption spectroscopy (XAS). Experimental XAS spectra were recorded at the tender X-ray beamline BL-16A1 (for the Mn K-edge at 6539 eV and Fe K-edge at 7112 eV) of the National Synchrotron Radiation Research Center (NSRRC) using the fluorescence detection mode with a Lytle detector. The incident and transmitted X-rays passing through the reference sample were simultaneously recorded in each XAS scan for calibration of the energy of the incident X-rays. Details of the structural parameter determination are provided in the ESI. For complementary discussions, both XRD and XAS characterizations were conducted on NFMO cathodes in selected charge/discharge states. In this study, the failure mode of the capacity of the NFMO cathode in a SIB was systematically interpreted by cross-referencing the results of structural inspections, ranging from microscopic to electronic spatial regimes. The electrochemical chemical properties of the experimental SIB cells revealed the occurrence of cycle fading during operation between 2.0 and 4.5 V. Focused ion 6


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beam section assisted scanning electron microscopy uncovered the formation of cracks due to volume differences as a result of Na+ intercalation and de-intercalation. X-ray diffraction revealed that the volume differences could be attributed to both the formation of ion channels and the phase transition between the P2 and O3 phases of NFMO. The results of X-ray absorption spectroscopy confirmed that changes in the symmetry of Fe sites were the likely initiating factor in the phase transition and ion channel formation. An irreversible chemical state was found at Fe sites in NFMO after long-cycle tests. During these tests, interactions with additives in the electrolyte led to dissolution of the Fe sites, accompanied by the collapse of neighboring Mn sites. These phenomena were consistently revealed by the formation of Fe and Mn oxide species in the hard carbon anode (i.e., the counterpart electrode), according to the results of the XPS depth profile composition.

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Results and Discussion Electrochemical performances of NFMO cathodes in a SIB coin cell operated at 2.0 to 3.8 V and 2.0 to 4.5 V in the formation stage To reveal the intercalation and de-intercalation of Na+ ions in NFMO crystal, polarization curves of SIBs operating between 2.0 to 3.8 V and 2.0 to 4.5 V were examined. Figure 1 compares the polarization curves of coin cells for the first two charge–discharge cycles. In Figure 1a (2.0 to 3.8 V), the potential plateau is insignificant in the first charge process to 3.8 V, suggesting that the de-intercalation of Na+ ions was mostly contributed from the near-surface region of the NFMO. In this event, as consistently revealed by XRD patterns of NFMO cathode at 3.8 V (Figure 3a), high numbers of Na+ ions were relocated to ion channels, leading to a low capacity of 14 mAh g-1 for the SIB coin cell. In the first discharge process to 1.8V, a rapid voltage drop resulted from intercalation of Na+ ions to the NFMO surface. When the potential was decreased from 3.3 to 2.3 V (region A), a linear capacity decay indicated an increasing energy barrier for Na+ intercalation upon the formation of ion channels in the NFMO crystal. This phenomenon can be attributed to increasing lattice compression in the opposite direction to lattice strain, generated by Na+ intercalation in the layered lattice planes of an NFMO. The plateau at 2.3 V indicated a low barrier for Na+ ion intercalation in NFMO crystal. Such a characteristic suggests the formation of ion channels by the positioning of sufficient amounts of Na+ between the layer lattice planes. The subsequent potential drop occurred when the coin cell discharged at higher than ~45 mAh g-1. In this region (region B), the potential decay slope was lower than that of region A. Such a feature indicates lower energy barriers for the formation and intercalation of Na+ ion storage sites than for those of the ion channel. In the second charge process, the linearly increased capacity (from 0 to ~76 mAh g-1) with potential from 2.0 to 3.3 V corresponded to ion de-intercalation from the storage site and the pseudo-steady channel built in the first discharge state. Increasing the potential above 3.3 V dramatically increased the slope in the polarization curve, suggesting a further increased Na+ ion de-intercalation barrier. This barrier developed due to the local

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distortion and an amorphous structure caused by the high amounts of de-intercalating Na+ ions from storage sites and channels. The polarization curve for a SIB coin cell operated at 2.0 to 4.5 V is presented in Figure 1b. Accordingly, the capacity–potential profile was identical to that of Figure 1a when the capacity was increased to 14.0 mAh g-1 with the potential increased from 2.0 V to 3.8V. The capacity–potential slope increased substantially up to 50 mAh g-1, indicating a further increased energy barrier for de-intercalation of Na+ ions from 3.8 to 4.3 V. After that, the applied potential exponentially increased and saturated at 4.5 V when the capacity was increased to ~68 mAh g-1. The symmetric profiles on the two sides of the reflection point Y refer to a typical diffusion control profile in crystal materials containing a low density of vacancies, therefore implying the formation of ion channels with proper alignments for ion–vacancy exchange. As compared to the capacity of Na+ ion intercalation in the first discharge process, the lower output capacity for de-intercalation of Na+ ions for both SIB coin cells can be explained by the formation of ion diffusion channels and certain residual Na+ ions in the storage site in the first charge process. Such a rationale is systematically explained by cross-referencing the results of physical structure inspections in later sessions.

Figure 1 Potential-capacity curves of a Na2/3[Fe1/3Mn2/3]O2/Na SIB coin cell with 1 M NaPF6 EC-DEC (1:1) electrolyte in the 1st charge/discharge and 2nd charge cycles (formation) in potential ranges of (a) 2.0 to 3.8 V and (b) 2.0 to 4.5 V at 0.1 C rate

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and 25ºC. Numbers in rad spheres refer to selected states of the SIB for XRD inspections. Electrochemical performance and corresponding microstructure inspections of NFMO cathode in SIB 18650 cells operated at 2.0 to 3.8 V and 2.0 to 4.5 V in long-cycle tests The discharge capacity of SIB 18650 cells (NFMO/separator/hard carbon) in long-cycle tests at potentials from 2.0 to 3.8 V and 2.0 to 4.5 V are compared in Figure 2a. As can be seen, in the first cycle, the discharge capacities were ~60.0 mAh g-1 for the 2.0 to 3.8 V cell and 144 mAh g-1 for the 2.0 to 4.5 V cell. For the SIB 18650 cell operated at 2.0 to 3.8 V, the capacity remained at ~60 mAhg-1 until the 100th cycle. On the other hand, the capacity of the SIB 18650 cell exponentially decreased to 61.5 mAh g-1 until the 100th cycle when operated between 2.0 to 4.5 V. The electrochemical properties of 18650 cells demonstrated responses similar to those of coin cells, which can be explained by the structural evolution of NFMO electrodes in selected states. Figure S1a in the electronic supplementary information (ESI) presents a top-view SEM image of a freshly-prepared NFMO cathode as a reference for distinguishing the microstructure evolutions of the experimental samples. As clearly indicated, the NFMO crystals were grown in hexagonal nano-disk of ~2.0–3.5 µm diameter and 120–150 nm thickness. Those nano-disks were stacked in parallel into a particle with a thickness of ~1.5 µm and showed a typical feature of NFMO in P2 phase. Figure 2b shows a SEM image of an NFMO electrode operated at 2.0 to 4.5 V after the first cycle. In this figure, the sharp nano-needles (denoted by yellow arrows) are precipitates of sodium carbonate crystal, and the nano-disks of 80–150 nm thickness are NFMO crystal. As denoted by the hollow spheres with yellow dashed lines, significant melting of the nano-disk indicated the dissolution and subsequent recrystallization of the NFMO crystal. In a SIB operated at 2.0 to 3.8 V for 100 cycles (Figure 2c), all nano-disks turned to condensed solid particles (denoted by the cross-section image of the FIB etching region) of several micrometers in size. The solid region in bulk indicates a homogeneous Na+ ion intercalation in the NFMO crystal. On the other hand, in a SIB operated at 2.0 to 4.5 V, the presence of cracks in the shell region and distinct white region in the core (denoted by the cross section 10


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image of the FIB etching region in Figure 2d) strongly indicated inhomogeneous Na+ ion intercalation/de-intercalation. Those uneven structure distributions and cracks clearly depicted two fatal issues for NFMO crystal as the cathode in a SIB when the operating voltage is higher than 3.8 volts. They are (1) a strong residual stress upon Na+ ion intercalation/de-intercalation in P2 crystal, and (2) an imbalance between the rates of Na+ ion intercalation/de-intercalation and phase transformation, and between the volume differences of the two phases (P2 and O3).15-16 Those hypotheses are further discussed with reference to analysis of the capacity decay rates of a SIB with NFMO cathodes in a long-cycle test (Figure S1b).

Figure 2 Discharge capacity in (a) linear plot versus cycle number of Na2/3[Fe1/3Mn2/3]O2/hard carbon 18650 full cell SIB cycled in different voltage ranges: 2.0 to 3.8 V and 2.0 to 4.5 V with 1 M NaPF6 EC-DEC (1:1) electrolyte at 25ºC and 0.1 C rate. SEM images of Na2/3[Fe1/3Mn2/3]O2 cathodes in SIB 18650 cells (b) operated at 2.0 to 4.5 V, after the first cycle, (c) operated at 2.0 to 3.8 V, after 100 cycles, and (d) operated at 2.0 to 4.5 V, after 100 cycles. 11



Crystal structure evolutions of NFMO cathodes in post-formation SIB coin cells at different voltages and SIB 18650 cells in charge-discharge cycles ranging from 2.0 to 3.8 V and 2.0 to 4.5 V XRD analysis provided explanations of the crystal structure evolutions of the NFMO electrochemical in selected electrochemical states. Figure 3 compares the XRD pattern of an NFMO cathode prepared at selected states of coin cells with the polarization curves demonstrated in Figure 1. The XRD patterns of as-prepared and NFMO cathodes at selected states in a coin cell operated at 2.0 to 3.8 V are compared in Figure 3a. As denoted by the red solid lines at pattern 1, the as-prepared sample was P-2 phase NFMO (i.e., planar crystal layers as ion channels perpendicular to the c-axis). For an NFMO in the fully charged state at 3.8 V, the diffraction lines of pattern 2 remained in the same position and featured a slightly decreased peak intensity as compared to those of pattern 1. Those characteristics indicated the occurrence of structural distortion without significant phase transition in the NFMO crystal upon Na+ ion de-intercalation in the charge process. The substantially decreased peak d intensity and presence of a broad diffraction peak at 26.2º indicated the formation of lattice planes with a space of 3.341 Å and an average coherent length (Davg) of 9.4 nm (Table S1). Such a phenomenon indicates the formation of ion channels upon de-intercalation of Na+ and certain lattice strains and volume expansion in an NFMO.17 Consequently, a considerable amount of Na+ was relocated from the lattice framework (i.e., storage sites or lattice point) to ion channels and could not contribute to the SIB capacity in the first charge process. At the first fully discharged state at 2.0 V (pattern 3), suppression of P2-characteristic peaks (c and d) as compared to those of pattern 1 consistently revealed the changes in structural symmetry in both long-range order and local atomic structures in the NFMO, with high amounts of Na+ ions stored at Fe sites and ion channels. When the coin cell was charged to 3.8 V in the 2nd charge process, the positions and intensities of the diffraction lines in pattern 4 again returned to those of pattern 1, demonstrating the reversibility of the NFMO structure in a coin cell operated at 2.0 to 3.8 V.

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The XRD patterns of the NFMO electrodes of SIB cells operated between 2.0 to 4.5 V are compared in Figure 3b. For the NFMO cathode in the first fully charged state at 4.5 V (pattern 5), the intensities of most of the characteristic diffraction lines in pattern 5 were compressed by more than 50% as compared to those of pattern 1, indicating a substantial Na+ ion relocation to ion channels that therefore disturbed the framework of the P2-NFMO crystal. Again, as discussed in pattern 2 of Figure 3a, relocated Na+ ions occupied the diffusion channels and were not able to (or could only slightly) contribute to the cell capacity in the first charge process. In the first fully discharged state at 2.0 V, the positions of the diffraction peaks in pattern 6 returned to those of the peaks in pattern 1. In this event, the substantially increased background intensity indicated the formation of a local distorted structure in NFMO crystal and possibly the formation of a solid electrolyte interface with high roughness. In the second fully charged state at 4.5 V, the absence of all P2 diffraction peaks in pattern 7 directly evidenced the formation of a long-range disordered structure and high Na+ ion retention in the diffusion channels, leading to the substantial capacity irreversibility of the NFMO in a long-cycle test. The correlations between the structure and capacity irreversibility of the NFMO crystal were further confirmed by comparing the XRD patterns of post long-cycle test cathodes and pristine NFMO cathodes in 18650 cells. Results are shown in Figure S2.

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Figure 3 XRD patterns of the prepared P2-Na2/3[Fe1/3Mn2/3]O2 cathode and electrodes in a SIB coin cell at selected states in (a) 2.0 to 3.8 V and (b) 2.0 to 4.5 V coin cells. In Figure 3a, patterns refer to NFMO in (1) as-prepared cathode, (2) the first fully charged state at 3.8V, (3) the first fully discharged state at 2.0V, and (4) the second fully charged state at 3.8 V. In Figure 3b, patterns refer to NFMO in (5) the first fully charged state at 4.5V, (6) the first fully discharged state at 2.0 V, and (7) the second fully charged state at 4.5 V.

Local structure evolutions of NFMO cathode in SIB coin cells and 18650 full cells The results on the crystal structure evolution and coin cell polarization curve illustrate that the capacity irreversibility was caused by stacking of Na ions at storage sites and ion channels in the inter-planar space in an NFMO crystal used as the cathode in a SIB. To further reveal the capacity fading mechanism in the local structure regime, Fe and Mn L-edge X-ray absorption spectroscopy analysis was employed. The results of X-ray absorption near-edge structure (XANES) spectra and extended X-ray absorption fine structure (EXAFE) revealed the evolution of electronic states and atomic structural distortion around Fe sites in the SIB coin cells in the high operation voltage range (1.8 to 4.3 V). The changes in the corresponding local structural parameters shine light on the capacity fading mechanisms in SIB 18650 full cells. Details of the quantitative structural interpretations are given in Figure S3 and S4. The stability of Mn and Fe sites in an NFMO cathode of a SIB 18650 full cell in a long-cycle test were confirmed by XAS analysis. Figure 4 compares the Fe and Mn K-edge XANES spectra of NFMO cathodes in 18650 cells operated at 2.0 to 3.8 V and 2.0 to 4.5 V in the 1st and 100th fully charged states (at 3.8 V or 4.5 V). Figure 4a compares the Fe K-edge XANES spectra of NFMO cathodes in the fully charged state at 3.8 V in the first (3.8V-1st) and 100th (3.8V-100th) cycles. Significant differences in the three XANES regions indicated that the local structure around the Fe atoms completely changed. In region A (details in inset), two quadrupole transition (e.g., t2g) peaks merged, revealing the transformation of Fe sites from Oh to Td symmetry 14


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upon de-intercalation and intercalation of Na+ ions in NFMO crystal.18-22 This merging illustrates a loss of the Fe-O coordination number (CNFe-O) due to breaking of the Fe-O bonds in the nearest atomic shell upon intercalation and de-intercalation of Na+ ions at Fe sites. The formation and dissolution of NaFeOx compounds is a possible pathway for Fe-O bond breaking and was consistently revealed by XPS in later sessions. This breakage distorted or collapsed the ion channels in the NFMO crystal, thus causing capacity irreversibility in the SIB. Changes in the near-edge profiles (region B) consistently depicted differences in the local structure symmetry at the Fe sites in the NFMO before and after the 100-cycle test of the 18650 cell. The reduction of the peak C intensity suggested an increasing electron density in the 3p-4s orbital, possibly due to substantial retention of Na+ ions (which trap consideration amounts of electrons in Fe-Na coordinates) at Fe storage sites. On the other hand, the Mn K-edge spectrum in region A (Figure 4b and inset) retained its original profile, implying the absence of local structural distortion at Mn sites after 100 cycles. A substantially enhanced pre-edge peak intensity indicated a strong chemical shift with a high density of electron vacancies in the 3d orbital. Such features are indications of a strong affinity of Na+ ions to Mn sites. However, it is important to note that the position and profile in the pre-edge region remained unchanged in the Mn K-edge after 100 cycles, showing that the symmetry of Mn sites was not affected by the intercalation and de-intercalation of Na+ ions and again confirming their stronger stability as compared to that of Fe sites in an NFMO cathode. In the 2.0 to 4.5 V cell (Figure 4c), the chemical shift in the 3d orbital and the local distortion around the Fe atoms were larger than those in the 2.0 to 3.8 V cell, as denoted by a larger positive deviation in the spectrum intensity in the pre-edge region after 100 cycles. The overlapped peak C intensity suggested a similar extent of the empty state in the 3p-4s orbital between the Fe atoms in NFMO before and after the long-cycle test. Such a phenomenon is understandable, considering that in the first cycle, high Na+ ion retention was found in the channels and storage sites. Notice that the chemical state of Mn sites in the 2.0 to 4.5 V cell was similar to those of the 2.0 to 3.8 V cell, except that the density of the empty state in the 3d orbital (pre-edge peaks in Figure 4d) was lower due to the greater extent of Na+ retention. Meanwhile, as compared to Figure 4b, a smaller decrease in the absorption peak intensity 15



consistently revealed the higher Na+ ion retention as compared to that of an 18650 cell operated at 2.0 to 3.8 V for 100 cycles.

Figure 4 X-ray absorption near-edge structure (XANES) spectra of NFMO cathode after the 1st and 100th cycles of the C-DC process in a SIB 18650 cell at (a) Fe K-edge and 2.0 to 3.8 V, (b) Mn K-edge and 2.0 to 3.8 V, (c) Fe K-edge and 2.0 to 4.5 V, and (d) Mn K-edge and 2.0 to 4.5 V.

The atomic structure evolutions of the Fe and Mn sites in the NFMO cathodes of 18650 cells in long-cycle tests were further elucidated by EXAFS analysis. Figure 5a compares the Fourier-transformed EXAFS spectra (radial structure function, RSF) of as-prepared NFMO and that of post-formation and long-cycle test cathodes in a 2.0 to 3.8 V 18650 cell at the Fe K-edge. The two radial peaks A and B in the RSF refer to contributions of Fe-O and Fe-Fe/Fe-Mn bond-pairs at the first and second nearest coordination shells around Fe atoms. The significant offset of peaks A and B to the high R side suggested an increment of interatomic distance (i.e., local structure expansion) between Fe and neighboring atoms (O, Fe, and Mn) in the 100th cycles. Such an expansion indicated the restructuring of Fe sites due to high amounts of Na+ 16


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ions intercalating in their neighbors. Meanwhile, the substantially enhanced peak B revealed an increasing coordination number (CN) in the second nearest shell of Fe after the first cycle. This CN increment was an indicator of local structure ordering at Fe sites due to the formation of Na+ ion channels. A lower radial peak intensity means a considerable extent of short-range disordering of the structure, in the form of either vacancies or local distortion in lattice planes, in as-prepared P-2 NFMO crystal. For the RSF of NFMO cathode after 100 cycles, the peak B intensity was dramatically reduced, indicating substantial local disordering around the Fe atoms. As consistently revealed by XANES analysis (Figure 4a), such a highly disordered structure was attributed to the transformation of Fe sites from Oh to Td symmetry or the formation of NaFeOx compounds upon Na+ intercalation and de-intercalation at Fe sites. For the case of a SIB 18650 cell operated at 2.0 to 4.5 V, the local structure evolutions of Fe sites were similar to those of a cell operated at 2.0 to 3.8 V. Some minor differences were caused by the larger extent of structural disorder, revealed by the lower intensities of peaks A and B as compared to those of RSF of as-prepared NFMO. The former indicated the larger steric barrier for packing Na+ ions around Fe atoms, and the latter implied the lower extent of lattice alignment due to the high Na+ ion stacking in ion channels (i.e., a large extent of capacity irreversibility). The Mn K-edge RSF spectra of an as-prepared NFMO electrode and those in SIB 18650 cells operated at 2.0 to 3.8 V for one and for 100 cycles are demonstrated in Figure 4c. As compared to those of the as-prepared NFMO, the intensities of the two radial peaks (A and B) were increased by the same extent in the first fully charged state at 3.8 V. In the absence of local structural distortion proved by XANES analysis, the enhancement of radial peak intensity suggested an increasing long-range ordering between Mn sites upon the formation of ion channels by intercalation and de-intercalation of Na+. In this event, the Mn sites might have mainly aligned along the Na+ ion channels. When the cycle number was increased to 100#, the intensity and position of peak A remained unchanged, suggesting that a sufficient amount of Na+ was stacked in storage sites and the local structure around Mn held. In this event, the substantially reduced intensity in peak B can be attributed to local disordering due to de-intercalation of Na+ from ion channels. For an 18650 cell operated at 2.0 to 4.5 V, the profiles of all the RSF spectra were identical, regardless of the number of charge– 17



discharge cycles. This means that the local chemical environment and symmetry of the Mn sites was not affected by the intercalation and de-intercalation of Na+ ions in their neighbors. Therefore, in the absence of local environment evolution around Mn sites, one can notice that the intercalation and de-intercalation and Na+ ions occurred mainly in the ion channels and particularly around Fe sites. This is also the explanation for the overloading and thus distortion or even dissolution fading of Fe sites in NFMO cathodes in a SIB 18650 cell. By cross-referencing the structural evolution results mentioned above, one can infer the following scenarios for a structure that would lead to capacity irreversibility of NFMO crystal in the cathodes of Na ion batteries. (1) As-prepared NFMO crystals contain certain defects and a short-range disordered structure; (2) Fe sites tend to transform from Oh to Td symmetry due to the chemical reaction of Na+ at their neighboring sites; (3) the local structure symmetry around Mn sites remains unchanged, regardless of the number of charge-discharge cycles; (4) the extent of chemical shift is higher at Mn sites than at Fe sites upon changes in the state of charge in a SIB; (5) the extent of core level chemical shift (i.e., Na+ ion affinity) is larger at Mn sites than at Fe sites in an NFMO crystal; (6) the structural distortion or transformation of Oh to Td symmetry of Fe sites occurs in the first charge–discharge cycle. The result is the dissolution of Fe and Mn atoms and thus suppression of the capacity reversibility of NFMO in a SIB. Those scenarios were further confirmed by in-operando neutron scattering and ex-situ XPS analysis in later sessions.

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Figure 5 Fourier transformed extended X-ray absorption fine structure spectra of NFMO cathode after the 1st and 100th cycles of the C-DC process in a SIB 18650 cell at (a) Fe K-edge and 2.0 to 3.8 V, (b) Fe K-edge and 2.0 to 4.5 V, (c) Mn K-edge and 2.0 to 3.8 V, and (d) Mn K-edge and 2.0 to 4.5 V.

The results of in-operando neutron diffraction further confirmed local and crystal structural distortion in NFMO upon Na+ ion de-intercalation in the charge stage of the first cycle. As shown in Figure 6 and Figure S5, the lattice space of (002) (d(200)) expanded from 4.755 to 4.806 Å with cell voltages of 2.0 to ~3.9 V. With a further increase in voltage to 4.3 V, d(200) remained at 4.806 Å. A plateau of the lattice space between 3.9 to 4.3 V implied a high energy barrier for Na+ ion diffusion. As consistently revealed by the results of XAS analysis and polarization curve in Figure 1, such a diffusion barrier rationalizes a considerable density of defect sites in the NFMO crystal. In the charge process, the polarization curve showed a similar profile to that of diffusion-controlled ion concentration distribution in a solid state crystal with the capacity increasing from 0 to 300 mAh. In this curve, an inflection point for 19



the different slopes of capacity to voltage plot was found, implying an interface for the ion diffusion barrier at 3.9V. The substantially reduced capacity to voltage slope revealed a slow Na+ to defect exchange region when the voltage was higher than 3.9 V. In this voltage range, a slight decrease of d(102) from 2.420 to 2.418 Å resembled local compression/distortion in lattice planes intersecting the (002) plane. As consistently elucidated by XAS analysis; such expansion could be attributed to local distortion occurring mostly at the Fe sites. With an increase in cell voltage to 4.5 V, d(002) expanded to 4.825 Å. In this stage, d(102) shifted back to its original width, revealing the relaxation of local distortion upon completion of the ion diffusion channels in the NFMO. By cross-referencing the results of XAS, XRD, and in-operando neutron diffraction, the scenarios for Na+ de-intercalation in NFMO crystal with respect to voltage changes in a SIB charge process include the following: (1) From 2.0 to 3.9 V, Na+ ions relocate from lattice points to ion channels. In this stage, Na+ ions tend to accumulate in lattice planes perpendicular to the c axis, as revealed by the expansion of d(002) from 4.755 to 4.806 Å. In this event, Na+ intercalation generates a lateral expansion in the (002) plane and thus compresses the lattice spaces of intersecting lattice planes (i.e., (102)); (2) From 3.9 V to 4.3 V, the capacity increases from 133 to 267 mAh. In the absence of d(000) expansion, Na+ relocates without generating additional distortion in (002) planes and is mainly in ion channels. Further increasing the voltage from 4.3 to 4.5 V slightly expands d(002) to 4.825 Å, which can be accounted for by further Na+ intercalation and structure alignment in ion channels. In this event, the expansion of d(102) (i.e., relaxation of lateral strain in ion channels) to its original value complementarily explains the interplay of lattice strains between intersecting lattice planes from 3.9 to 4.5 V.

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Figure 6 In-operando neutron diffraction results and the corresponding polarization curve of an 18650 cell operating between 2.0 to 4.5 volts in the charge process.

Hypothetical pathways for the Fe and Mn restructuring were further confirmed by the XPS depth profiles of the hard carbon (HC) anode in 18650 cells operating at 2.0 to 3.8 and 2.0 to 4.5 V after long-cycle tests. The XPS depth profiles of two 18650 cells are compared in Figure 7. As can be seen, the atomic ratios were ~2.0 +/- 0.2% for Fe and 1.7 to 7.9 +/- 0.3% for Mn with the sputter time increased from 0 to 20 mins for the HC anode of the 2.0 to 4.5 V cell. On the other hand, residual Fe and Mn species were absent in the HC electrode of the 2.0 to 3.8 V cell. Those phenomena rationalized the formation of amorphous NaFeOx clusters upon Na+ de-intercalation from Fe sites (Scheme 1). In this case, the Fe sites turned from Oh to Td symmetry and possibly formed amorphous NaFeOx in the NFMO cathode. By interacting with intermediate species of PF additives, the two structures would dissolve in electrolyte, leading to plating of FeOx and MnOx species on the HC electrode and thereby causing capacity fading of the NFMO cathode in a SIB. 21



Figure 7 Atomic concentrations of Mn and Fe on the hard carbon anode in Na2/3[Fe1/3Mn2/3]O2/hard carbon SIB 18650 cell after 100 cycles in voltage ranges of 2.0 to 3.8 V and 2.0 to 4.5 V, respectively, with 1 M NaPF6 EC-DEC (1:1) electrolyte at 25ºC and 0.1 C rate.

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Scheme 1 Fading pathways of NFMO crystal in a SIB cell. Intercalation of Na+ in storage sites induces two reactions: (1) ion exchange between Na+ and vacancy (∆), and (2) formation of local FeOx and MnOx compounds in NFMO crystal. Diffusion of Na+ induces lateral strains perpendicular to the ion channels (ab-plane). FeOx and MnOx in the NFMO surface will dissolve in electrolyte and then plate the hard carbon anode. In NFMO bulk, FeOx and MnOx will form local amorphous clusters and block the ion channels. Both pathways cause the irreversibility of the NFMO cathode in SIBs.

23



Conclusion P2-Na2/3[Fe1/3Mn2/3]O2 crystal with a maximum capacity of ~150 mAh g-1 as a cathode in a sodium ion battery was synthesized by a solid state annealing method. By employing ex-situ X-ray absorption spectroscopy analysis, we have demonstrated that the capacity irreversibility of such a cathode in a SIB is caused by the transformation of Fe sites from octahedral (Oh) to tetragonal (Td) symmetry, which occurs due to interactions with Na+ ions in charge and discharge processes. Meanwhile, Mn sites show higher affinity for Na+ storage than do Fe sites, and their local symmetry remains unchanged even in a long-cycle test of 100 cycles. XPS depth profile analysis revealed plating of Fe on the hard carbon anode. By cross-referencing experimental evidence, we found that the Oh to Td transformation can be ascribed to the breakage of Fe-O bonds due to the interaction between the Fe and neighboring Na+ ions, with the subsequent formation of NaFeOx and dissolution of Fe and Mn in the electrolyte. Such a phenomenon collapses the NFMO crystal and thereby suppresses its capacity reversibility in a SIB.

Supporting information Supporting Information Available: SEM image of pristine NFMO crystal, XRD patterns of NFMO electrodes, Fe and Mn K edges XAS analysis of NFMO electrodes in selected charge stages of SIB, in-operando neutron diffraction spectra of NFMO electrode in 18650 cell, charge-discharge curve of hard carbon in SIB, and XPS depth profile of experimental NFMO electrodes are provided.

Acknowledge The authors are grateful for the assistance of the staff of the National Synchrotron Radiation Research Center (NSRRC), Hsinchu, Taiwan, in conducting various synchrotron-based measurements. T.-Y. Chen and C.-C. Chang acknowledge the 24


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funding support from the Ministry of Science and Technology, Taiwan (MOST 106-2112-M-007-016-MY3, 105-2221-E-006-189-MY3,

MOST MOST

105-3113-E-006-019-CC2, 106-3113-E-006-003-CC2,

106-2912-I-024-502).

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and

MOST MOST


Reference 1.

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Q.; Chen, L. Q.; Huang, X. J., A Zero-Strain Layered Metal Oxide as the Negative Electrode for Long-Life Sodium-ion Batteries. Nat. Commun. 2013, 4, 7. 10. Yabuuchi, N.; Komaba, S., Recent Research Progress on Iron- and Manganese-based Positive Electrode Materials for Rechargeable Sodium Batteries. Sci. Technol. Adv. Mater. 2014, 15 (4), 29. 11. Clement, R. J.; Bruce, P. G.; Grey, C. P., Review-Manganese-Based P2-Type Transition Metal Oxides as Sodium-Ion Battery Cathode Materials. J Electrochem Soc. 2015, 162 (14), A2589-A2604. 12. Wang, H.; Yang, B.; Liao, X.-Z.; Xu, J.; Yang, D.; He, Y.-S.; Ma, Z.-F., Electrochemical Properties of P2-Na2/3[Ni1/3Mn2/3]O2 Cathode Material for Sodium Ion Batteries when Cycled in Differentvoltage Ranges. Electrochim Acta 2013, 113 200-204. 13. Yabuuchi, N.; Kajiyama, M.; Iwatate, J.; Nishikawa, H.; Hitomi, S.; Okuyama, R.; Usui, R.; Yamada, Y.; Komaba, S., P2-type Nax[Fe1/2Mn1/2]O2 Made from Earth-abundant Elements for Rechargeable Na Batteries. Nat. Mater. 2012, 11, 512-517.

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14. Gonzalo, E.; Han, M. H.; Amo, J. M. L. o. d.; Acebedo, B.; Casas-Cabanasa, M.; Rojo, T., Synthesis and Characterization of Pure P2- and O3-Na2/3Fe2/3Mn1/3O2 as Cathode Materials for Na Ion Batteries. J. Mater. Chem. A 2014, 2, 18523-18530. 15. Sun, H.-H.; Manthiram, A., Impact of Microcrack Generation and Surface Degradation on a Nickel-Rich Layered Li[Ni0.9Co0.05Mn0.05]O2 Cathode for Lithium-Ion Batteries. Chem. Mater. 2017, 29 (19), 8486-8493. 16. Yoon, C. S.; Jun, D.-W.; Myung, S.-T.; Sun, Y.-K., Structural Stability of LiNiO2 Cycled above 4.2 V. ACS Energy Lett. 2017, 2 (5), 1150-1155. 17. Kondrakov, A. O.; Schmidt, A.; Xu, J.; Geßwein, H.; Mönig, R.; Hartmann, P.; Sommer, H.; Brezesinski, T.; Janek, J., Anisotropic Lattice Strain and Mechanical Degradation of High- and Low-Nickel NCM Cathode Materials for Li-Ion Batteries. J. Phys. Chem. C 2017, 121 (6), 3286-3294. 18. Oliveira, M. M.; Schnitzler, D. C.; Zarbin, A. J. G., (Ti,Sn)O2 Mixed Oxides Nanoparticles Obtained by the Sol-gel Route. Chem. Mater. 2003, 15 (9), 1903-1909. 19. Parlebas, J. C.; Khan, M. A.; Uozumi, T.; Okada, K.; Kotani, A., Theory of Many-body Effects in Valence, Core-level and Isochromat Spectroscopies along the 3d Transition-metal Series of Oxides. J. Electron. Spectrosc. Relat. Phenom. 1995, 71 (2), 117-139. 20. Wu, Z. Y.; Zhang, J.; Ibrahim, K.; Xian, D. C.; Li, G.; Tao, Y.; Hu, T. D.; Bellucci, S.; Marcelli, A.; Zhang, Q. H.; Gao, L.; Chen, Z. Z., Structural Determination of Titanium-oxide Nanoparticles by X-ray Absorption Spectroscopy. Appl. Phys. Lett. 2002, 80 (16), 2973-2975. 21. Newville, M., IFEFFIT: Interactive XAFS Analysis and FEFF Fitting. J Synchrotron Radiat. 2001, 8, 322-324. 22. Durham, P. J., X-Ray Absorption: Principles, Applications, Techniques of EXAFS, SEXAFS and XANES (Chemical Analysis: A Series of Monographs on Analytical Chemistry and Its Applications). John Wiley & Sons: New York, 1998; p 53-81.

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Table of contents graphic

Caption: Distortion followed by bond braking induces Fe dissolution and consequently capacity fading of NFMO crystal as cathode in SIB.

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Local chemical reaction of Na+ to O transforms Fe sites from octahedral to tetragonal symmetry in Na[Fe1/3M]O2 crystal. In surface, such reaction dissolves Fe and neighbouring Mn to chelate with CHn in electrolyte and then plat in graphite surface. In bulk, amorphous NaFeOx is formed to block ion channels. Both the two pathways induce capacity irreversibility of a cathode in Na ion battery. 206x175mm (150 x 150 DPI)


X-ray Absorption Spectroscopy and In-Operando Neutron Diffraction

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