The Study of Higher Discharge Capacity, Phase Transition and

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The Study of Higher Discharge Capacity, Phase Transition and Relative Structural Stability in Li2FeSiO4 Cathode upon Lithium Extraction using Experimental and Theoretical Approach and Full cell Prototype Study Shivani Singh, Anish Raj K, Manas Ranjan Panda, Raja Sen, Priya Johari, Anil K Sinha, Sher Singh Meena, and Sagar Mitra ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.9b01145 • Publication Date (Web): 08 Aug 2019 Downloaded from pubs.acs.org on August 11, 2019

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ACS Applied Energy Materials

The Study of Higher Discharge Capacity, Phase Transition and Relative Structural Stability in Li2FeSiO4 Cathode upon Lithium Extraction using Experimental and Theoretical Approach and Full Cell Prototype Study

Shivani Singh†, Anish Raj K† Manas Ranjan Panda†,‡, Raja Sen#, Priya Johari#,*, A. K. Sinha┴, Sher Singh Meena§ and Sagar Mitra†,*

†Electrochemical

Energy Laboratory, Department of Energy Science and Engineering, Indian

Institute of Technology, Bombay, Powai, Mumbai 400076, India ‡IITB #Department

Monash Research Academy, Bombay, Mumbai-400076, India

of Physics, School of Natural Sciences, Shiv Nadar University, Gautam Budha Nagar, Greater Noida, Uttar Pradesh-201314, India.

┴Indus

Synchrotron Utilization Division, Raja Ramanna Centre for Advanced Technology, Indore-452013, India

§Solid

State Physics Division, Bhabha Atomic Research Centre, Mumbai, 400085, India

Corresponding Authors

*Email: [email protected] (Priya Johari) *Email: [email protected] (Sagar Mitra)

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ABSTRACT: We have revisited the study of nanostructured lithium iron silicate-based cathode for safe lithiumion battery and to understand the reaction mechanism from first cycle to second cycle. Ex situ Mössbauer and X-ray absorption near edge structure spectroscopy (XANES) measurements have been carried out on electrodes charged at various voltages to investigate the electrochemical activity of Fe3+/Fe4+ redox couple to confirm the existence of Fe4+ and its role in defining the structural and electrochemical properties. The first charge and discharge lead to a structural change, which results in potential plateau shift after the first charge. To validate this understanding; ex situ Synchrotron X-ray diffraction (SXRD) along with Rietveld refinement results and first-principles density functional theory-based analysis have been performed, which also support the change in the crystal structure of the material with cycling. The in situ electrochemical impedance spectroscopy demonstrates phase transformation in de-lithiated iron silicate as lithium concentration changes during the charging process, which has been correlated with change in the density of states calculated by density functional theory. Finally, a full-cell prototype has been demonstrated first time by using lithium iron silicate cathode as cathode and graphite as anode and the full-cell shown the capacity retention of 92% after 50 cycles at 1C rate. KEYWORDS: cathode material; Li2FeSiO4, phase transition, SXRD, XANES, density functional theory, lithium-ion batteries


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1. INTRODUCTION Multi-electron charge transfers in polyanionic-based silicates such as Li2FeSiO4 leads to higher theoretical capacity, ensuing in nearly double discharge capacity, i.e., 333 mA h g-1, related to commercially existing substitutions.1 This material is greener, safer and low cost due to the abundance of Fe and Si in the earth’s crust.2 Li2FeSiO4 can be prepared in three different space groups (Pmn21,

3

Pmnb

4

(orthorhombic) and P21/n

1

(monoclinic)) by changing experimental

parameters viz. calcination time and temperature. In the crystal structure, entirely cations (Li+, Fe2+, Si4+) are tetrahedrally attached by oxygen in reported polymorphs.5–6 Once the cathode reacts with lithium in half-cell configuration, the first Li-ion de-intercalation for Fe2+/Fe3+ redox couple occurs roughly at 3.0 V, whereas Fe3+/Fe4+ redox couple reactivity occurs at around 4.75 V, in all the three polymorphs.2,5–9 It is well known and established for silicates in the literature that the value of potential plateau shifts from 3.1‒2.8 V after first charging cycle

3,9,10

and this lowering of

potential plateau after the first cycle is attributed to phase transformation in Li2FeSiO4. Several experimental and computational efforts have been made to understand the observation of shift in potential plateau after the first charge cycle in Li2FeSiO4.5,6,11,12 Earlier experimental observations on potential plateau shift have been assigned to structural rearrangement, i.e. cation mixing with Li and Fe. The creation of anti-site defect upon cycling is another reason proposed in the literature.13 Moreover, experimental literature reports the successful extraction of more than one lithium, resulting in greater than 166 mA h g-1 discharge capacity. It is well known that the obtained capacity mainly depends on the morphology, particle size, and synthesis route.1,10,12,14–16 Synthesizing material by solution method results in nanomaterial which helps in achieving more capacity because Li diffusion path length reduces substantially in the case of nanomaterials.17 As 3 ACS Paragon Plus Environment


demonstrated by D. Rangappa et.al. that by synthesizing ultra-thin sheets of Li2FeSiO4, successful extraction of two lithium at 45 ˚C up to 20 cycles is possible.10 In their work, the obtained capacity of ~340 mA h g-1 by full lithium de-insertion has been assigned to electrochemically active Fe2+/Fe3+ and Fe3+/Fe4+ redox couples. Such high discharge capacity was attributed to the electrochemical active Fe3+/Fe4+ redox couple. Few different mechanisms exist in the literature for second lithium extraction, i.e., the charge compensation beyond one lithium extraction. Masese et.al. have demonstrated that the formation of ligand hole in O-2p band is responsible for the extra capacity achieved during the second oxidation step for Li2FeSiO4, which compensates the charge after first lithium extraction.18 While Dongping Lv et. al. have assigned extra capacity to the electrochemical activity of Fe3+/Fe4+ redox couple by analyzing the ex situ Mössbauer spectroscopy and in situ X-ray absorption (XAS) experiments.12 Furthermore, theoretical studies using ab-initio density functional theory (DFT) have also been undertaken to understand the mechanism for second lithium reversibility.18 In short, broadly two mechanisms have been proposed in the literature to explain the extra capacity.12,18 Moreover, the capacity loss upon cycling in case of pure Li2FeSiO4 has also been reported in the previous literature,19–22 which might be because of very poor electronic/ionic conductivities19–22 or instability associated with the Fe4+ oxidation state of iron.23 But, the discrepancy regarding the cyclability of the second lithium exists in the literature. Therefore, in order to develop an understanding to explain the robust reason behind the extra capacity in the initial cycles as well as the phase transformation with a change in the lithium concentration upon cycling, we have undertaken a comprehensive study of lithium iron silicate, experimentally as well as computationally. In our earlier work, an initial discharge capacity of around 240 mA h g-1 has been observed, indicating the extraction of more than one lithium-ion.24 This extra capacity was attributed to Fe3+/Fe4+ redox couple. In present work, ex situ


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Mössbauer spectroscopy, SXRD and XANES measurements have been carried out on electrodes charged at various voltages viz. at OCV ~ 4.0V, 4.5V and 4.8V, to observe the structural changes associated with the varying electrochemical properties. Thus, this study mainly focuses on the experimentally synthesized Li2FeSiO4 and investigates the change in the crystal structure of Li2FeSiO4 as well as the change in the oxidation state of iron upon lithium extraction. We have also attempted to clarify the nature of phase transition between Li2FeSiO4/Li2-xFeSiO4 compounds (i.e., between lithium-rich and lithium poor phases of lithium iron silicate) using in situ electrochemical impedance spectroscopy, which has been correlated to the calculated total and partial density of states (DOS) for different lithium concentration. Additionally, we have demonstrated a full cell performance by combining Li2FeSiO4 (LFS) as cathode and graphite as anode. The LFS//graphite full-cell has shown the capacity retention of 92% after 50 cycles at 1C rate. To the best of our knowledge, the full cell using Li2FeSiO4 as cathode has not yet been studied in the available literature. 2. EXPERIMENTAL METHODS 2.1. Synthesis of Li2FeSiO4. Sol-gel process was used to synthesize nano Li2FeSiO4. Initially, P123 (Sigma-Aldrich) was fully dissolved in double distilled water and ethanol, subsequently, stoichiometric quantity of iron nitrate (98%, Merck-India) and lithium acetate (99.5%, SD fine chemical limited, India) were added to the solution. Then, the stoichiometric quantity of tetraethyl orthosilicate (TEOS-99%, Sigma-Aldrich) was also dissolved to this solution. The solution was dried at 50 °C for overnight to evaporate the water and ethanol, and then heated at 100 °C under vacuum. Finally, the dried gel precursor made into powder and calcined at 700 °C in N2/H2 (95%/5%) environment in order to



obtain pure Li2FeSiO4 phase. Li2FeSiO4/multi-walled carbon nanotubes (MWCNTs) composite was processed via solution method to increase the particle-particle connection. 2.2. Materials Characterization. The surface morphology and microstructure of the synthesized sample was viewed by FEG-SEM (JEOL-7600F) and HR-TEM (JEOL-2100F), respectively. The Fourier transformed infrared spectroscopy (3000 Hyperion Microscope with Vertex 80 FT-IR System Bruker, Germany) in ATR mode was carried out in the range of 400-2000 cm-1 to get the metal-oxygen bond information and analyze the variations in vibrational band intensity and frequency in the sample upon cycling. Room temperature Mössbauer spectra (MS) were recorded by a Mössbauer spectrometer functioned in constant acceleration mode (triangular wave). Co-57 was employed as source in Rh matrix of strength 50 mCi. The standardization of the velocity has been performed by means of an enriched α-57 Fe metal foil. The inner line width of the calibration spectrum is 0.23 mm/s. The recorded MS were fitted using a WIN-NORMOS site fit program. For all ex situ measurements, electrodes were dissembled from the Swagelok cell and washed properly with dimethyl carbonate (DMC, Sigma Aldrich) in an argon-filled glovebox (Mbraun, Germany) and dried well prior to experiments, ambient condition exposure was minimized as much as possible. The ex situ SXRD and XANES experiments were carried out on angle-dispersive X-ray diffraction (ADXRD) beamline (BL-12) at INDUS-2 Synchrotron Source (2.5 GeV, 120 mA), Raja Ramanna Centre for Advanced Technology (RRCAT), Indore, India.25, 26 The X-ray of wavelength (0.71184 Å) was used for SXRD measurement. XANES measurement was performed on LFS electrodes before cycling (i.e., as prepared electrode), and at different delithiated states at Fe K edge (7112 eV). XANES spectra were recorded in steps of 1 eV in fluorescence mode.25, 26


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For all electrochemical measurements, the slurry was made by adding Li2FeSiO4 active material (78%) with Super-P carbon (12%), PVDF binder (10%) and NMP-acetone mixture used as solvent. Then the well-mixed slurry was cast on an Al- foil used as current collector. The casted slurry was dried for 12h in under vacuum at 120 °C to eliminate solvents; circular disk of 12 mm electrode was cut to prepare the cell. The average active material loading was 1.6-2.3 mg cm-2. For the electrochemical performance of cathode, Swagelok cell was used to assemble in an Ar-packed glove-box (Mbraun, Germany). 1M LiPF6 salt in a (1:1) solution of ethylene carbonate (EC) and dimethyl carbonate (DMC) (LP-30, Merck Germany) was used as electrolyte. All electrochemical measurements were performed at constant temperature of 20 °C in the potential window of 1.54.8 V vs. Li/Li+ using Arbin BT-2000, USA test system. In situ electrochemical impedance spectroscopy (EIS) experiments were carried out during galvanostatic charge-discharge at various potentials at 20 ⁰C. The EIS at each intermediate potential was collected over frequencies ranging from 1 MHz‒50 mHz with Bio-logic VMP-3 by applying voltage perturbation of ΔV = 5 mV. The full-cell was constructed using nano Li2FeSiO4 as cathode and graphite as anode in CR-2032-coin cell using LP-30 as electrolyte while borosilicate glass fiber (Whatman/GF/D) was used as a separator. For the full-cell construction, the Li2FeSiO4 electrode (i.e., cathode) was used after two cycles at 0.1 C rate. The two cycles at 0.1 C rate were performed to activate and stabilize the electrode. The Galvanostatic charge-discharge for full-cell was carried out using Arbin instrument (BT-2000, USA) in the voltage window of 1.0‒4.5 V and 1.0‒4.0 V.

2.3. Computational details.



Structure optimization and total energy calculations of different LixFeSiO4 phases (with x =2, 1, 0.5, and 0) were carried out by first-principles spin-polarized density functional theory (DFT) as employed in the Vienna Ab-initio Simulation Package (VASP).27,28 The electron-ion interactions in the calculations were represented by projector-augmented-wave (PAW) pseudopotentials,29 where Li: 2s12p0, Fe: 3p63d74s1, Si: 3s23p2, and O: 2s22p4 electrons were treated as valence electrons, and all other electrons were treated as core. The generalized gradient approximation (GGA) parameterized by Perdew-Burke-Ernzerhof (PBE)30 was adopted for the description of the exchange-correlation of valence electrons. It is well known from earlier studies31 that cathode materials having transition metals exhibit strong electron correlation effects, therefore, demands an onsite Coulomb self-interaction correction for d-orbital. Thus, to account the strong correlation of the d-electrons of Fe, a Hubbard-like correction term32 (GGA+U) was included in the total energy functional by considering a J correction term of 1 eV and a U value of 5 eV, giving an effective value of 4 eV (Ueff = U-J).6 These values are well described in the literature and have been obtained by fitting the experimental oxidation enthalpies of Fe.33 They are found to estimate a reasonable bandgap and redox couple voltage for Li2FeSiO4/LiFeSiO4. In this study, starting from experimental lattice constants, all three polymorphs of Li2FeSiO4, i.e., Pmn21, Pnma, and P21/n are fully relaxed until the Helman–Feynman forces acting on atoms were < 0.01 eV/Å. The plane-wave kinetic energy limit was assumed as 500 eV, while the reciprocal space determination for k-points generation was set to 0.05×2Π Å-1 with uniform Γ-centered meshes. All the abovementioned parameters ensure that the total energy calculations are well converged with fluctuation in energy to be less than 1.0 meV/atom. A finer k-mesh with reciprocal space resolution of 0.02×2Π Å-1 was adopted for the calculations of the density of states (DOS) of pristine and deintercalated phases of Li2FeSiO4. A supercell of 4 formula units was considered for Pnma and


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P21/n, whereas 2 formula units were considered for Pmn21 based Li2FeSiO4. In order to imitate the de-intercalation process of Li2FeSiO4 and to understand the structural evolution from Li2FeSiO4 to FeSiO4 upon elimination of Li, we have systematically extracted Li from different (four) Wyckoff sites. In every case, the structure was allowed to relax completely. A schematic of the adopted approach is presented in Figure S1 of Supporting Information (SI). These optimized structures corresponding to various configurations for a given concentration were analyzed, primarily on the basis of ground-state energy. The configuration with minimum ground state energy for a given concentration represents the most stable structure. To analyze the electrochemical properties of Li2FeSiO4, the average de-intercalation voltage was calculated by using the equation given below.34

where, x2 and x1 represent the number of lithium atoms per formula unit present in the compositions before and after the de-intercalation process. However, the above equation was obtained by further simplifying Gibb’s free energy, on neglecting the volume and entropy terms in the equation G = E + PV – TS, as their magnitude are much smaller than the total energy.34 G(LixFeSiO4) and G(Li) are the Gibbs free energy of LixFeSiO4 and the Gibb’s free energy of Li (Li-Im-3m) per Li atom, respectively. Thus, we approximated G with the total energy (E) of each composition.

3. RESULTS AND DISCUSSION 9 ACS Paragon Plus Environment


The reversible extraction and insertion of more than one lithium into the host matrix results in higher capacity in lithium iron silicate. The impact of morphology, particle size, and preparation route on obtained capacity can be clearly seen in the reported literature.1,10,14,17 The theoretical calculation supports that the potential for Fe3+ to Fe4+ transition in iron silicates is above 4.7V vs. Li metal. Few experimental observations based on XANES and Mössbauer spectroscopy studies also support the fact that Fe4+ state can exist. The Mössbauer spectroscopy can well define the transition from Fe2+ to Fe3+ oxidation state and easily distinguish it, however, the conversion of Fe3+ to Fe4+ is quite difficult to monitor because of their close values. Therefore, the literature on higher capacity and associated redox process is under active surveillance and not many convincing results or justifications are yet available. Here, we have performed a series of ex situ surfacesensitive experiments to understand the oxidation state of iron, along with the change of oxidation states of oxygen through DFT supported Bader charge analysis. Experimental and DFT studies together have been used to investigate the phase transformation with respect to cycling and lithium concentration. In the end, we conclude the redox mechanism that is responsible for extra capacity at the higher potential in case of lithium iron silicate. Moreover, it is well known that lithium-iron silicate cathode is best known for fundamental study, but no attempts to make full-cell LIB are available in the literature. This is, to the best of our knowledge, the first attempt to make full-cell against graphite and show excellent and comparable performance. 3.1. Crystal structure and microstructural analysis of synthesized Li2FeSiO4. To start with, the crystal structure analysis of synthesized Li2FeSiO4 was performed by synchrotron X-ray powder diffraction (SXRD). Figure 1 shows the Rietveld refined fitted SXRD pattern of Li2FeSiO4. SXRD profile of synthesized Li2FeSiO4 before cycling has been fitted well with bi-phase fittings having space groups Pmn21 (93.5%) and P21/n (6.5%) respectively. The 10 ACS Paragon Plus Environment

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Rietveld fitting reveals that a small amount of monoclinic phase is present along with the orthorhombic phase in the synthesized sample. The goodness of fitting (χ2) value of 1.22 confirmed that the experimental data is in good agreement with the fitted profile.

Intensity (arb. unit)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Rp : 2.13 % Rwp: 2.51 % Rexp: 1.32 % 2



5

10

15

20

25

2 (degrees)

30

: 1.22

35

40

Figure 1. Rietveld refinement plot along with the fitted parameters SXRD pattern of synthesized Li2FeSiO4. To understand the particle morphology, scanning electron microscopy (SEM) technique was employed. The SEM images (Figure S2) of synthesized Li2FeSiO4 are given in supporting information (SI) is observed to be in the range of 32‒140 nm with an irregular spherical shape. This synthesized sample was used for further in situ and ex situ studies.

3.2. Ex situ Mössbauer spectroscopy and X-ray Absorption near Edge structure (XANES) Spectroscopy Analysis upon Cycling.



In recent literature of Li2FeSiO4, it is observed that the first charge capacity to be less than the first discharge capacity and the capacity obtained in subsequent cycles.24,35,36 In current work, we also obtained first charge capacity (value) is lower, which might be due to the oxidative nature of silicates after air exposure.35,37 However, the obtained discharge capacity indicates the extraction of 2nd lithium from the host matrix. In our previous work also, we have achieved more than 165 mA h g-1 discharge capacity in initial cycles indicating extraction of more than one lithium. The charge plateau above 4.3 V vs. Li/Li+ has been attributed to Fe3+/Fe4+ redox couple explaining that more than one lithium is electrochemically active. To confirm this hypothesis, we have done ex situ Mössbauer spectroscopy and XANES spectroscopy measurements on electrodes charged at various voltages (OCV, 3.8 V, 4.8 V vs. Li). Typical charge-discharge profile of Li2FeSiO4 at a current rate of 11 mA g-1 (C/15, 1C = 165 mA h g-1) in the potential window of 1.5‒4.8V vs. Li/Li+ is shown in Figure 2a. The charge profile exhibits two plateaus, first one around 3.0 V and second above 4.3 V vs. Li/Li+ which can be attributed to Fe2+/Fe3+ and Fe3+/Fe4+ redox couples, respectively. The Mössbauer spectra at different voltages during charging are also shown in Figure 2c and 2d. The isomer shift, quadrupole splitting, and area for fitted spectra have been summarized in Table 1. Li2FeSiO4 has been oxidized to LiFeSiO4 (36 %) before cycling which can be seen by fitted parameter at OCV. The Mössbauer spectrum at OCV consists of two doublets with isomer shift of + 0.99 and + 0.31 mm s‒1, which can be assigned to tetrahedral Fe2+ and Fe3+, respectively.38,39 The presence of Fe3+ at OCV indicates the partial oxidation of Li2FeSiO4, which is also well supported by the charging profile of Li2FeSiO4 as shown in Figure 2a, where we can clearly see that plateau around 3.0 V vs. Li/Li+ is shorter in first charging than the second charging curve.


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Table 1. Fitted parameters of Mössbauer spectra of the charged electrode at different state of charges. Sample

OCV 3.8 V 4.8 V

Iron Sites (Doublet) Fe2+ Fe3+ Fe2+ Fe3+ Fe3+ Fe3+ or Fe4+

Isomer shift (δ) mm/s

Quadrupole splitting (∆) mm/s

Line width (Г) mm/s

Relative Area (RA) %

0.99 0.31 0.87 0.23 0.45 0.16

2.52 0.70 2.45 0.66 0.81 0.66

0.35 0.77 0.30 0.55 0.34 0.44

63.95 36.05 29.45 70.55 25.67 74.33

During the charging process, the Mössbauer spectrum collected at 3.8 V has also been fitted by two spectra with isomer shift of +0.87 and + 0.23 mm s-1 indicating Fe2+ and Fe3+, respectively. The doublet area of Fe3+ increases from 36% to 70% upon charging up to 3.8 V with a charge capacity of ~ 50 mA h g-1. On charging the electrode further up to 4.8 V, a new doublet with decreased isomer shift of +0.16 mm s-1 is observed, which might be ascribed to the formation of Fe4+, indicating Fe3+/Fe4+ redox couple is active at higher voltage, i.e., > 4.3 V. The doublet area corresponding to a new doublet with decreased isomer shift of +0.16 mm s-1 in our study is 74% with a charge capacity of ~ 123 mA h g-1. This can be assigned to the high spin Fe3+ or low spin Fe4+. From the above discussion, we can conclude that Fe3+/Fe4+ redox couple might be partially electrochemically active in initial cycles.38–40 The difference in capacity of about 10-15 mA h g-1 for Li2FeSiO4 obtained during the charging process from charge-discharge profile (Fig. 2(a), 2(b), 2(c), and 2(d)) and from quantitative analysis of the doublet area change from Mössbauer spectra during charging, can be attributed to low signal to noise ratio in Mossbauer spectra. Moreover, to further confirm the change in oxidation state and co-ordination geometries of iron, Li2FeSiO4 samples at OCV and at different states of de-lithiation have also been characterized by using ex situ XANES



at room temperature (Figure 2e and 2f) by charging electrode to different voltages, i.e., 4.0 V, 4.5 V and 4.8 V.

Figure 2. (a) The charge-discharge profile of Li2FeSiO4 sample for the first two cycles at the current density of 11 mA g-1 in potential gap of 1.5 - 4.8V; The Mössbauer spectra of Li2FeSiO4 (b) at OCV (c) charge at 3.8 V (d) charge at 4.8 V, (e) Normalized Fe K-edge XANES spectra of all electrode samples with Fe metallic foil, FeO, and Fe2O3 standards, (f) Fe K-edge XANES spectra of all electrode samples showing pre-edge features along with Fe metallic foil, FeO, and Fe2O3 standards. 14 ACS Paragon Plus Environment

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XANES is an element-specific characterization technique because the X-ray absorption edges of different elements have different values of energies. The Fe metal foil has been used to calibrate the photon energies with the X-ray absorption coefficient µ recorded in fluorescence mode.25,26 Figure 2e shows the edge step normalized XANES spectra of Fe K-edge XANES for all the samples along with Fe standards FeO (Fe2+), and Fe2O3 (Fe3+). Features in the Figure 2e are shown as “P” which is due to dipole forbidden 1s-3d transition, “E” corresponds to 1s-4s monopole transition, while the feature marked “W” relates to the dipole 1s-4p transition.25,41 These features have been observed to shift towards higher energy with the increase in oxidation state. The values for change in the pre-edge and edge positions are shown in Table S1 of supporting information at different states of de-lithiation. Figure 2f displays the pre-edge portion of the Fe K-edge spectra and the data show smooth and regular progress of the pre-edge energy to upper energy as Li concentration decreases, which shows that the average oxidation state of Fe varies from Fe2+ to an upper state. The shifting of upper energy suggested that the nuclear bonding around the core ion is stronger so, from this observation, we can see that the fully delithiated sample (having Fe ions in Fe3+ or upper state) has tight nuclear bonding than the fully lithiated sample (containing Fe ions in +2 state). The main absorption edge of Fe K-edge XANES spectra for the electrode samples at OCV and 4.00 V lies in between the energy position of standard FeO and standard α-Fe2O3, confirming the charge state of Fe ions for these two samples lying between +2 to +3. Whereas, in case of delithiated samples charged up to 4.5 V and 4.8 V, the observed energy position lies at a higher value than that of the spectra of standard α-Fe2O3, i.e., indicating the oxidation state of Fe ions for these two electrode samples to be slightly higher than + 3. This suggests the presence of Fe4+ ions in the electrode samples charged to 4.50 V and 4.80 V. However, these energy positions of charged electrode samples (up to 4.5 V and 4.8 V) for the Fe ions has been measured somewhat



lesser than the Fe K-edge energy (7128 eV) reported in literature.25,42 This hints the existence of only a small amount of Fe4+ ions in samples, which is in good agreement with Mössbauer spectroscopy analysis, as well. Further, to obtain the average shifting of Fe-charge states a linear interpolation of charge states of the electrode sample along with standard has been discussed in SI (Figure S3). 3.3. Change in Crystal Structure with Cycling. Figure 2a shows that the voltage plateau of the 2nd charge is dissimilar and lesser than the 1st charge plateau while 1st and 2nd discharge plateaus are nearly identical. This occurrence is described as the structural reordering taking place in early charge profile.13,43 To reconfirm ex situ SXRD measurements have been performed while cycling. The ex situ SXRD reveals a new structure after de-lithiation, which has been shown in Figure 3. In consecutive cycling, the initial pattern has not been recovered. Two possibilities can be interpreted from present results. First being that Li2FeSiO4 is transforming to a new phase and the other being that it is changing to amorphous phase during lithiation or else crystallite size is decreasing substantially. To confirm this hypothesis, the Rietveld fitting of the obtained SXRD has been performed. Figure 3 shows that the fitted SXRD profile patterns by Rietveld refinement, which were collected at different states of de-lithiation by charging electrode to 4.0 V, 4.5 V and 4.8 V, with respect to Li/Li+.


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YObs

YCalc

(d)

YObs - YCalc

Bragg _Position Rp : 2.82%

4.8V

Rwp: 4.64% Rexp: 3.25% 

Intensity (arb. unit)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(c)

4.5V

2 : 1.73

Rp : 2.21% Rwp: 4.32 % Rexp: 3.37 % 

(b)

4.0V

2

: 1.98

Rp : 1.98 % Rwp: 2.01 % Rexp: 2.37 % 

(a)

OCV

2

: 0.83

Rp : 2.13 % Rwp: 2.51 % Rexp: 1.32 % 

5

10

15

20

25

30

2 (degrees)

2

: 1.22

35

40

Figure 3. SXRD pattern and Rietveld refinement fitted profile of (a) synthesized sample, (b) charge to 4.0 V, (c) charge to 4.5 V (d) charge to 4.8 V. The SXRD profile of the samples at 4.0 V during the de-lithiation process is well fitted with biphase fittings having space groups Pmn21 and P21/n. Further, after charging electrodes to 4.5 V and 4.8 V, the obtained SXRD profile is fitted with a single-phase having space group Pnma. The values of lattice constant, the goodness of fitting (χ2) and phase percentage for all the samples are organized in Table 2.



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Table 2. Lattice parameters, values of Rp, Rwp and Rexp from the Rietveld refinement of the SXRD pattern of the different delithiated samples. Sample

Phase

Name LFS@OCV

[email protected]

Space

a(Å)

b(Å)

c(Å)

Rp

Rwp

Rexp

χ2

Phase (% age)

Pmn21

6.2894

5.3837

5.0094

2.13

2.51

1.32

1.22

93.5%

P21/n

8.2631

5.0350

8.2526

Pmn21

6.2632

5.3787

4.9659

P21/n

8.2727

5.0299

8.2445

group Double

Double

6.5% 1.98

2.01

2.37

0.83

94.28% 5.72%

[email protected]

Single

Pnma

10.3527

6.5667

5.0189

2.21

4.32

3.37

1.98

100%

[email protected]

Single

Pnma

10.3704

6.5538

5.0287

2.82

4.64

3.25

1.73

100%

The fitted data values for all samples shown in Table 2, confirm that the experimental data agrees well with the fitted profile. As per the current observation, a clear idea about the phase transition achieved and a similar type of phase transition has been observed by other groups as well by using ex situ SXRD.13 In the present study, the irreversible structural change after the first cycle upon charging can be concluded from ex situ SXRD measurements. The irreversible structural change results in the potential plateau shift after the first cycle. Moreover, the crystallite size is obtained around 5.5 nm as calculated by Scherrer’s formula44 in synthesized material, which reduces to 1.52 nm after cycling. This change in crystallite size upon de-intercalation shows clear evidence of a substantial decrease in crystallite size along with irreversible structural change with cycling. The change in structure was further studied by FTIR. FTIR experiments have been performed in transmittance and ATR approach as shown in Figure S4a of supporting information. Furthermore, theoretical studies using ab-initio density functional theory (DFT) have also been undertaken to understand 18 ACS Paragon Plus Environment

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the mechanism for second lithium reversibility and investigate the change in crystal structure of Li2FeSiO4 upon lithium extraction, which has been discussed in the next section. 3.4. Change in Crystal Structure with Cycling Crystal Structure and Phase Analysis of Simulated Li2FeSiO4 Polymorphs. Experimentally observed polymorphs of Li2FeSiO4 having symmetry Pmn21, Pnma, and P21/n are again studied in this work using DFT calculations. It can be seen from Figure 4 that all three cations (Li+, Fe2+, Si4+) in each polymorph are tetrahedrally attached with oxygen. The only difference remains in their interconnectivity of different LiO4, FeO4 and SiO4 tetrahedra, which results in a different crystal symmetry.5–9,11

Figure 4. Optimized crystal structure of Li2FeSiO4 polymorphs; (a) Pmn21, (b) Pnma, (c) P21/n, (d) demonstration of orientation of LiO4, FeO4 and SiO4 tetrahedra in Pmn21, (e) representation of 19 ACS Paragon Plus Environment


the edge-sharing between LiO4 and FeO4 tetrahedra in Pnma, (f) the pattern of Fe-Si-Fe-Si sequence in SiO4 /FeO4 chains in Pnma, (g) representation of the edge-sharing between LiO4 and FeO4 tetrahedra in P21/n, and (h) the pattern of Fe-Si-Fe-Si sequence in SiO4 /FeO4 chains in P21/n. In case of Pmn21, corner-sharing identical tetrahedra are found to build "tetrahedral-slabs" in the (001) plane, stacked perpendicular along the c-axis. In each slab, two different one-dimensional chains of tetrahedra, one with pure LiO4 and another with an alternate stacking of FeO4 and SiO4 are running along [010] direction. These two different one-dimensional chains of tetrahedra repeat themselves along [100] direction. Here, in each upward-directed tetrahedron, cations occupy half of the tetrahedral sites in a slightly distorted hcp oxygen array. In all three polymorphs, each corner of FeO4 tetrahedra is connected to the two Li+ and one Si4+ cation. However, in Pnma and P21/n based Li2FeSiO4 the orientation of tetrahedra is not unidirectional like Pmn21. The succession of one-dimensional chains are found to run along [100] direction in Pnma, while the same array of chains is noticed along [-101] direction in P21/n, as compared to [010] direction in Pmn21. The alteration of Fe and Si (Fe-Si-Fe-Si) is found common in all three structures (Pmn21, Pnma, and P21n), however, the orientation of FeO4 and SiO4 tetrahedra is seen different in different polymorphs. It is identified as up-down-up-down in Pnma (Figure 4f), and up-up-down-down in P21n (Figure 4h), instead of up-up-up-up in Pmn21 (Figure 4d). In Pmn21 only corner-sharing occurs between tetrahedra which can be clearly seen from Figure 4d. While, in case of Pnma along with corner-sharing, a given FeO4 tetrahedron shares two edges with adjacent LiO4 tetrahedra, whereas in P21/n only one edge is shared between FeO4 and LiO4 tetrahedra, as evident from Figure 4e and 4g, respectively. The edge sharing is not seen between FeO4 and SiO4 tetrahedra because of the substantial difference in the Fe‒O and Si‒O bond lengths. Also, the repulsion between the Fe2+‒Si4+ ions is more as compared to Li1+‒Fe2+ ions. In all polymorphs, the SiO4 tetrahedra are 20 ACS Paragon Plus Environment

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found to be least distorted, while the distortion in FeO4 and LiO4 tetrahedra is found to be much more pronounced as compared to SiO4 tetrahedra (Table S2, S3, and S4 of supporting information). The value of distortion in FeO4 and LiO4 tetrahedra is highest in case of Pnma and least in Pmn21, whereas its value lies in between both structures in case of P21/n. This trend in distortion can be explained by edge-sharing in FeO4 and LiO4 tetrahedra as no edge-sharing are present in Pmn21, whereas one/two edge is shared in P21/n/Pnma, respectively. The optimized lattice parameters and volume per formula unit of all three polymorphs of Li2FeSiO4 are given in Table 3, with the experimental values in bold.9 It can be seen that for all structures the DFT calculated lattice parameters show decent agreement with the experimental results. This validates our choice of exchange-correlation functional and Hubbard correction. Table 3. The calculated lattice parameters and volume per formula unit for all the three polymorphs (Pmn21, Pnma, and P21/n) of Li2FeSiO4. System Pmn21 Pnma P21/n

Li2FeSiO4 Synthesized Li2FeSiO4 Experimental9 Li2FeSiO4 Experimental 9

a (Å)

b (Å)

c (Å)

α (°)

β (°)

γ (°)

6.3239 6.2894 6.3462 6.2855 8.2913 8.2312

5.3880 5.3837 10.7428 10.6594 5.0941 5.0216

5.0027 5.0094 5.1126 5.0368 8.3136 8.2316

90.00

90.00

90.00

Volume (Å3/f.u.) 85.2307

89.99

89.99

90.00

87.1422

90.00

99.18 99.20

90.00

86.6613

The simulated X-ray diffraction patterns are also presented for all relaxed crystal structures by using Powdcell software

45

and are compared with the standard ICSD pattern. In Figure S5 of

supporting information, we can see the simulated pattern of Li2FeSiO4 is matching well with ICSD for all three phases of Li2FeSiO4 (Pmn21, Pnma and P21/n). Considering the agreement between the calculated and measured lattice parameters and matching of simulated XRD patterns with the 21 ACS Paragon Plus Environment


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ICSD data, we, therefore, considered these relaxed structures for further study, i.e., to calculate the density of states and to create the different position of lithium in de-intercalated phases Li2FeSiO4. 3.5. Change in Crystal Structure, Bader charge, and Relative Stability of Li2FeSiO4 Polymorphs upon Lithium De-intercalation. To understand the de-lithiation process in Li2FeSiO4, we systematically removed Li atoms from different symmetry sites and performed the DFT calculations for all three space groups. The most stable, optimized crystal structures of the delithiated product in all three symmetries are given in Figure S6 of supporting information. The calculated lattice parameters and volume per formula unit of delithiated Li2-xFeSiO4 (x= 0, 1, 1.5, 2.0) for all the three polymorphs are listed below in Table 4. We can clearly see from Table 4 that the lattice parameters change with the extraction of lithium. The optimized structures of LiFeSiO4, Li0.5FeSiO4, and FeSiO4 were used to generate simulated XRD patterns for delithiated products. As we move from Li2FeSiO4 to LiFeSiO4 the highest intensity peak at ~ 24º shifts to the left with a relative increase in lattice parameters. Figure 5a clearly depicts the change in the XRD pattern with cycling. The structural change with delithiation results into change in overall space group symmetry along with the change in lattice parameters. The above findings strongly support our experimentally reported change in space group symmetry in silicate with charging. The change in XRD upon charging for Pnma and P21/n are given Figure S7 and S8 of supporting information, respectively. Table 4. The observed lattice parameters (DFT) and volume per formula unit for all the three polymorphs (Pmn21, Pnma and P21/n) of delithiated Li2FeSiO4. System Li2FeSiO4

a (Å)

b (Å)

c (Å)

α (°)

β (°)

γ (°)

6.3239

5.3880

5.0027

90.00

90.00

90.00


Volume (Å3/f.u.) 85.2307

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Pmn21

Pnma

P21/n

LiFeSiO4 Li0.5FeSiO4 FeSiO4 Li2FeSiO4 LiFeSiO4 Li0.5FeSiO4 FeSiO4 Li2FeSiO4 LiFeSiO4 Li0.5FeSiO4 FeSiO4

6.0863 6.1203 6.0796 6.3462 5.9561 6.0168 6.0221 8.2913 10.1045 9.3435 7.1841

5.6355 5.6563 5.7253 10.7428 11.3706 11.3776 11.8991 5.0941 5.0395 5.1075 5.3945

5.0456 5.1087 5.3806 5.1126 5.2804 5.3144 5.4028 8.3136 6.6532 7.0459 9.3213

90.00 85.14 89.97 89.99 89.99 89.99 90.00 90.00 90.00 90.00 89.99

89.22 88.78 89.99 89.99 89.99 89.99 90.00 99.18 89.04 94.03 93.29

90.00 102.2 90.06 90.00 90.00 74.11 90.00 90.00 90.00 90.00 89.99

86.5234 86.0516 93.6442 87.1422 89.4053 87.4782 96.7908 86.6613 84.6878 83.8543 90.1633

Figure 5. (a) Simulated X-ray diffraction pattern of Li2-xFeSiO4 (x = 0, 1.0, 1.5, 2) in Pmn21 based structures (b) The total energy vs. volume of Li2FeSiO4, LiFeSiO4, Li0.5FeSiO4, and FeSiO4 in all three space groups (Pmn21, Pnma, P21/n). The thermodynamic stability of Li2FeSiO4 polymorphs with cycling was evaluated by comparing the calculated total energy. Figure 5b shows the calculated total energies as a function of volume per formula unit (f.u.). Our calculation shows that the total energy for the optimized Li2FeSiO4 in Pmn21, Pnma and P21/n are -53.161, -53.132 and -53.163 eV per formula unit, respectively, which indicates P21/n to be most stable structure. The energy difference between Pmn21 and P21/n is 23 ACS Paragon Plus Environment


around 1.78 meV/f.u. It is also well known that smaller distortion (i.e., low spread between minimum and maximum bond length) leads to structural stability.9 Analyzing the structures on the basis of distortion, Pmn21 is found to be the most stable structure. The distortion values for Pmn21, Pnma, and P21/n are shown in Tables S2, S3, and S4 of supporting information. On considering both distortions as well as energy, Pnma is found to be the least stable polymorph in case of Li2FeSiO4. To analyze the stability of intermediate phases LiFeSiO4 and Li0.5FeSiO4 in all three space groups, their formation energy was calculated per formula unit. The formation energies are found to be –895.18, –909.94, and –874.85 meV/formula unit for Pmn21, Pnma, and P21/n based LiFeSiO4, respectively, whereas –416.84, –463.33, –410.20 meV/formula unit for Pmn21, Pnma, and P21/n based Li0.5FeSiO4, respectively, by using the following equation.46 E F  (Li x FeSiO 4 )  [ x 2 (Li 2 FeSiO 4 )  (1  x 2 )(FeSiO 4 )] .

The negative energy indicates that Li0.5FeSiO4 and LiFeSiO4 are stable at low temperature with respect to phase separation in Li2FeSiO4 and FeSiO4. Reversibility and the performance stability of material is also related to the stability of the intermediate phase. Based on the absolute value of formation energy in all the three space groups of LiFeSiO4, P21/n has lower stability, while Pnma has the highest stability, which indicates P21/n has better cyclic performance and stability as compared to Pmn21 and Pnma. Cyclic performance based on intermediate phase follows the trend, P21/n > Pmn21 > Pnma. Figure 5b clearly demonstrate that cycled Pnma is the most stable structure on comparing the total energies for the optimized delithiated Li2-xFeSiO4 in Pmn21, Pnma and P21/n. Therefore, it can be concluded that the symmetry of Li2FeSiO4 is changed to Pnma, upon cycling. This also matches with our experimentally obtained results wherein the space group of synthesized material is changed to Pnma, after the charging process (Figure 3 and Table 2). Here, it is also noteworthy to mention that anti-site mixing between Li and Fe during the charging of 24 ACS Paragon Plus Environment

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Li2FeSiO4 is considered to be a very common phenomena for such orthosilicate compounds, where the mixing proportion is found to be strongly dependent on the cycling rate.35,59,60 For example, in a recent study Masses et al.59 revealed that at a low current rate of lithium extraction from Li2FeSiO4 (i.e., C/50, where 1C =166 mA/g), Li/Fe anti-site mixing leads the delithiated phases towards their thermodynamically stable structure, as they observed a monoclinic to orthorhombic phase transition upon removal of one lithium from Li2FeSiO4. However, at a higher current rate (C/20), delithiated structures generally try to retain their parent symmetry and crystallize into a local minima structure with little cationic mixing. Such phase transformation is also demonstrated through first-principles calculations.46,61 For further validation from the perspective of our experiments, we also modeled LiFeSiO4 with various schemes for anti-site mixing and relaxed the structures by allowing cell shape as well as volume to vary, together with ionic relaxation. Basically, we considered two different proportion of Li/Fe anti-site mixing: (i) one Li and one Fe sites and (ii) two Li and two Fe sites are interchanged by applying possible permutations and combinations. The relaxed structures and their respective energies are depicted in Figure S12 of SI. On analyzing the figure, we interestingly find that at lower proportion of anti-site mixing most of the examined structures are energetically less favorable than the pristine Pnma-LiFeSiO4, while at higher mixing a monoclinic structure is found to be of comparable energy with the ground state energy of Pnma-LiFeSiO4 (~10 meV/atom). This reveals that an anti-site mixing can be expected into the delithiated structures of Li2FeSiO4, depending on the proportion of mixing or current rate. Table 5. Average Bader charge on atoms in Pmn21 based Li2FeSiO4. Values given in parenthesis are the normalized charge calculated by considering the charge on Si to be +4. Bader charge (Pmn21) System →

Li2FeSiO4

LiFeSiO4 25 ACS Paragon Plus Environment

Li0.5FeSiO4

FeSiO4


Atom ↓ Li Fe Si O

0.85 (1.10) 1.34 (1.74) 3.08 (4.00) -1.53 (-1.99)

0.86 (1.11) 1.76 (2.28) 3.09 (4.00) -1.43 (-1.85)

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0.87 (1.13) 1.79 (2.32) 3.09 (4.00) -1.33 (-1.72)

1.83 (2.36) 3.10 (4.00) -1.23 (-1.59)

To confirm the oxidation state of iron we have performed Bader charge analysis on materials upon lithium extraction. The effective charge on atoms can be calculated by taking the difference between Bader charge and total number of valence electrons for each atom. The change in average Bader charge with changing lithium concentration is tabulated in Table 5. The effective charges on Li, Fe, Si, and O in Li2FeSiO4 are calculated as 0.85 e, 1.34 e, 3.08 e, and -1.53 e, respectively. The value of effective charge is smaller than the classical oxidation states of Li+, Fe2+, Si4+, O2‒ in Li2FeSiO4, which is due to the covalency of Fe‒O‒Si linkage. Now, as we move from Li2FeSiO4 to LiFeSiO4, then to Li0.5FeSiO4 and finally to the FeSiO4 upon delithiation, the charge on lithium and silicon does not change much whereas, a change can be seen in case of iron and oxygen. To ascertain the change in the oxidation state of Fe and O with cycling, we recalculated the Bader charges of these two atoms by normalizing the charge on Si to be +4. This consideration clearly depicts that as we moved from Li2FeSiO4 to LiFeSiO4 the oxidation state of iron changes from +1.74 to +2.28, while -1.99 to -1.85 for oxygen atoms. Thus, in accordance with the experimental study, our DFT study also confirms that with first lithium cycling iron goes towards higher oxidation state i.e., +2 to +3 state. Simultaneously, oxygen is also found to take a part in the redox activity on moving towards lower oxidation state (i.e., -2 to -(2-δ) state). However, the change in the Bader charge is found to be more prominent for Fe (0.42 |e|) during the first Li extraction, while the second Li-ion extraction mostly triggered by the oxygen (0.20 |e|) redox. The same phenomena have further been confirmed from the analysis of change in density of states of the 26 ACS Paragon Plus Environment

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delithiated structures in the subsequent section. The detailed Bader charge on atoms for Pmn21, Pnma, and P21/n are given in supporting information (Tables S5, S6 and S7 of supporting information). 3.6. In situ Electrochemical Impedance Spectroscopy and Density of States from DFT with Cycling. The charge-discharge profile of Li2FeSiO4 consists of two plateaus, showing a phase transition in Li2-xFeSiO4 as lithium concentration is decreasing. To understand the phase transition between lithium rich and lithium poor phase during de-intercalation and intercalation, in situ electrochemical impedance spectroscopy has been performed during the charging and discharging process. Electrochemical impedance spectroscopy (EIS) provides time-dependent information about the properties of electrochemical processes such as electrode kinetics and phase transformation. Combined effects of impedances on Li2FeSiO4 electrode, lithium electrode, and electrolyte have been shown by EIS of Li2FeSiO4/Li half-cell during charging and discharging (Figure 6a and 6b). The observed change in the EIS during cycling can be attributed to the changes in Li2FeSiO4 electrode, as the contribution from lithium electrode and electrolyte are believed to be constant. Nyquist plots of Li2FeSiO4 at regular interval during cycling (current rate = 11 mA g-1) in potential window of 1.5-4.8 V vs. Li/Li+ are shown in Figure 6a and 6b. The EIS of Li cell consists of a depressed semicircle, wherein the high frequency and medium frequency semicircle overlap and a straight sloppy line lie at lower frequency. The obtained EIS at different potentials has been fitted by an equivalent circuit and showed in Figure 6a. The bulk resistance (Rb) of the cell imitates a mutual resistance of the electrolyte, separator, and electrode. RSEI and QSEI are lithium-ion migration resistance and capacitance of SEI layer formed on the surface of the electrode, which corresponds to the high-frequency semicircle. The middle-frequency semicircle 27 ACS Paragon Plus Environment


Rct and Qdl agree to charge transfer resistance and double-layer capacitance. The sloppy line at small frequency is linked to Warburg impedance, which corresponds to the diffusion of lithiumions from electrode-electrolyte interfaces. The value of bulk resistance (Rb) has increased as we move from OCV to 3.5 V during charging, which indicates a change in the electrode structure, which has also been supported by ex situ SXRD shown in Figure 3. It is also observed in Figure 6c that the value of Rct decreases during the charging process which means that charge-transfer resistance is decreasing with the decrease in lithium concentration in the system. This observation can be related to the transformation of insulator to conductor as lithium concentration is decreasing in the system and well established in LiCoO2 cathode material literature.47,48 To get more insight into the insulator-semiconductor and phase transformation phenomena of lithium iron silicate, the spin-polarized total density of states (DOS) together along with the projected DOS (PDOS) onto the d-orbitals of Fe atoms and the p-orbitals of O atoms are calculated for four different lithium concentrations of Li2-xFeSiO4 (x=0, 1, 1.5, 2). The PDOS and DOS for different lithium concentration (x=0, 1, 1.5, 2) of Pmn21 based Li2-xFeSiO4 are shown in Figure 6e-h. A schematic has also been drawn to visualize the change in band gap with the decreasing Li concentration in Figure 6d. Figure 6e demonstrates that in total DOS of Li2FeSiO4 (Pmn21) the main contribution near Fermi level comes from a strong hybridization between Fe-d and O-p orbitals. Here Fe atoms remain in 2 oxidation state and can be explained by the electronic configurations e↑↓↑ t2↑↑↑. A precise DOS peak has been noticed under the Fermi level in the spindown channel in all the three structures which are due to partially filled down spin Fe-3d states (Figure 7e, Figure S9a and Figure S10a of supporting information).


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Figure 6. Nyquist plots of Li2FeSiO4 at varying potentials; (a) during charging and (b) during discharging (c) variation of Rct with potential during charge and discharge of Li2FeSiO4, (d) The reported band gap of Li2FeSiO4, LiFeSiO4, and FeSiO4 calculated by GGA+U method for orthorhombic (Pmn21), Total and partial density of states of (e) Li2FeSiO4, (f) LiFeSiO4, (g) Li0.5FeSiO4, and (h) FeSiO4, calculated for orthorhombic (Pmn21) phase. Our calculations predict the band gap of pristine of Li2FeSiO4 having symmetry Pmn21, Pnma, and P21/n to be 2.83, 2.70, and 2.80 eV, respectively, indicating the insulating behavior of Li2FeSiO4, which is in excellent contract with the results obtained from preceding experimental and theoretical studies.49-51 In case of LiFeSiO4, i.e., on removing one lithium from Li2FeSiO4 introduces a vacancy in the lattice, which therefore shifts the Fermi level towards conduction band and reduces the band gap. The value of band gap decreases from 2.8 eV (Li2FeSiO4) to 2.1 eV (LiFeSiO4) in Pmn21 and 2.21 eV and 2.24 eV for the Pnma and P21/n based LiFeSiO4, respectively. Table 6. Calculated band gap for Li2-xFeSiO4 (x=0, 1, 1.5, 2.0) in case of Pmn21, Pnma and P21/n.

System

Band gap Eg (eV) Reported values in literature

Pmn21 (Li2-xFeSiO4) Pnma (Li2-xFeSiO4) x = 0, 1, 1.5, 2.0 x = 0, 1, 1.5, 2.0 x=0 x = 1 x = 1.5, x = 0 x = 1 x = 1.5, 2.0 2.0

P21/n (Li2-xFeSiO4) x = 0, 1, 1.5, 2.0 x = 0 x = 1 x = 1.5, 2.0

2.83

2.66

2.18

3.351

2.451

2.7549

2.6749

Half metal Half metal51 0.1049

2.72

2.21

Half metal

2.24

Half metal

The values of band gap for all three symmetries are given in Table 6. As we move from Li2FeSiO4 to Li2-xFeSiO4 (x=1, 1.5, 2.0), the increased ionic bonding characteristics of Fe-O influence the 30 ACS Paragon Plus Environment

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band gap as Si-O bonding interaction remains invariant with the change of lithium concentration (Tables S8, S9, and S10 of supporting information). A clear transition in the oxidation state of Fe is also noticed from the Bader charge analysis on moving from Li2FeSiO4 to LiFeSiO4. The Fe2+ state changes to Fe3+ which reduces the average Fe-O bond length and provides a strong wave function overlap. This also leaves a signature in total DOS of LiFeSiO4 through the exodus of the Fe-3d sharp-peak under the Fermi level due to fully empty minority-spin states in the system. Extraction of second lithium from LiFeSiO4, further changes the oxidations states of iron from Fe3+ to Fe(3+δ)+.52 In case of Li0.5FeSiO4 and FeSiO4, in addition to the unfilled minority states of Fe-3d, the majority Fe-3d states become partially filled, which results into de-stabilization of dorbitals and shifting of these partially filled up-spin states beyond the Fermi level, making the system half-metallic in nature. A similar movement is also noticed for O-2p state, reflecting strong hybridization between Fe-3d and O-2p states (Figure 6 e-h). This supports the increase in electronic conductivity observed with the de-lithiation process. The phase transition from insulator to conductor in Li2FeSiO4 obtained from in situ EIS can thus be correlated with the calculated bandgap of lithiated and delithiated phase. We have also calculated the density of states (DOS) for Li2-xFeSiO4 (x=0, 1, 1.5, 2.0) in case of Pnma and P21/n (Figure S9 and S10 of supporting information), the similar transition has been observed in another space group as well. As discussed above, the O-2p holes close to the Fermi level are also found to increase significantly in the upspin state as we move from Li0.5FeSiO4 to FeSiO4. Masese et al. also reported the formation of ligand hole in the O-2p band to be responsible for extra capacity achieved during the second oxidation step for Li2FeSiO4, which compensate the charge after 1st lithium extraction.18 The theory of ligand hole formation by Masese et. al.,18 is again checked by measuring the density of states of LixFeSiO4 through DFT.53 In the case of FeSiO4 O-2p, bands move overhead Fermi level



which confirms the oxygen involvement in the charge compensation route52 as moving from LiFeSiO4 to FeSiO4.53 From our lab, Sarkar et. al., have also demonstrated extraction of more than one Li+ for Li2Ru1-xIrxO3 system, showing a plateau above >4.0 V, which can be ascribed to the loss of oxygen from lattice.54 Moreover, the capacity contribution from oxygen evolution can be clearly calculated by dQ/dV plots.54 This is, however, not possible in case of silicate materials as the position of oxygen evolution and second redox couple are overlapping each other. Moreover, silicates are also known to have good thermal stability. It is also shown by A. Manthiram et.al. through differential scanning calorimetry (DSC) study that Li2-xFeSiO4 is much better in terms of thermal stability than layered oxides and comparable (slightly inferior) to LiFePO4,16 which make Li2FeSiO4 an attractive candidate for large-scale application.16 Moreover, through the DFT study, Amador and co-workers demonstrated that on the removal of second lithium from Li2-xFeSiO4, the charge compensation is through anionic oxidation.46 However, they have completely denied the possibility of oxygen evolution as oxygen release from lattice is energetically impossible because of strong lattice stabilization.46 From the above discussion, we can clearly interpret that the capacity contribution beyond 166 mA h g-1 might be coming from both, i.e., electrochemical active Fe3+/Fe4+ redox couple and O-2p hole formation as we cannot completely deny the possibility of anionic oxidation. 3.7. Average De-intercalation Voltage and Electrochemical Performance. Table 7 presents the observed typical voltages for Fe2+/Fe3+ and Fe3+/Fe4+ redox couple for all three space groups. In Li2FeSiO4, Li de-intercalation is a two steps process; first, at ~ 3.0 V and second at ~4.75 V. Voltage difference among different polymorphs can be seen, which might be because of Fe–O bond lengths (Tables S8, S9, and S10 of supporting information). A slight change in the bond length value is observed as we move from Pmn21 to Pnma. As explained by Manthiram 32 ACS Paragon Plus Environment

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et.al.55 and Padhi et.al.,56-57 the difference in the position of redox couple of isostructural compounds can be correlated by Fe-O bond covalence.55–57 The stronger Fe-O covalence results in the higher energy of redox couple, and hence, lower OCV. Increase in bond length results in more ionic Fe-O bond, leading to lower energy of anti-bonding state, hence increasing oxidation potential value for the reaction. The value of the first oxidation potential is highest in case of Pmn21 and lowest in case of Pnma indicating that the bond length contributes significantly to the voltage trend. The second oxidation potential is again highest for Pmn21 than for Pnma, but lowest for P21/n which is not linearly dependent upon bond length (Tables S8, S9, and S10 of supporting information), as reported by Sirisopanaporn et al. and factors like size, orientation, and distortion ensure effect balance potential9 restrained while the oxidation of iron in all polymorphs of Li2FeSiO4. The intercalation voltage value varies inversely with the distortion of FeO4 tetrahedra. Comparing the distortion values of Pnma and P21/n, the oxidation potential is found lowest for P21/n. This again indicates that the initial removal of second lithium is having more probability in the case of monoclinic phase.6 On observing the value of distortion after full de-lithiation the value is minimum for P21/n, which again supports the above conclusion that monoclinic, i.e., P21/n is electrochemically superior to orthorhombic phases of Li2FeSiO4, i.e., Pnma and Pmn21 in the initial cycle.6

Table 7. Calculated typical voltages for all Fe2+/Fe3+ and Fe3+/Fe4+ redox couple in Pmn21, Pnma and P21/n based Li2FeSiO4 Li2FeSiO4

Fe2+/Fe3+ 33 ACS Paragon Plus Environment

Fe3+/Fe4+


Pmn21 Pnma P21/n

3.02 2.95 3.01

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4.81 4.76 4.75

The differential capacity (dQ/dV) graphs for the 1st and 2nd cycle of Li2FeSiO4 at the current rate 11 mA g-1 in the potential range of 1.5-4.8V vs. Li/Li+ are showed in Figure 7a and 7b, respectively. The two redox peaks at ~ 3V and ~ 4.6 V can be noticeably observed from Figure 7a. The first anodic peak at 3 V can be endorsed to the Fe2+/Fe3+ redox couple. The second anodic peak nearly 4.5 V ascribed to the Fe3+/Fe4+ redox couple. Figure 7b depicts the dQ/dV plot for the second cycle, where we can see the anodic peak resembles Fe2+/Fe3+ redox couple shifted to 2.8 V, which indicates the structural transformation. Moreover, the intensity of peak corresponding to the first redox couple is also found to be less in the first cycle than the second cycle, which confirms the partial oxidation of material before cycling.


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Figure 7. Differential capacity (dQ/dV) plot for Li2FeSiO4 in a half-cell assembly at current rate of 11 mA g-1 in the potential window of 1.5-4.8V vs. Li/Li+ for (a) 1st cycle, (b) 2nd cycle, (c) Discharge capacity versus cycle number plot tested at current rate of 16.5 mA g-1, 33.2 mA g-1 and 83 mA g-1 in the potential range of 1.5-4.8V vs. Li/Li+, (d) The volume change with deintercalation in all three space groups (Pmn21, Pnma, P21/n), (e) Charge-discharge profiles and (f) Cycling performance of full cell composed of Li2FeSiO4 as cathode and graphite as anode in the voltage range of 1.0 - 4.5V at 1C current rate, (g) Cycling performance of full cell composed of Li2FeSiO4 as cathode and graphite as anode in the voltage range of 1.0 - 4.5V at 1C current rate, and (h) Ragone plot of Li2FeSiO4-Graphite full-cell shows energy versus power on gravimetric basis. The cyclic profile of pure Li2FeSiO4 composite cathode material showed in Figure 7c, at current rate of 16.5 mA g−1, 33.2 mA g−1, and 83 mA g-1 (C/10, 1C=166 mA h g−1) in potential range of 1.5-4.8V at 20 °C. The 1st discharge capacity at the current rate of 16.5 mA g−1 is 223 mAh g-1, which shows the reversibility of 1.35 Li+ ion. The capacity declines to 126 mA h g-1 after 50 cycles. While, the first discharge capacity at the current rate of 33.2 mA g−1 is observed as 210 mAh g-1, decreases to 136 mA h g-1 after 50 cycles. The cell drops ~ 41.6% of its early discharge capacity after 50 cycles at the rate of 16.5 mA g−1 in potential range of 1.5-4.8V at 20 °C, whereas this loss is around 35% when cell is run at the higher current rate of 33.2 mA g−1. This loss is further decreased to 28 % after 50 cycles while cell is cycled at more current rate of 83 mA g-1. The 1st discharge capacity (169 mA h g-1) reduced to 121 mA h g-1 after cycling. Electrochemically compelled phase change with structural instability related with the Fe4+ oxidation state of iron58 makes the utilization of the Fe3+/Fe4+ redox couples challenging in lithium iron silicate cathode material, shows into capacity fade in subsequent cycles. 36 ACS Paragon Plus Environment

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As known, volume change leads to stress and strain in the film during cycling. On calculating the volume change after 1st and 2nd de-lithiation for Pmn21 based Li2FeSiO4 is calculated to be 1.51% and 9.87%, respectively (Figure 7d). As evident from the Figure 7d, the volume changes after 1st de-lithiation for Pmn21, Pnma, and P21/n in Li2FeSiO4 is calculated as 1.51%, 2.59%, and -2.27%, respectively, and 9.87%, 11.07%, 4.04% after 2nd Lithium extraction. We can see from the plot that volume change during first lithium extraction is low as compared to second lithium extraction (around 10-11%) for orthorhombic based Li2FeSiO4. The cation-cation repulsion increases after removal of second lithium in each polymorph, which hints to the rise in the volume of the unit cell. The increased volume generates stress in the material leading to the capacity fading in the subsequent cycle. Therefore, the removal of second lithium is a limitation of Li2FeSiO4 as cycling beyond Fe2+/Fe3+ redox couple activity results in rapid capacity fading. The Full‐cell evaluation of Li2FeSiO4 cathode against Graphite anode (LFS-Graphite) has been studied after analyzing the individual half-cell reactions. The full-cell performance with different potential limitation on cathode side is studied by using Li2FeSiO4 as cathode. The Figure S11 of supporting information demonstrations initial charge-discharge of the LFS-Graphite fullcell carried out in the voltage range of 1.0‒4.5V at a current rate of C/40 (C=166 mA g-1). The initial two cycles were performed at very slow rate as these are considered as formation cycle to achieve the better cyclability. Figures 7e and 7f show the cycling profile as well as the cycling performance of the full cell, which is performed under same potential window with a current rate of 1C. The cycling performance is done at high rate since the capacity loss is less at high rate compared to slow scan in half-cell configuration of Li2FeSiO4. Figure 7f demonstrates that the full cell gives charge capacity ~124 mA h g-1 with a reversible discharge capacity of 95 mA h g-1 in third cycle. The difference in the charge to discharge capacity in the 3rd cycle is considered as the 37 ACS Paragon Plus Environment


polarization loss due to increase in the current density. The discharge capacity after 50 and 100 cycles are obtained as 80 mA h g-1 and 65 mA h g-1, respectively. This shows capacity retention of 84% after 50 cycles and 68% after 100 cycles when obtained capacity is normalized to the discharge capacity achieved in the third cycle. To improve the reversibility and capacity retention, a slight modification in the voltage window at cathode side is done in full-cell. The Figure 7g shows the cyclic performance of LFS-Graphite full cell carried out in the voltage window of 1.0 4.0 V. The first formation cycle is done at the rate of C/20 followed by the first discharge and subsequent cycles charge and discharge at 1C rate. In the first cycle, charge capacity is obtained as ~ 219 mA h g-1 at the rate of C/20 with a discharge capacity of ~ 87 mA h g-1 at 1C rate. The first discharge capacity is lower than the first charge capacity, which is due to sudden change in the current rate as well as an irreversible loss in the formation cycle, which is mainly believed to be due to the formation of SEI. The obtained discharge capacity in the second cycle is 63 mA h g-1. This cell shows capacity retention of 92% after 50 cycles and 73% after 100 cycles. Figure 7h shows Ragone plot created on the weight of cathode material as LFS-Graphite full cells are cathode limited. Predictably, the gravimetric energy density in LFS-Graphite full cell drops with the rise in power density, which means that an energy storage device which can fulfill the dual demand is very challenging to realize in practical. These results indicate that for practical applications, we need to optimize cells parameters as per the application requirements, i.e., whether we need high energy density or high-power density. 4. CONCLUSION Li2FeSiO4 was successfully synthesized via sol-gel reaction. Ex situ Mössbauer and XANES spectroscopy measurements of charged electrodes confirmed the charged state of Fe and its role for the structural stability of the material. SXRD and ex situ FTIR spectroscopy demonstrate the 38 ACS Paragon Plus Environment

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nature of phase transition between Li2FeSiO4/Li2-xFeSiO4 compounds (i.e., between lithium rich and lithium poor phases of lithium iron silicate). Moreover, an understanding to explain the reason behind the extra capacity in the initial cycles, capacity fading in the subsequent cycles and the phase transformation associated with the lithium concentration upon cycling has been developed thorough experimental and computational studies for Li2FeSiO4 charged electrode at the discussed potentials. By relating the density functional theory (DFT) calculations along with experimental observations, we have explained the reason for capacity fading mechanism. The improved electrochemical performance corresponding to Fe3+/Fe4+ couple can be attributed to the reduced charge transfer impedance and the decrease of Li-ion diffusion path length. Finally, the studied full-cell prototype by using Li2FeSiO4 (LFS) as cathode and graphite as the anode has been studied which showed the great potential for the practical application of this material. Overall, the abovediscussed results have established a one-to-one correlation between the experimental and theoretical observations along with the potential for the future applications of the full-cell prototype studies. AUTHOR INFORMATION Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS Authors acknowledge the financial and instrumental facilities provided by National Centre for Photovoltaic Research and Education (NCPRE-Phase II), IIT Bombay and Sophisticated Analytical Instrument Facility (SAIF) for the instrumental facility. MRP thanks IITB-Monash research Academy for providing Doctoral fellowship. We also thank RRCAT Indore for providing 39 ACS Paragon Plus Environment


SXRD and XANES experimental facilities. We thank Archna Sagdeo, Anuj Upadhyay, Abhaya and M. N. Singh for their assistance in SXRD and XANES measurements. The high-performance computing facility available at the School of Natural Sciences, Shiv Nadar University, was used to perform all calculations. REFERENCES (1) Gong, Z. L.; Li, Y. X.; He, G. N.; Li, J.; Yang, Y. Nanostructured Li2FeSiO4 electrode material synthesized through hydrothermal-assisted Sol-gel process. Electrochem. SolidState Lett. 2008, 11, 60–63. (2) Islam, M. S.; Dominko, R.; Masquelier, C.; Sirisopanaporn, C.; Armstrong, A. R.; Bruce, P. G. Silicate cathodes for lithium batteries: Alternatives to phosphates. J. Mater. Chem. 2011, 21, 9811. (3) Nytén, A.; Abouimrane, A.; Armand, M.; Gustafsson, T.; Thomas, J. O. Electrochemical Performance of Li2FeSiO4 as a new Li-battery cathode material. Electrochem. commun. 2005, 7, 156–160. (4) Sirisopanaporn, C.; Boulineau, A.; Hanzel, D.; Dominko, R.; Budic, B.; Armstrong, A. R.; Bruce, P. G.; Masquelier, C. Crystal structure of a new polymorph of Li2FeSiO4. Inorg. Chem. 2010, 49, 7446–7451. (5) Saracibar, A.; Van der Ven, A.; Arroyo-de Dompablo, M. E. Crystal structure, energetics, and electrochemistry of Li2FeSiO4 polymorphs from first principles calculations. Chem. Mater. 2012, 24, 495–503. (6) Zhang, P.; Hu, C. H.; Wu, S. Q.; Zhu, Z. Z.; Yang, Y. Structural properties and energetics 40 ACS Paragon Plus Environment

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