Preferentially Oriented TiO2 Nanotubes as Anode Material for Li-Ion

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Preferentially Oriented TiO Nanotubes as Anode Material for LiIon Batteries: Insight into Li-Ion Storage and Lithiation Kinetics Andrea Auer, Engelbert Portenkirchner, Thomas Götsch, Carlos Valero Vidal, Simon Penner, and Julia Kunze-Liebhäuser ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b11388 • Publication Date (Web): 03 Oct 2017 Downloaded from http://pubs.acs.org on October 5, 2017

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ACS Applied Materials & Interfaces

Preferentially Oriented TiO2 Nanotubes as Anode Material for Li-Ion Batteries: Insight into Li-Ion Storage and Lithiation Kinetics Andrea Auer,† Engelbert Portenkirchner,† Thomas Götsch,† Carlos Valero-Vidal,‡ Simon Penner† and Julia Kunze-Liebhäuser†*. †

Institute of Physical Chemistry, Leopold-Franzens-University Innsbruck, Innrain 52c, Innsbruck, 6020, Austria.

‡

Advanced Light Source (ALS) and Joint Center for Energy Storage Research (JCESR), Lawrence Berkeley National Laboratory, 1 Cyclotron Rd., Berkeley, CA 94720, USA.

KEYWORDS: lithium ion battery, anode material, TiO2 nanotubes, preferential orientation, intercalation kinetics, electrochemical impedance spectroscopy (EIS).

ABSTRACT: Self-organized TiO2 nanotubes (NTs) with a preferential orientation along the [001] direction are anodically grown by controlling the water content in the fluoride containing electrolyte. The intrinsic kinetic and thermodynamic properties of the Li intercalation process in the preferentially oriented (PO) TiO2 NTs and in a randomly oriented (RO) TiO2 NT reference are determined by combining complementary electrochemical methods, including electrochemical impedance spectroscopy (EIS), cyclic voltammetry (CV) and galvanostatic cycling. PO TiO2 NTs demonstrate an enhanced ACS Paragon Plus Environment


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performance as anode material in Li-ion batteries due to faster interfacial Li insertion/extraction kinetics. It is shown that the thermodynamic properties, which describe the ability of the host material to intercalate Li-ions, have a negligible influence on the superior performance of PO NTs. This work presents a straightforward approach for gaining important insight into the influence of the crystallographic orientation on lithiation/delithiation characteristics of nanostructured TiO2 based anode materials for Li-ion batteries. The introduced methodology has high potential for the evaluation of battery materials in terms of their lithiation/delithiation thermodynamics and kinetics in general.

INTRODUCTION Lithium ion batteries (LIBs) are widely used in portable electric devices due to their high energy and power density. Further improvement of LIBs is crucial for large-scale and high power applications. Especially TiO2 has been studied intensively over several years as an alternative high capacity anode material as it has important advantages compared to state of the art, commercially available materials in terms of cost effectiveness, safety and environmental capability.1–4 In general, nanostructuring of the electrode material was found to enhance the kinetics of the lithiation reaction.2,5 In particular, selforganized and highly ordered TiO2 nanotubes (NTs) are promising due to their parallel oriented pore structure which allows for one-dimensional electronic and ionic conductivity. In addition, the thin walls provide short pathways for Li-ion diffusion and high tolerance to structural changes, i.e. in volume, occurring during charge/discharge cycling.1 Furthermore, there is no need for addition of conductive carbon and binder for the TiO2 NTs since they possess a sufficiently high electronic conductivity. Additionally, the Ti metal substrate itself serves as current collector.6–8 These advantages and the fact

ACS Paragon Plus Environment

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that upon annealing, i.e. thermal modification of the crystal structure under mild conditions, the anodically grown NTs used in this study do not show any changes in morphology,7,9 make them good model systems for fundamental research. The development of TiO2 crystals with specific high surface energy facets is particularly desirable since in LIBs the interaction between Li-ions and the electrode surface, followed by ion transport across the surface, is a key step. Properties of anatase (001) crystals as anode material in LIBs have been studied earlier,10–13 revealing an enhanced lithiation performance compared to randomly oriented anatase crystals, due to a more favorable Li insertion and a facile Li diffusion along the [001] direction. In contrast to nanoparticulate TiO2 crystals,14–16 anatase TiO2 NT arrays with a preferential orientation along the [001] direction have only been characterized in terms of their photovoltaic properties.17–19 To the best of our knowledge, only one study reports on their lithiation performance.20 However, this work compares highly oxygen-deficient, preferentially oriented NTs with fully oxidized randomly oriented TiO2 NTs.20 Since it is well known that oxygen vacancies enhance the lithiation/delithiation performance of TiO2 electrodes,7,21,22 the observed effects cannot solely be ascribed to the preferential orientation. In the present study, self-organized TiO2 NTs are anodically grown in a 2 wt% H2O containing electrolyte and annealed at 450 °C in air. The resulting NTs are mostly vacancy-free and possess exceptional crystallographic preferential orientation along the [001] direction (PO TiO2 NTs). To ensure an accurate comparison, reference TiO2 NTs with similar morphology are produced, using 10 wt% H2O in the electrolyte, and annealed under identical conditions, leading to randomly oriented NTs (RO TiO2 NTs). This work reports new insights into the influence of the crystallographic orientation of TiO2 NTs on the Li insertion and extraction processes as well as their kinetics. A combination of complementary electrochemical methods, such as cyclic voltammetry, galvanostatic cycling and electrochemical impedance spectroscopy, were used to probe lithiation and delithiation characteristics online.



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EXPERIMENTAL SECTION Preparation of TiO2 NTs. Anatase TiO2 NTs were prepared by a two-step anodization process20,23,24 and annealing in air adapted from Ref.17. Prior to anodization, titanium disks (99.6 %) were mechanically polished (0.5 µm MasterPrep, Buehler) and cleaned by ultrasonication in aceton, ethanol and deionized (DI) water. The first anodization step was carried out by exposing the polished surface (314.2 mm2) to an electrolyte containing ethylene glycol (99.5 % EMSURE, Merck) and 0.2 wt% NH4F (99.99 %, Merck, dried under vacuum and stored in an Ar filled glove-box with a water and oxygen content below 0.1 ppm) and applying an anodic potential of 20 V in an Ar purged two-electrode cell to ensure that no additional moisture can infiltrate. The water content of the electrolyte was either 2 wt%, leading to preferential orientation (PO), or 10 wt%, resulting in a randomly oriented (RO) reference NT array. After 1 h of oxidation, the initially-grown NT array was removed by a cathodic pulse of -3.0 V vs. carbon paper (Alfa Aesar, Toray carbon paper, TGPH-60) in a 1 M H2SO4 (Merck, EMPROVE, 95-98 %). The second anodization was carried out in the same electrolyte, which was kept under Ar, at a potential of 20 V for 75 min. After anodization, the sample was removed from the electrochemical cell and thoroughly rinsed with DI water. To convert the as-grown amorphous NTs to anatase TiO2 NTs, they were annealed in ambient air (open tube furnace) at 450 °C for 60 min with a ramp of 2.5 °C min-1 immediately after anodization. Materials characterization. X-ray powder diffraction (XRD) measurements were performed with a Siemens D5000 X-ray diffractometer, using Cu Kα emission. Diffractograms were acquired between 15 and 75 degrees (2 theta) with a step size of 0.02 degrees (2 theta) and an acquisition time of 2 s per step. Scanning electron micrographs (SEM) were taken with a Zeiss Crossbeam NVision 40 with an


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acceleration voltage of 15 kV, equipped with an InLens detector. Transmission electron microscopy (TEM) was performed on a Zeiss EM 10C, operated at 100 kV, as well as a FEI Tecnai F20 S-TWIN (high-resolution) analytical (scanning) transmission electron microscope (200 kV), equipped with an EDAX Apollo XLT2 silicon-drift detector for energy-dispersive X-ray spectrometry (EDX). For crosssectional images and TEM, the specimens were prepared by carefully scratching the TiO2 NTs from the Ti metal and transferring them onto a supporting grid under environmental conditions. Electrochemical measurements. Electrochemical measurements were carried out in a commercially available three electrode EL-Cell (ECC-Ref Cell, EL-Cell). Lithium foil (99.9 %, Alfa Aesar) was used as counter (CE) and reference electrode (RE). The self-organized NT arrays, which were anodically grown on Ti metal sheets and annealed in air, were directly mounted as working electrode (WE) in the electrochemical cell, without adding conductive carbon or any binder. WE and CE were separated by one layer of glass fiber separator (18 x 1.55 mm, EL-Cell) soaked with electrolyte. The electrolyte (99.9 %, Solvionics) used in all measurements was 1 M LiPF6 in a 1:1 (w/w) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC). All potentials herein are given with respect to the Li/Li+ reference electrode. All cells were assembled in an Ar filled glove-box (MBraun, UNIlab, H2O and O2 < 0.1 ppm). The assembled, air-tight cells were connected to a potentiostat (BioLogic VMP3) outside the glove box, and all measurements were performed at room temperature. Cyclic Voltammetry (CV) measurements were performed with scan rates of 0.05 mV s-1 and 5 µV s-1 between 3.0 and 1.1 V. Galvanostatic Cycling with Potential Limitation (GCPL) was carried out between 3.0 and 1.1 V. For PO TiO2 NTs current densities of around 13, 26, 2600, 5300 and 13000 µA g-1, corresponding to lithiation/ delithiation rates of C/20, C/10, 10C, 20C and 50C were applied. For RO TiO2 NTs current densities of around 8, 17, 1600, 3400 and 8400 µA g-1 corresponding to lithiation rates of C/20, C/10, 10C, 20C and 50C were applied. The mass of the active material was



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estimated by measuring the volume of the NT via SEM and by using the TiO2 anatase density of 3.84 g cm−3 to calculate the mass of the solid fraction. The current dependent gravimetric specific capacity for lithiation/ delithiation was determined from these measurements, as specified in Ref.6. Electrochemical impedance spectroscopy (EIS) was performed after three CV cycles to ensure quasireversible conditions. Spectra were collected in the potential range between 3.0 and 1.1 V for both lithiation and delithiation. A step size of 40 mV was applied in the vicinity of the Li insertion/extraction peak (1.9 -1.2 V), whereas from 3-1.9 V and from 1.2-1.1 V, step sizes of 70 mV and 50 mV were used. Each potential was kept constant for 5 h to ensure steady state conditions before the impedance measurement, ranging from 100 kHz to 5 mHz with a peak to peak amplitude of ±5 mV.

RESULTS AND DISCUSSION Morphology and Crystal Structure. Self-organized TiO2 NTs with a preferential orientation along the [001] direction are anodically grown by controlling the water content in a NH4F containing electrolyte. Figure 1 depicts scanning electron micrographs (SEM) of the NT arrays grown with 2 wt% (PO TiO2 NTs) and 10 wt% (RO TiO2 NTs) H2O. The NT arrays are highly ordered and covered by a thin capping layer, which is typically observed in organic electrolytes with small water contents.25–27


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Figure 1. SEM images (secondary electrons (SE)) of top and cross-sections of a),b) preferentially oriented TiO2 NTs (PO TiO2 NTs) and of c),d) randomly oriented TiO2 NTs (RO TiO2 NTs).

The cross-sectional images show the underlying tubular morphology, which confirms the formation of NTs underneath the capping layer. The NTs have an approximate average tube length of 1.50 µm and 1.05 µm, a wall thickness of 10 nm and 15 nm at the top and 18 and 27 nm at the bottom and a diameter of 57 and 82 nm for 2 wt% and 10 wt% H2O, respectively. XRD measurements (Figure 2) reveal the existence of the TiO2 anatase structure for both NT samples after annealing in air at 450 °C.28 In case of the PO TiO2 NTs (2 wt% H2O), a high peak intensity of the (004) plane can be observed, which indicates an unusually high percentage of (001) planes in the growth direction of the TiO2 NTs. This suggests a preferential growth of the NTs along the anatase [001]



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direction, which is confirmed by TEM measurements (see also Figure S1 in the supporting information, SI).

Figure 2. XRD patterns for a) PO TiO2 NTs and b) RO TiO2 NTs after annealing at 450°C in ambient air for 1 h.

In Figure 3, TEM images of the PO TiO2 NTs are shown. In the bright field overview image (Figure 3a)) the upper part of a NT array is visible. The wall thicknesses are determined to be around 11 nm, which confirms the dimensions observed in the SEM images (Figure 1). High-resolution TEM (HRTEM, Figure 3b)) confirms the collinearity of the NT axis and the c-axis of the anatase unit cell: In the image, the lattice fringes of the (101) planes (0.350 nm compared to the expected value of 0.351 nm,28 see also the inset fast Fourier transform (FFT) of the marked region) are pronounced. These fringes appear rotated by 22.6° relative to the growth direction of the NT (orange arrow in the image and orange line in the FFT). This agrees well with the theoretical value of 22.7° between the [001] direction and the [10-1] direction, corroborating that the NT axis is in the [001] direction. The other reflexes visible in the FFT are also ascribed to anatase: The lattice spacing of 0.240 nm with the corresponding spot in the top right corner of the FFT stems from the (10-3) planes (theoretical:


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0.243 nm), and the (004) planes are located at 0.240 nm (0.238 nm as per the crystal structure).28 The angles between these fringes in the image reveal that all of them originate from the same crystallite: There is an angle of 110.9° measured between the (101) and the (004) planes which agrees well with the calculated angle of 111.7°. Furthermore, the angle between (101) and (10-3) is 70.5° (compared to 70.7°), and (10-3) and (004) span an angle of 39.1°, which is close to the expected value of 39.9°.

Figure 3. TEM images of PO TiO2 NTs: a) bright field image of a NT bundle from the array and b) high-resolution TEM (HRTEM) image of a single nanotube. Inset in b) shows the fast Fourier transform (FFT) of the marked region.



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The overview image in Figure 3a) reveals that the walls of the NTs are not smooth, but rather corrugated. This is also shown in the HRTEM image: There is a noticeable kink in the outer wall, which is especially visible on the right side of the tube (marked by the light blue line). The lower surface is parallel to the (101) lattice planes, and the upper one perpendicular to the (004) planes. The angle between the two facets is 22.6°, which is the same as the angle measured between the [10-1] and the [001] directions; thus, the facets can be determined to be {100} and {101} truncations of the anatase crystals, indicating that the tubes consist of small crystallites that are arranged as shown in the scheme in Figure S2 in the SI. Preferential orientation along the [001] direction of self-organized TiO2 NTs is a known phenomenon,17 but no clear explanation for its occurrence is provided so far. It has previously been ascribed to the variation of the water content in the electrolyte and interpreted to result from the subsequent concentration variation of hydroxyl groups on the TiO2 surface during NT growth.17,20,29 However, studies of anatase TiO2 single crystals30 suggest that the preferential orientation originates from the fluoride content in the electrolyte, which is a known capping agent that lowers the surface energy of the (001) surface of anatase TiO2. During the growth of TiO2 NTs, fluoride species from the electrolyte accumulate at the phase boundary of the metal/metal-oxide interface.31,32 According to the flow mechanism,33 these fluoride species are pushed up the pore walls due to the plasticity of the oxide to form a fluoride rich layer, consisting of Ti-F or Ti-O-F species, all along the outer NT wall.31 The dissolution of these species, however, is strongly dependent on the H2O content.31 Hence, it is expected that, upon NT growth in an electrolyte containing only 2 wt% of H2O, little dissolution of fluoride species is taking place. During subsequent annealing, the usually volatile fluoride species may be trapped in the porous structure of the TiO2 NTs and act as a capping agent that stabilizes the {001} facets at the crystallization temperature. For the hydrothermal synthesis of anatase crystals with a high


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percentage of {001} facets, a dissolution-precipitation (or dissolution-recrystallization) mechanism has been proposed.34,35 This mechanism is, however, not applicable to the non-hydrothermal transition of amorphous to PO anatase TiO2 NTs. The proposition of a solid state rearrangement is therefore rather reasonable.36 This mechanism, postulated previously for the reaction of gas phase fluorine and amorphous TiO2 NTs, suggests that the dangling bonds of the amorphous surface are terminated by Fions, which energetically favor the {001} facet.35 Via annealing, fluoride free (see energy dispersive X-ray (EDX) spectrum, in Figure S3 in the SI), stable {001} facets are formed.35 With a water content of 10 wt% and more, the majority of the soluble Ti-(O)-F species is dissolved, and no preferential orientation along the [001] direction is obtained. This hypothesis is supported by an experiment where the fluoride species have been removed after anodization in the electrolyte containing 2 wt% H2O, by washing the sample with deionized H2O for 20 min. Annealing of the washed sample, leads to a random NT orientation (see Figure S4 in the SI). This confirms the crucial role of fluorides for the formation of PO TiO2 NTs.

Cyclic voltammetry and galvanostatic cycling. Cyclic voltammetry is used to determine the lithiation/delithiation characteristics of both NT arrays. Figure 4a) shows cyclic voltammograms (CVs) of PO TiO2 and RO TiO2 NTs measured with a scan rate of 0.05 mV s-1. Both electrodes exhibit peak couples at 1.74 V/ 1.99 V (PO TiO2 NTs) and at 1.73 V/ 2.02 V (RO TiO2 NTs) corresponding to the transition of α-Li poor LixTiO2 (0.01 < x ≤ 0.21)5 with anatase structure to the orthorhombic β-Lititanate (Li~0.55TiO2) phase; their positions are in good agreement with those reported in the literature.7,37–39 The lithiation/delithiation of anatase TiO2 is a two phase process, which has already been discussed elsewhere.40 In addition to the main peaks, a small peak pair is visible at ~1.46 V



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(lithiation) and ~1.74 V (delithiation) for both materials, which indicates a second phase transition from Li-titanate to fully lithiated LiTiO2 at these potentials.5,7,41 This second phase transition only happens in crystallites with a size < 10 nm due to kinetic limitations of Li+ diffusion in the fully lithiated phase.5,41 Upon delithiation, a peak broadening for both the PO TiO2 and RO TiO2 NTs can be observed. This broadening is however more pronounced in the case of RO TiO2 NTs. Furthermore, the PO TiO2 NTs show higher reversibility over repeated lithiation/delithiation, which is evident due to the smaller difference in the reduction and oxidation peak current densities.

Figure 4. CVs of PO (orange) and RO (blue) TiO2 NTs between 3.0 and 1.1 V with a) a scan rate of 0.05 mV s-1 and b) a scan rate of 5 µV s-1 in an electrolyte consisting of 1 M LiPF6 in a 1:1 (w/w) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC).

These results indicate enhanced kinetics for the PO TiO2 NTs. In addition, CVs for both NT arrays at a scan rate of 5 µV s-1 are shown in Figure 4b). In general, it is observed that, at lower lithiation/delithiation rates, the kinetic limitations (i.e. peak broadening and hysteresis) are reduced. The differences in the peak current densities between the PO and RO NTs, however, are still present.


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Figure 5. Galvanostatic charge-discharge curves of a) PO TiO2 NTs and b) RO TiO2 NTs recorded at different C-rates (C/20, C/10 and 10C) between 3.0 and 1.1 V.

Galvanostatic cycling is carried out between 3.0 and 1.1 V (Figure 5). The charge-discharge curves measured at different C-rates show a distinctive plateau at around 1.78 V, corresponding to the phase transition from Li-poor TiO2 to Li-titanate. In analogy to the small peak at around 1.5 V in the CVs (Figure 4), a small plateau is observed at the same potential in the galvanostatic curves corresponding to the second phase transition to Li1TiO2. A more detailed description of the specific regions of the chargedischarge curves can be found in Ref.7. Figure 6 shows the gravimetric capacities and corresponding coulombic efficiencies of PO and RO TiO2 NTs as a function of cycle number and C-rates. Both systems show stable cyclability in the investigated potential range between 3.0 and 1.1 V. RO TiO2 NTs exhibit a slightly higher reversible capacity of 185 mAh g-1 (Li0.55TiO2) at slow C-rates (C/20) compared to PO TiO2 NTs with 172 mAh g-1 (Li0.51TiO2). A possible explanation for this difference can be the uncertainty in the calculated active mass of the TiO2 NTs. Both the C-rate as well as the gravimetric capacity are effected



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by the determined mass. However, since this error stays constant, the changes in the lithiation performance are significant. At a faster C-rate of 10C, 94 mAh g-1 (Li0.28TiO2) are recorded for PO TiO2 NTs, and 79 mAh g-1 (Li0.23TiO2) for RO TiO2 NTs (cycle 5). For the PO NTs at 10C, directly after switching the C-rate, coulombic efficiencies of 95.5 % (cycle 5) are measured, the efficiency at the same C-rate increases to 100.0 % at cycle 100. For the RO NTs at 10C, directly after switching the C-rate, coulombic efficiencies of 89.4 % (cycle 5) are measured, the efficiency at the same C-rate increases to 100.0 % at cycle 100. At even faster C-rates of 50C for both RO and PO TiO2 NTs a capacity of 38 mAh g-1 is obtained. At these fast rates it can be assumed, that the bulk properties have a minor influence on the lithiation/delithiation and only the surface characteristics are important, therefore the same capacity values are obtained. After 300 cycles, both TiO2 NT samples exhibit reasonable cycling performance at slower rates (C/10), shown in Figure 6b). However, the cycling stability as well as the coulombic efficiency (99.8 % at cycle 320) of the PO TiO2 NTs are much higher, whereas the efficiency of the RO TiO2 NTs only reaches 98.8 % (cycle 320). XRD, measured after delithiation (Figure S5 in the SI), confirms that the preferential orientation is still present, even after 320 cycles. The crystallographic properties of the PO and RO TiO2 NTs stay the same, however two additional peaks can be found, most likely due to the electrolyte residue crystallizing on the sample surface. These findings indicate not only superior kinetics of the PO TiO2 NTs upon a change to faster lithiation/delithiation rates, but also imply that these kinetic effects result from differences in the bulk properties, since at 50C no differences in the measured capacities are obtained.


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Figure 6. Gravimetric capacities (squares) and corresponding coulombic efficiencies (circles) of PO TiO2 NTs (orange) and RO TiO2 NTs (blue), a) measured as a function of cycle number at different Crates and b) measured at C/10 after the fast cycles at 50C. Closed squares correspond to discharge (lithiation) and open squares to charge (delithiation) cycles.

From theoretical42 and fundamental experimental10 work, the origin of the enhanced performance of the PO TiO2 NTs over RO TiO2 NTs can be rationalized by a lower energy barrier in the surface transmission of Li-ions across the (001) surface. Furthermore, the more open structure of the lattice should lead to a more facile diffusion along the nanotube c-axis.

Electrochemical impedance spectroscopy. To further understand the origin of the superior kinetics of PO TiO2 NTs and to distinguish between kinetic and thermodynamic properties of the electrode material upon lithiation and delithiation, electrochemical impedance spectroscopy (EIS) is employed. This allows for differentiation of processes taking place at different time scales during Li insertion and extraction.



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Figure 7. Nyquist plots of a) PO TiO2 NTs and b) RO TiO2 NTs. c) CV of PO TiO2 NTs with colored points indicating the potentials of EIS data acquisition shown in a) and b).

Figures 7a)-b) show characteristic Nyquist plots of PO and RO TiO2 NT arrays in the potential range between 1.9 and 1.1 V (for potentials between 3.00 and 2.19 V see Figure S6 in the SI), indicated as color coded points in the CV shown in Figure 7c). Electrical circuit elements are determined after fitting the EIS data with the electrical equivalent circuit (EEC) depicted in Figure 8.43–45 Through careful analysis of the EIS data (see also Bode plots in Figure S7 in the SI), three well defined patterns are found, which are associated with specific processes taking place during lithiation/delithiation. At the highest frequencies from approximately 100 kHz to 25 Hz, a depressed semicircle can be observed, which corresponds to the parallel combination of surface film resistance Rf and surface film capacitance Cf. This first time constant relates to the Li+ migration through thin surface films, formed on the electrode.44,46 The thin surface films originate from the decomposition (and adsorption) of compounds from the carbonate-based LiPF6-containing electrolyte,47 and their resistance remains almost unchanged over the whole potential range (see Figure S8 in the SI). They mainly consist of inorganic precipitates such as LiF and Li2CO3.47,48


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The second medium-frequency arc relates to the interfacial charge transfer, represented by the charge transfer resistance, Rct, and the double-layer capacitance, Cdl, in parallel. The behavior of Rct and Cdl is highly potential dependent which has already been reported previously for a different electrode material.44 It is noteworthy that the charge transfer corresponds to a transfer of Li+ ions from the TiO2 NT to the surface film and vice versa, and does not take place at the electrode/electrolyte interface, due to the presence of the surface films on the active material. At lower frequencies, a small Warburg region, which is attributed to solid state diffusion in the active material, is observed. At even lower frequencies, capacitive behavior is dominant, which is displayed by a steep line with an angle higher than 45° in the Nyquist diagrams pointing out the initial part of a capacitive element. This capacitance Cint arises from the accumulation of Li+ ions in the bulk of the active material. Since it is observed in the lowest frequency range from about 10 to 5 mHz, it is equivalent to the derivative of the lithiation/delithiation curve at quasi-equilibrium,49 and thus delivers thermodynamic information of the intercalation reaction in TiO2 NTs.

Figure 8. Equivalent circuit used for fitting the EIS data. Rs: solution resistance, Rf: surface film resistance, Cf: corresponding capacitance (modelled with a constant phase element (CPE)), Cdl: double layer capacitance of the surface film/electrode interface (modelled with a CPE), Rct: charge transfer resistance, Zw: Warburg element, Cint: low-frequency internal chemical capacitance.



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The main parameters extracted from fitting the EIS data to the EEC displayed in Figure 8 are shown in Figure 9. The Rct for both PO and RO TiO2 NTs exhibits the highest value at 3.0 V and a fast decrease upon further lithiation. The PO TiO2 NTs show a higher degree of reversibility upon Li insertion and extraction: while at 3.0 V the RO TiO2 NTs start at a lower Rct (3 Ω g versus 7 Ω g for PO TiO2 NTs), their final Rct at 3.0 V delithiation is higher (7 Ω g versus 3 Ω g for PO TiO2 NTs). At potentials around the phase transition, the interfacial transfer of Li-ions for both PO and RO TiO2 NTs takes place at low resistance. In case of the RO TiO2 NTs, the resistance increases again upon further lithiation, exhibiting a maximum of 3 Ω g at 1.1 V. As shown in the CV and in the galvanostatic lithiation/delithiation curves, a distinctive volume fraction of the NTs (particle size < 10 nm) is converted to fully lithiated Li1TiO2.5 In Li1TiO2, diffusion of Li-ions is so slow that, once formed at the surface, additional Li cannot enter to further contribute to the lithiation reaction.50 This will increase the Rct at lower potentials, as the surface is passivated. Interestingly, for PO TiO2 NTs, Rct stays at its minimum until a lithiation potential of 1.1 V is reached, which indicates a better ability of the Li-ions to cross through the surface film/electrode interface. The crystallite sizes of the prepared NT samples, determined from XRD measurements (see Figure S10 in the SI), are identical for PO and RO TiO2 NTs. Hence, the observed differences cannot be ascribed to the impact of crystallite sizes.5


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Figure 9. Potential dependent electrical parameters of PO (orange) and RO (blue) TiO2 NTs derived from EIS fits: a) charge transfer resistance Rct, b) double layer capacitance Cdl and c) internal capacitance Cint for lithiation (closed circles) and delithiation (open circles).

The corresponding Cdl exhibits a pseudo-capacitive behavior with a broad peak at around 1.8 V due to the phase transition during lithiation. Figure 9c) depicts the internal capacitance, Cint, which accounts for the ability of the active electrode material to intercalate and accumulate Li-ions.44,49,51 The peakshaped potential dependency of Cint closely resembles the CV shapes at very low scan rates in Figure 4b). Interestingly, the peak height and hysteresis of Cint are completely identical for both, PO and RO TiO2 NTs, while this is not the case in the slow scan CV. The reason for this obvious difference is that EIS spectra are measured after long potentiostatic steady state (5 h); therefore, a quasi-equilibrium is established prior to EIS measurements, which is not the case for the CVs. The plot of Cint versus the electrode potential still shows a small thermodynamic hysteresis between lithiation and delithiation peak. The origin of this hysteresis can be interpreted as an effect arising from the use of a many-particle electrode.52,53 Since the TiO2 NTs are polycrystalline and the lithiation process proceeds sequentially, particle by particle, several equilibria persist at the same time.52,53 For example, in anatase TiO2 particles, either the anatase or the Li-titanate phase exists, and both phases are never coexisting in equilibrium.40 Therefore, even at zero-current, a small voltage gap of around 20 mV has been found between charge and discharge.52 In the present experiments, a hysteresis of 40 mV is measured, but the accuracy is limited due to the relatively large potential step size. The fact that no difference in the chemical capacitances, Cint, of PO and RO TiO2 NTs is found strongly suggests that the enhancement in performance of PO TiO2 NTs results from the faster interfacial Li insertion and extraction kinetics (Rct).



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These results are supported by theoretical work, reporting the charge transfer to be the overall ratelimiting step, based on the highest calculated energy barrier.12 However, in the nanotube system, the charge transfer may not be the only rate-determining step; therefore, it is also necessary to consider solid state diffusion in the bulk material. The diffusion coefficient D is estimated from EIS measurements at different, selected potentials (Figure 10). To determine the Warburg diffusion, the real impedance part is plotted versus ω-1/2 and the slope of the linear part is extracted (see Figure S9 in the SI). From the slope, the Warburg coefficient Aw is determined, which is related to the diffusion time constant τdif and the diffusion coefficient D by the equations51

࣎ࢊ࢏ࢌ = ૛ሾ࡭࢝ ࡯࢏࢔࢚ ሿ૛ , ࡰ=

࢒૛ , ࣎ࢊ࢏ࢌ

(1) (2)

where Cint is the differential capacitance (=internal capacitance) and l is the crystallite size determined from XRD measurements (see Figure S10 in the SI), which is 24.1 nm and 24.0 nm for PO and RO TiO2 NTs, respectively.

The potential dependence of the determined Warburg coefficients Aw is shown in Figure 10a). For PO TiO2 NTs, a higher Warburg coefficient is obtained at high potentials, before it decreases and shows a minimum which stays completely constant with a value of ~60 Ω s-1/2 from 1.78 V to 1.10 V. The RO TiO2 NTs, on the contrary, show a lower Aw value at 3.00 V, that reaches a minimum at the same potential (1.78 V), but slightly rises again to ~170 Ω s-1/2 at 1.10 V. The corresponding plot for the


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chemical diffusion coefficient D is shown in Figure 10b). D is inversely proportional to Aw (equations 12), therefore maxima of D could be expected at potentials, where Aw exhibits minima. This is, however, not observed since at these potentials PO and RO TiO2 NTs display high internal capacitances, which results in distinct minima of D near the phase transition. This minimum of D observed for TiO2 NTs is in good agreement with values reported in the literature for similar anode materials,51,54,55 and can be explained with the lattice gas model where strong attractive interactions between the Li-ions exist.56

Figure 10. Plots of a) the Warburg coefficient Aw and b) the diffusion coefficient D versus E for PO (orange) and RO (blue) TiO2 NTs, determined by extracting the Warburg coefficient from the EIS data and using equations 1-2.

The peak for D is much broader than the corresponding peak for Cint. Recently, it was found that during the Li induced phase transiton of anatase TiO2 particles, non-equilibrium conditions exist.40 After the solid solution stage of the lithiation, the β-Li-titantate phase starts to form. The migration of the boundary between the α and β phase is faster than the diffusion of Li-ions, which consequently leads to a Li concentration below the equilibrium value in the β phase. After this non-equilibrium stage, the



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diffusion-controlled equilibrium phase transformation proceeds until the whole particle is converted.40 This diffusion control at the phase transition potential (under quasi-equilibrium conditions) could explain the appearance of a broad peak in the diffusion coefficient, rather than a distinct, narrow peak as was found for Li1-xCoO2.44,54 A similar trend of D is observed with minimum values (1.78 V) of 2.9 x 10-15 cm2 s-1 and 5.5 x 10-16 cm2 s-1 for PO and RO TiO2 NTs, respectively. The similarity of the Li-ion diffusion in the investigated systems, strongly reinforces the previously made assumption that the slow charge transfer is the actual rate-determining step for Li insertion (extraction) into (from) anatase TiO2 NTs. However, the small differences of D indicate that solid state Li-ion diffusion might also contribute to the faster lithiation/delithiation kinetics. It is important to note that the determination of the Warburg coefficient Aw from EIS data with a wide frequency range, is rather challenging and lacks accuracy. To accurately determine the Li-ion diffusion in PO and RO TiO2 NTs, the application of additional electrochemical methods such as the Potentiostatic Intermittent Titration Technique (PITT) can be beneficial.44,55,57 The precise determination of the solid state diffusion coefficients in these nanotubular systems will be a focus of further studies in our group.

CONCLUSIONS The synthesis of TiO2 NTs with a specific and controllable crystallographic orientation is highly desirable for gaining fundamental understanding of the performance of Li-ion batteries based on selforganized TiO2 NTs. In this study, PO TiO2 NTs with preferential orientation along the [001] direction in anatase TiO2 demonstrate enhanced kinetics during lithiation/delithiation processes compared to RO


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TiO2 NTs. It is shown that a preferential orientation along the [001] direction facilitates the Li-ion insertion and extraction through a reversible decrease of the charge transfer resistance Rct. The peak height and hysteresis of the internal capacitance Cint, which is a thermodynamically governed measure of the amount of Li-ions intercalated in the TiO2 bulk, are approximately identical for both, PO and RO TiO2 NTs. This confirms that the enhancement in performance is a kinetic, rather than a thermodynamic, effect. In summary, by separating the kinetic and the thermodynamic properties of selforganized, anodically grown TiO2 NTs during lithiation and delithiation, it could be demonstrated that the preferential orientation along the [001] direction enhances the reversibility and the rate capability of this anode material.

ASSOCIATED CONTENT Supporting Information. TEM images with corresponding schematic representation of PO and RO NTs, STEM-EDX spectra, XRD pattern of washed samples, XRD pattern after lithiation, Nyquist and Bode plots of the impedance data, surface film resistances, determination of Warburg coefficients and determination of crystallite size from XRD. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author * E-mail: [email protected] Author Contributions



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All authors have given approval to the final version of the manuscript.

ACKNOWLEDGMENT We thank the DFG (project KU 2397/3-1) for financial support. T. Götsch and E. Portenkirchner acknowledge funding by the Austrian Science Fund (FWF) via grant F4503-N16 and grant P29645.

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