Probing the pseudo-1-D ion diffusion in lithium

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Cite this: Phys. Chem. Chem. Phys., 2016, 18, 22323

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Probing the pseudo-1-D ion diffusion in lithium titanium niobate anode for Li-ion battery† Suman Das,a Dipak Dutta,a Rafael B. Araujo,b Sudip Chakraborty,c Rajeev Ahujac and Aninda J. Bhattacharyya*a Comprehensive understanding of the charge transport mechanism in the intrinsic structure of an electrode material is essential in accounting for its electrochemical performance. We present here systematic experimental and theoretical investigations of Li+-ion diffusion in a novel layered material, viz. lithium titanium niobate. Lithium titanium niobate (exact composition Li0.55K0.45TiNbO51.06H2O) is obtained from sol–gel synthesized potassium titanium niobate (KTiNbO5) by an ion-exchange method. The Li+-ions are inserted and de-inserted preferentially into the galleries between the octahedral layers formed by edge and corner sharing TiO6 and NbO6 octahedral units and the effective chemical diffusion coefficient, is estimated to be 3.8 1011 cm2 s1 using the galvanostatic intermittent titration technique (GITT). Calculations based on density functional theory (DFT) strongly confirm the anisotropic Li+-ion diffusion in the interlayer galleries and that Li+-ions predominantly diffuse along the crystallographic b-direction. The preferential Li+-ion diffusion along the b-direction is assisted by line-defects, which are

Received 27th June 2016, Accepted 14th July 2016 DOI: 10.1039/c6cp04488c

observed to be higher in concentration along the b-direction compared to the a- and c-directions, as revealed by high resolution electron microscopy. The Li–Ti niobate can be cycled to low voltages (E0.2 V) and show stable and satisfactory battery performance over 100 cycles. Due to the possibility of cycling to low voltages, cyclic voltammetry and X-ray photoelectron spectroscopy convincingly reveal

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the reversibility of Ti3+ 2 Ti2+ along with Ti4+ 2 Ti3+ and Nb5+ 2 Nb4+.

Introduction Lithium-ion batteries have been widely recognized to be a very promising energy storage device for various small and large scale applications.1,2 The conventional lithium-ion battery, comprising of a transition metal lithium insertion cathode3 (e.g. LiCoO2, LiMn2O4, LiFePO4) and a graphite anode,4 has been demonstrated to possess very limited capabilities, especially at high operating current densities. One of the major reasons behind the poor performance of lithium-ion batteries at high operating current densities is the graphite anode. At high current densities, (nano)graphite, or as a matter of fact, for any lower dimensional carbon (e.g. nanotube, graphene) electrolyte degradation is even higher than for conventional graphite, leading to a

Solid State and Structural Chemistry Unit, Indian Institute of Science, Bangalore-560012, India. E-mail: [email protected] b Applied Materials Physics, Department of Materials and Engineering, Royal Institute of Technology (KTH), S-100 44 Stockholm, Sweden c Condensed Matter Theory Group, Department of Physics and Astronomy, Uppsala University, Box 516, S-75120 Uppsala, Sweden † Electronic supplementary information (ESI) available: Table S1 (comparison of 2y and d-spacing value) and Fig. S1–S10 (ICE-OES result, TEM images, E vs. Ot plot, cyclic voltammetry plot, dI/dV versus V plot, GITT data, battery cycling and Coulombic efficiency). See DOI: 10.1039/c6cp04488c

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an unstable electrode|electrolyte interface.5 As the lithium intercalation in graphite takes place at a potential close to the lithium plating voltage, Li-dendrites are expected to grow at a faster rate at higher currents and may pose serious safety hazards during device operation. In these circumstances, the scope of graphite is very limited at practical operating currents required for large sized lithium-ion batteries for potential deployment in electric vehicles and power grids. In the case of sodium-ion batteries, conventional graphite has to be chemically and structurally modified for reversible sodium insertion and deinsertion. Thus, there is a strong interest in materials exploration, especially in terms of layered materials, to seek alternatives to graphite. In recent times, various carbonaceous compositions6–10 and non-carbonaceous11,12 materials have been proposed as potential alternatives to conventional graphite anode. Numerous potential non-carbonaceous anode systems storing lithium reversibly via intercalation, conversion, and alloying reactions have been demonstrated.13 Generally, lithium storage using noncarbonaceous systems occurs in the proximity of 1 V, which is far above the lithiation potential in graphite (E0.1 V). However, the reduced overall cell voltage is typically offset by other beneficial features related to non-carbonaceous electroactive systems. Many of the potential non-carbonaceous anodes reversibly store far more lithium than graphite (0.5 Li per formula weight LixC6)

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and display superior rate capability. In this context, Ti-based compounds14 have shown considerable promise as potential alternatives to graphite or graphitic materials. Anatase and B-phase titania (TiO2) composites with carbon in varying geometrical dimensions (1–3 D) are very promising electrode materials.15,16 The Ti4+|Ti3+ redox couple with regard to Li+|Li appears at 1.8 V and thus should lead to a reduction in the full cell voltage when used with conventional transitional metal cathodes. However, TiO2 stores more Li (nearly 1 Li per formula weight of TiO2) and displays superior rate capability compared to graphite.17 These are also safer than graphite, as the lower cut-off voltage is E1 V. A major issue in TiO2 and many of the alternative electrode materials is the sluggish kinetics of lithium diffusion through the intrinsic 3-D crystal structure. Approaches to obtain fast Li-kinetics and thus high rate capability TiO2 has been achieved via the design of composites of TiO2 with electronic conduits such as carbon in complex morphologies using multi-step synthesis procedures.18–22 Thus, simpler and novel materials synthesis design is desired to produce a stable and high rate capable Ti-based anode for lithium-ion batteries. The current rate capability of an electrode is intrinsically correlated to the diffusion coefficient of the working ion in its intrinsic crystal structure. Spatial dimensionality is known to play a critical role in charge transport and scales inversely with the diffusion coefficient. Lithium ion diffusion pathways have been studied extensively in various cathode23–28 and anode materials.29–33 A rare and interesting observation is preferential diffusion in a 3-D structured material. Among the various anode materials, anisotropic lithium diffusion has been reported in layered Li1+xV1xO2,31 and LiNbS2.32 We probe here this fundamental physical property, viz. the influence of dimensionality on charge transport, in a novel anode material for prospective applications in lithium and sodium-ion batteries. We present here a novel and alternative Ti-based compound, viz. a layered lithium titanium niobate (abbreviated Li–Ti-niobate) as a prospective alternative anode material to graphite. The structure is unique as the interlayer spacing in the Li–Ti-niobate is flexible and has the ability to accommodate cations of various sizes without a significant strain on the crystal structure. Unlike all previous Ti-based systems which predominantly exhibit 3-D Li+ ion transport, the Li–Ti-niobate displays preferential pseudo-1-D Li+ ion transport. The introduction of mixed valency cations, viz. Ti4+ and Nb5+ also resulted in lowering the lithium intercalation potential from 1.8 to 1.7 V. The Li–Ti-niobate anode performance is further improved by downsizing the particles and subsequently coating them with carbon as an electrical conduit.

Results and discussion The schematic representation in Scheme 1a shows the discharge process of MTiNbO5 (M = K/Li). The lithium ions are stored anisotropically, specifically in the interlayer galleries of the 2D-anode material. The structure of the parent compound KTiNbO5 was first elucidated by A. D. Wadsley34 (cf. Scheme 1b)

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Scheme 1 (a) Schematic representation of the discharge process of MTiNbO5 (M = K/Li) and the preferential direction of Li+ ion diffusion pathways; (b) crystal structure of MTiNbO5 (M = K/Li).

and it exhibits a layered structure with an orthorhombic crystal system, space group: Pnma, a = 6.459 Å, b = 3.792 Å and c = 18.472 Å. The negatively charged octahedral [TiNbO5] layer is formed by TiO6 and NbO6 units. The stacking of the octahedral layers results in the given structure of the compound. The charge deficit of the layers is compensated by the positively charged K+ ions which occupy the interlayer space (Scheme 1b). The Ti4+ and Nb5+ ions in the KTiNbO5 have the same co-ordinates with a random arrangement. The unit cell of KTiNbO5 contains two octahedral layers and two sets of interlayer K+ ions. The K+ ions present in the galleries of this layered structure are weakly held by the negatively charged layers and this provides ample opportunity to exchange the K+ ions with Li+ ions, leading to the formation of Li–Ti-niobate. The powder X-ray diffraction pattern of the as-synthesized KTiNbO5 is shown in Fig. 1a(i). All reflections of the assynthesized KTiNbO5 can be indexed to the orthorhombic crystal system (space group Pnma, JCPDS reference code: 01-072-1047). No impurity peaks are detected in the synthesized KTiNbO5. The powder XRD (PXRD) pattern of Li–Ti-niobate (cf. Fig. 1a(ii)) is similar to that of KTiNbO5, with the exception of shifts of the basal (00l) reflections (l = 2, 4,. . .) (cf. Table S1, ESI†). The shifting of the (00l) peaks towards lower 2y values indicates the expansion of the lattice spacing due to the replacement of K+ ions by bigger sized hydrated Li+ ions. The hydration of the Li+-ions is confirmed by thermogravimetric analysis of the Li–Ti-niobate (Fig. 1b). For Li–Ti-niobate, the initial weight loss of around 2% at 160 1C is due to the water adsorbed on the surface of the compound. The subsequent weight loss of around 6% at 400 1C corresponds to the removal of water of hydration from the interlayer space of Li–Ti-niobate. The exact composition of the Li–Ti-niobate, as obtained from the ICP-OES (cf. Fig. S1, ESI†) and the TGA (Fig. 1b) results is Li0.55K0.45TiNbO51.06H2O. Raman spectroscopy is also performed here to investigate microstructural information of the as-synthesized niobate compounds.


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Fig. 1 (a) Powder XRD patterns of (i) KTiNbO5, and (ii) Li–Ti-niobate. The green vertical lines indicate the reference patterns of KTiNbO5 (JCPDS No. 01-072-1047). Inset: Shift of the (002) basal reflection when K+ ion is exchanged with Li+ ions. (b) Thermogravimetric analysis of Li–Ti-niobate (red line) and its derivative plot (blue line). (c) Raman spectra of (i) KTiNbO5 and (ii) Li–Ti-niobate.

In Fig. 1c(i), KTiNbO5 shows bands at 267 cm1 and 342 cm1 corresponding to the stretching and bending modes of vibrations of the TiO6 octahedra. The bands at 444 cm1 and 652 cm1 correspond to Ti–O–Ti stretching vibrations of edge shared TiO6 octahedra. The 544 cm1 and 891 cm1 peaks correspond to the Nb–O–Nb and NbQO stretching modes of vibration respectively.35,36 Interestingly, in the case of Li–Ti-niobate, slight changes in the Raman spectra (Fig. 1c(ii)), such as the splitting of peaks corresponding to NbQO, Ti–O–Ti and Nb–O–Nb vibrations, are observed. The changes are a direct consequence of the ion-exchange reaction, indicating the introduction of distortion in the TiO6 and NbO6 octahedra following ion-exchange. The oxidation states of the elements present in Li–Ti-niobate are confirmed through X-ray photoelectron spectroscopy (Fig. 2). The binding energy values are obtained after fitting the experimental data (red dots). In Fig. 2a and b, the peaks at 459.0 eV, 464.7 eV and 207.02 eV, 209.08 eV correspond to the 2p1/2, 2p3/2 peaks and 3d3/2, 3d5/2 peaks of oxygen coordinated Ti4+ and Nb5+ ions respectively. These binding energy values are similar to reported values for HTiNbO5.37,38 Fig. 2c shows the 1s peak of Li+ ions at 56.7 eV and the peaks at 292.5 and 295.3 eV in Fig. 2d are attributed to the presence of K+ ions in the sample. In Fig. 2e the peak at 528.8 eV is due to the O 1s. The satellite peak at 530.7 eV is due to oxygen which appears due to physically adsorbed water in the system.39 The morphology and shape of the as-synthesized Li–Ti-niobate particles are studied using transmission electron microscopy, TEM (Fig. 3). From the TEM images it is clear that the Li–Tiniobate particles are pseudo-rectangular in shape (Fig. 3a). The particle size analysis (inset: Fig. 3a) reveals the average aspect ratio of the pseudo-rectangles to be 1.3 (length = 34 nm and width = 26 nm). All spots in the observed selected area electron diffraction (SAED) pattern of Li–Ti-niobate (Fig. 3b) can be indexed to the Bragg reflections as observed in the XRD patterns (Fig. 1a). Representative high resolution transmission electron microscopy (HRTEM) images of Li–Ti-niobate are shown in Fig. 3c and d. The sharp spots in the first Fourier transform (FFT) patterns (insets: Fig. 3c and d) as obtained from these HRTEM patterns indicate the highly crystalline nature of the particles. These spots can also be well indexed


Fig. 2 X-ray photoelectron spectra of Li–Ti-niobate. The peaks correspond to the presence of (a) Ti-2p, (b) Nb-3d, (c) Li-1s, (d) K-2p, (e) O-1s elements.

(Fig. S2, ESI†) to the Bragg reflections shown in Fig. 1a. The atomic lattice fringes with a spacing of 0.9 nm (Fig. 3c) can be indexed to the (002) planes stacked along the crystallographic c-direction, and these are highly ordered. Similar observations from several Li–Ti-niobate particles suggest the arrangement of (002) titanium-niobate galleries in the Li–Ti-niobate are stacking-fault free. To account for the electrochemical performance of Li–Tiniobate, comprehensive understanding of the diffusion of Li+ in the intrinsic structure assumes utmost importance. In order to ascertain the chemical diffusion coefficient (DLi+) and the number of Li+ ions involved in the intercalation–deintercalation process in Li–Ti-niobate, galvanostatic intermittent titration technique (GITT) experiments are performed. GITT measurements are based on Fick’s law of diffusion and are related to the concentration gradients in the system. Since the rate of electrochemical reactions at the electrode surface and charge transport


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Fig. 3 (a) Bright field transmission electron microscopy (inset: size distribution), (b) selected area electron diffraction (SAED), (c and d) high resolution transmission electron microscopy (HRTEM) images (insets: respective fast Fourier transform patterns) of Li–Ti-niobate.

into the electrode material depend on the translational diffusion of the ions, the faster diffusion of Li+ ions in Li–Ti-niobate will imply improved electrochemical performance. DLi+ in Li–Tiniobate is calculated using eqn (1):40 4 mVM 2 DEs 2 (1) DLi ¼ ; t L2 DLi ; pt MS DEt where m is the mass of the active electrode material, VM and M are the molar volume and molar mass of Li–Ti-niobate, respectively, S is the electrode–electrolyte contact area, t is the constant current pulse time, L is the characteristic length of the electrode, and DEs and DEt are the change in the equilibrium potential and change in potential during the current pulse, respectively. The GITT charge–discharge behavior of Li–Ti-niobate is shown in Fig. 4a. Using eqn (1) and Fig. 4b, the chemical diffusion coefficient of the electrode material is found to be 3.8 1011 cm2 s1. In addition, the linearity of the equilibrium voltage (E) versus Ot has been shown in Fig. S3 (ESI†), which is important for the validation of eqn (1). It is well known that diffusion varies inversely with spatial dimension, i.e. D p (2n)1, where D is the diffusion coefficient and n the is dimensionality of space.41 So, if there exists a condition providing a configuration of favorable sites which can sustain long range ion transport, diffusion in 1-D will be the fastest and will lead to the highest diffusion coefficient. As Li–Ti-niobate is a 2-D material (layers are stacked along the crystallographic c-direction) the diffusion of Li+ along the c-direction is not expected to be favorable. Thus, it is essential to know whether there is any preferential direction for diffusion of the Li+ ions in Li–Ti-niobate. We have performed the climbed nudge elastic band (cNEB)42 method to predict the activation energy barrier with a sufficient


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Fig. 4 (a) Galvanostatic intermittent titration technique (GITT) charge– discharge curve of Li–Ti-niobate in the voltage range 0.2–2.75 V and (b) polarization curve for single GITT discharge process. (c) Energy minimized crystal structure of K0.5Li0.5TiNbO5, (d) Li+ diffusion along the a-direction of K0.5Li0.5TiNbO5 and (e) Li+ ion diffusion along the b-direction of K0.5Li0.5TiNbO5.

number of images for each Li+ ion hopping between the equilibrium positions. Fig. 4c shows the energy minimized structure of Li0.5K0.5TiNbO5. Along the Z direction, i.e. c-direction, there is a layer formed by Ti, Nb and O that induces an electrostatic potential that acts as a hindrance for ion diffusion in the Z direction. Thus, the possible diffusion directions are mainly X (or a) and Y (or b) for this system. It is also important to highlight that the energetic difference presented in the energy profiles of the Li diffusion, Fig. 4d and e, is a result of distinct environments felt by the Li+ ions in the initial and final states of the accounted hopping. This is due to some degree of disorder being introduced between the K+ and Li+ ions, and also the disorder relative to Ti and Nb cations in the formed layers. From the activation energy profile shown in Fig. 4d for diffusion along the a-direction, it can be seen that Li+ ions can initially move with an activation energy of 0.67 eV. In sequence, the Li+ ion gets trapped in a local minimum with a much smaller energy than the initial configuration. Then, Li+ ions have the chance to move forward in this direction with an activation energy of about 0.41 eV. From this analysis, a general activation energy for the migration of Li+ ions in the a-direction amounts to 1.67 eV, producing a diffusion coefficient of 3 1030 cm2 s1 at room temperature. On the other hand, once the analysis is done, taking into account the two steps of the ionic migration as shown in Fig. 4d, the diffusion coefficients are derived to be 1 1016 cm2 s1 and 3 1014 cm2 s1, respectively, for the barriers of 0.67 eV and 0.41 eV. The energy profile following the b-direction, as depicted in Fig. 4e, shows a first step with activation energy of 0.37 eV and 0.22 eV for the second step. In this case, the overall activation energy of 0.46 eV is derived from the cNEB calculations, leading to a diffusion coefficient of 5 1012 cm2 s1 at room temperature. As in the case for the a-direction, the derived Li diffusion coefficients are 3 1012 cm2 s1 and 2 1011 cm2 s1. In summary,


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the cNEB method predicts Li+-ion diffusion in the a–b plane, i.e. along layers. However, the dominant Li+-ion diffusion is along the b-direction because of its comparatively smaller energy barrier compared to that observed along the a-direction. It is has been proposed that ion diffusion through shortcircuit diffusion pathways, such as linear, planar and surface defects, can be faster compared to the lattice.43,44 The preferential Li+-ion diffusion along the crystallographic b-direction in the interlayer galleries as predicted from the theoretical studies of Li–Ti-niobate may also be attributed to the presence of linedefects, as observed in the lattice fringes of the HRTEM images of Li–Ti-niobate (Fig. 5). It is anticipated that the amount of dislocation defects (viz. line defects) along this direction is larger compared to that in the two other crystallographic directions. To ensure this, several HRTEM images are collected (Fig. 5a, b and Fig. S4(a–c), ESI†) and the corresponding FFT patterns are shown as insets in the respective figures. From these FFT patterns, the planes along the crystallographic (002), (200) and (020) directions are extracted (Fig. 5c–e and Fig. S4(d–f), ESI†) using the inverse FFT (IFFT) technique. It is found that the dislocation defects exist along the b-direction planes (Fig. 5e and Fig. S4f, ESI†) in all the extracted images, and are much larger in number compared to that along the two other crystallographic directions (Fig. 5c and d; Fig. S4d and e, ESI†). To probe the electrochemical properties of Li–Ti-niobate, cyclic voltammetry (CV) and galvanostatic charge–discharge cycling experiments are performed. In the cyclic voltammogram of Li–Ti-niobate (Fig. 6a) the cathodic peak that appears at B0.45 V (vs. Li+/Li) is assigned to the SEI (solid–electrolyte interface). The SEI is known to form as a result of electrolytic degradation on the electrode surfaces and appears only in the 1st discharge cycle. Three other cathodic peaks that appear

Fig. 5 (a and b) The HRTEM images of Li–Ti-niobate showing lattice fringes (inset: FFT patterns generated from the full image for (a) and the red marked region for (b), respectively). (c) The inverse FFT patterns generated by extracting the (002) plane from the inset of (a). (d and e) Inverse FFT pattern generated by extracting the (200) and (020) planes from the inset of (b). The insets in (c–e) represent patterns that are masked to obtain the respective inverse FFT patterns.


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Fig. 6 (a) Cyclic voltammogram of Li–Ti-niobate vs. Li|Li+. The inset contains the dI/dV versus V plot of the 1st cycle that shows three different redox couples. (b) Voltage profile of Li–Ti-niobate in the 0.2–2.75 V range. (c) Capacity vs. cycle no. plot of (i) Li–Ti-niobate in the 1–3 V range, (ii) Li–Ti-niobate in the 0.2–2.75 V range and (iii) cd-Li–Ti-niobate in the 0.2–2.75 V range. Inset: No. of lithium ions stored in two different voltage ranges. (d) Rate capability of cd-Li–Ti-niobate in the 0.2–2.75 V range.

at 2.4 V, 1.25 V and 0.65 V are assigned to the Nb5+/Nb4+, Ti4+/Ti3+ and Ti3+/Ti2+ redox-couples respectively (Fig. 6a (inset): dI/dV versus V and Fig. S5, ESI†). The appearance of the Ti4+/Ti3+ peak at a lower voltage of 1.25 V is not unusual and similar values have been reported for several Ti-based systems.45,46 The corresponding anodic peaks that appear at 2.67 V, 1.75 V and 0.85 V are assigned to the Nb4+/Nb5+, Ti3+/Ti4+and Ti2+/Ti3+ redox couples respectively. The peak that appears at around 2.1 V may be attributed to the degradation of the remnant interlamellar water molecules. The Ti3+/Ti2+ peak is reversible; however, it becomes weaker over successive cycles (Fig. S6, ESI†). In order to probe the existence of Ti3+/Ti2+, galvanostatic charge– discharge cycling of Li–Ti-niobate is performed in the voltage range of 0.2–2.75 V at a current rate of 0.1C (1C rate = 353 mA g1, considering a three electron process). The voltage profiles are shown in Fig. 6b. In this voltage range, Li–Ti-niobate shows (Fig. 6c(ii)) a 1st cycle discharge capacity of 348 mA h g1 which is nearly 98.6% of the theoretical value (=353 mA h g1 considering to a 3e process). The 1st cycle charge capacity is observed to be 132 mA h g1. The 2nd cycle discharge and charge capacities are found to be 165.6 mA h g1 and 124.7 mA h g1 respectively. The drastic drop of capacity between the 1st and 2nd cycles is attributed to the irreversible capacity loss due to SEI formation. However, the capacity attained stability following this, with retention of 93.8% between the 15th and 40th cycles. To ascertain whether the Ti3+/Ti2+ redox couple has any significant effect on the total capacity of Li–Ti-niobate, the battery performance is tested in the range 1–3 V (Fig. 6c(i)), as conventionally done for titanium based oxides47 and LiTiNbO5 by reported Colin et al.48 By doing so, it is observed that the 1st discharge capacity values decreased substantially to 313 mA h g1. In the subsequent cycles, the capacity decreased at a faster rate compared to the cycling in the voltage range 0.2–2.75 V. In the 15th and 40th cycles, the capacities are 86 mA h g1 and 78 mA h g1, respectively,


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which strongly suggests that the Ti3+/Ti2+ redox couple has a significant contribution to the total capacity of Li–Ti-niobate and that it is advantageous to cycle this material to a lower cut-off voltage, as low as 0.2 V, similar to the reported observations.49 The deviation in capacity from the theoretical value in both the voltage ranges of 0.2–2.75 V and 1–3 V can also be correlated with the number of Li+ ions involved in the electrochemical reaction of Li–Ti-niobate as obtained from the GITT. In the 2nd cycle, the number of Li+ ions intercalated in Li–Ti-niobate in the ranges 0.2–2.75 V and 1–3 V are 1.84 (instead of 3) and 1.15 (instead of 2) (inset: Fig. 6c and Fig. S7, ESI†), respectively. This corresponds well with the 2nd cycle capacities being 166 mA h g1 and 140 mA h g1 instead of the theoretical values of 353 mA h g1 (3e process) and 235 mA h g1 (2e process). The change in the chemical state of the electroactive ions in Li–Ti-niobate is probed by an ex situ XPS study (Fig. 7). The changes in the oxidation states of Ti4+ and Nb5+ are shown in Fig. 7a–c and d–f respectively. Fig. 7a and d show the XPS-spectra of the Li–Ti-niobate prior to discharge. The binding energy values are in good agreement with the literature.50 After a single discharge up to 1 V (Fig. 7b), the oxidation state of titanium changes from 4+ (binding energy values 459.0 and 464.7 eV) to 3+, with binding energy values of 457.2 eV and 462.2 eV for 2p1/2 and 2p3/2, respectively. However, when the sample is discharged up to 0.2 V, Ti2+ ions (Fig. 7c, binding energy values 455.5 eV and 460.2 eV for 2p1/2 and 2p3/2, respectively) are present along with Ti3+ ions. The oxidation

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state of the niobium in the pristine Li–Ti-niobate, changes from 5+ (cf. Fig. 7d, 207.02 and 209.8 eV for 3d3/2 and 3d5/2 peaks, respectively) to a mixed 5+ and 4+ (205.2 and 207.1 eV) at the 1 V discharged state (cf. Fig. 7e). However, when discharged to 0.2 V (Fig. 7f) all the niobium ions exhibit the 4+ oxidation state. It is well known that downsizing a particle decreases the effective diffusion path traversed by Li+ ions inside the material and consequently improves battery performance. The pseudorectangular as-synthesized Li–Ti-niobate particles (aspect ratio = 1.3, length = 34 nm and width = 26 nm) were downsized to 6 2 nm (Fig. S8a, ESI†). However, the performance of the cell with downsized Li–Ti-niobate was found to be poor (Fig. S9, ESI†). This can be accounted for on the basis of the introduction of a random distribution of defects in 3-dimensions and a large number of grain boundaries that form high particle– particle resistance pathways for Li+ ion transport. Furthermore, the absence of facile electron transport pathways may have a substantial effect in the capacity fading of miniaturized particles. To circumvent this, the downsized particles (d-Li–Ti-niobate) are coated with conductive layers of carbon (cd-Li–Ti-niobate) (Fig. S8b, ESI†). This resulted in improved performance (Fig. 6c(iii)) in the voltage range 0.2–2.75 V. Even though the cd-Li–Ti-niobate shows initial capacity fading (capacity loss = 4.65% per cycle), after the 10th cycle the battery shows excellent stability. From the 20th to the 50th cycle, the capacity loss per cycle is 0.46% and from the 50–100th cycle, the capacity loss per cycle is only 0.28%. A similar trend is also observed with the Coulombic efficiency (Fig. S10, ESI†). From the 10–100th cycle, the Coulombic efficiency of the cell is stable, attaining a value of 99%. The cd-Li–Ti-niobate battery also shows excellent rate capability (Fig. 6d) and can run at a wide range of currents (C/40 to 3C). Stable capacities of 182 mA h g1, 133 mA h g1 and 73 mA h g1 are obtained when the cell is run at current rates of C/40, C/10 and 1C respectively. A capacity of 165 mA h g1 is recovered at a rate of C/40 following the cell being charged at a fairly high current of 3C (capacity at 3C = 40 mA h g1).

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Fig. 7 Ex situ XPS study of Li–Ti-niobate in a lithium ion battery. (a) Ti before discharge, (b) after discharge up to 1 V and (c) discharge up to 0.2 V respectively, (d) Nb before discharge, (e) after discharge up to 1 V and (f) discharge up to 0.2 V respectively.


In summary, we have systematically discussed here via experimental and theoretical investigations the Li+ ion transport in Li–Ti-niobate, a novel non-carbonaceous anode for lithium-ion batteries. Unlike conventional anatase TiO2 or similar oxides which exhibit a Ti4+/Ti3+ redox couple, the Li–Ti-niobate exhibited two Ti-redox couples, viz. Ti4+/Ti3+ and Ti3+/Ti2+. This facilitated charging and discharging of the Ti-based compound to a voltage much lower than 1 V. The Li-ion diffusion in Li–Tiniobate is highly anisotropic and is essentially pseudo onedimensional, as predicted by the theoretical calculations. The theoretically estimated diffusion coefficients along the preferential direction agree very well with the effective diffusion coefficients estimated by experimental methods. Additionally, the ion diffusion is influenced by the presence of line defects, being highest in the direction with the highest amount of line-defects.


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The Li–Ti-niobate showed improved cyclability and rate capability on downsizing and carbon coating. The layered structure of Li–Ti-niobate is envisaged to be also highly beneficial for the insertion of other types of working ions, mainly alkali ions, viz. sodium and potassium. It also has the potential for the chemical design of electrodes which can reversibly store energy using a combination of storage mechanisms.


Acknowledgements Authors acknowledge the DST Nano Mission for financial support (SR/NM/NS-88/2010), the Centre for Excellence in Nano Science and Engineering (CENSE), Indian Institute of Science, Bangalore and the Solid State and Structural Chemistry Unit, Indian Institute of Science, Bangalore for infrastructural support. S. D. acknowledges the University Grants Commission (UGC) for providing Senior Research Fellowship. R. B. A., S. C. and R. A. would like to acknowledge Carl Tryggers Stiftelse for Vetenskaplig Forskning (CTS), Swedish Research Council (VR), Erasmus Mundus and StandUP for financial support. SNIC, HPC2N and UPPMAX are acknowledged for providing computing time.

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