New Insights on the Reversible Lithiation Mechanism of TiO 2 (B)

Subscriber access provided by ISTANBUL UNIVERSITESI

Article

New Insights on the Reversible Lithiation Mechanism of TiO(B) by Operando X-ray Absorption Spectroscopy and Xray Diffraction Assisted by First Principles Calculations 2

Marcus Fehse, Mouna Ben Yahia, Laure Monconduit, Frederic Lemoigno, Marie-Liesse Doublet, Florent Fischer, Cecile Tessier, and Lorenzo Stievano J. Phys. Chem. C, Just Accepted Manuscript • Publication Date (Web): 05 Nov 2014 Downloaded from http://pubs.acs.org on November 5, 2014

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

The Journal of Physical Chemistry C is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 29

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

The Journal of Physical Chemistry

New Insights on the Reversible Lithiation Mechanism of TiO2(B) by Operando X-ray Absorption Spectroscopy and X-ray Diffraction Assisted by First Principles Calculations Marcus Fehse,∗,† Mouna Ben Yahia,†,‡ Laure Monconduit,†,‡ Frédéric Lemoigno,†,‡ Marie-Liesse Doublet,†,‡ Florent Fischer,¶ Cécile Tessier,¶ and Lorenzo Stievano∗,†,‡ ICGM CNRS and Université Montpellier 2, Place E. Bataillon, 34095 Montpellier, France, Réseau sur le Stockage Electrochimique de l’Energie, CNRS FR3459, France, and SAFT, Direction de la Recherche, 111-113 Bd Alfred Daney, 33074 Bordeaux, France E-mail: [email protected]; [email protected] Phone: +33 (0)4 67 14 33 46. Fax: +33 (0)4 67 14 33 04

Abstract Operando X-ray absorption spectroscopy (XAS) and X-ray diffraction (XRD) measurements provide new insights on the mechanism of lithium insertion into TiO2 (B). The investigation of the evolution of electronic, long range, and local structure during electrochemical cycling, indicates a purely monophasic insertion mechanism upon ∗

To whom correspondence should be addressed ICGM CNRS and Université Montpellier 2 ‡ RS2E ¶ SAFT Direction de la Recherche †

1 ACS Paragon Plus Environment


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

lithium insertion, while global and local structure are only slightly modified. While XRD reflects an anisotropic lattice expansion, EXAFS reveals a wide distribution of Ti-O bond length, in line with the presence of two distinct distorted octahedral Ti environments, in agreement with previous DFT calculations. Upon lithium insertion, these Ti-O coordination shells undergo significant modifications which are enhanced once the insertion of 0.4 Li is exceeded, connoting a two regime process which is in good agreement with the electrochemical signature of this material. DFT calculations and local chemical bond analyses were coupled with experimental results, thus providing additional insights into the structural response of TiO2 (B) upon lithiation.

Keywords LIB, Anode, EXAFS, XANES, DFT calculations

Introduction In the search for suitable high power battery negative electrode materials, the bronze phase also known as TiO2 (B) is a potential candidate with distinct advantages over other titaniabased insertion materials. While its increased theoretical capacity (336 mAh/g), compared to commercialised LTO (Li4 Ti5 O12 , Captheo 175 mAh/g), has been known for quite some time, 1–3 its superior rate capability compared to the largely studied anatase phase has been more recently claimed in various papers. 4–8 These assets, in combination with the intrinsic benefits of titania-based materials, such as environmental benignancy, cost efficiency, long cycle life and high safety due to elevated insertion potential, endow TiO2 (B) to be a promising alternative negative electrode material. 9 Despite the numerous experimental studies showing favourable insertion properties of TiO2 (B), there are comparatively few works dedicated to the understanding of the physicochemical foundations of such performances. In a recent publication, we showed that the


Page 2 of 29

Page 3 of 29

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


superior rate capability of TiO2 (B) can be attributed to lower internal resistance at increased cycling rate compared to other more common TiO2 polymorphs such as anatase. 8 Other research groups, on the other hand, attribute this superiority to the low density or to the unique open channel structure. 3,10–12 Zukalová et al. were the first to introduce the concept of pseudocapacitive storage for describing the Li+ storage mechanism in TiO2 (B). 5 They describe a faradaic surface storage mechanism which is not restricted by slow solid-state diffusion. Such a mechanism is enabled by the monoclinic C2/m open channel structure which allows not only fast Li+ transport, but also provides easily accessible Li+ insertion sites within the channels. 5,13 Armstrong et al. studied the Li coordination sites in a partially lithiated sample with average composition Li0.15 TiO2 (B) by neutron diffraction. 14 This sample was found to be a disordered mixture of TiO2 (B) and of the partially lithiated phase Li0.25 TiO2 (B), where Li+ occupies only the channel site, C. These data where backed-up by DFT calculations, which also suggested the higher stability of the C sites compared to the other two, A1 and A2, located on the (001) plane and between the oxygen atoms bridging adjacent layers in the (001) plane, respectively. However, they also found that upon further lithiation the C insertion site becomes less and less stable compared to the other two. More recently, the assertion that these C-sites are the energetically preferred lithium insertion ones, has been questioned by the comprehensive computational study of Morgan & Madden. 6 In the present work, we confirm the study of Morgan & Madden by directly correlating the titanium local distortions to the site preference and to the structural changes expected. It is challenging to answer this question without the support of experimental studies, since various parameters, such as the degree of filling, 14 or the morphology, 15 are claimed to have an effect on the stability of the different insertion sites. In one of the few experimental studies aiming at elucidating the insertion mechanism of TiO2 (B), and conducted by ex situ XAS, Okumura et al. claim that Li storage in nanostructured TiO2 (B) occurs in the space charge layer (SCL). Therefore it neither inflicts changes



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

in lattice nor in the electronic structure resembling more to a fully capacitive double layer storage than to a pseudocapacitive faradaic storage. 16 For bulk TiO2 (B), however, they suggest a two regime insertion mechanism with preferred C site filling, 17 in line with the cited results of Armstrong et al. 14 which have been questioned by more recent publications. 6,18,19 In their recently published ex situ NMR study Hoshina et al. suggest that surface storage in nanoscaled (Ø ≈ 100 nm) TiO2 (B) occurs only up to 0.18 Li. 20 Upon further lithiation lithium is stored primarily in A2 site. Furthermore they claim that lithiation exceeding 0.45 Li leads to the filling of A1 site causing a growing local lattice distortion eventually impeding Li diffusion. In our recently published work we demonstrate that by coupling measurement results from operando XRD and XAS a reliable and comprehensive picture of the lithiation and delithiation process of TiO2 anatase can be obtained. 21 In this work we present a thorough study by operando XAS and XRD of the insertion of lithium in TiO2 (B). We show insights on the evolution of electronic, local, and lattice structure during lithiation in order to elucidate the physico-chemical mechanisms and processes that are the basis for the favourable cycling properties of TiO2 (B). These experimental results are supported by local chemical bond analysis investigated within the conceptual DFT framework to shed light on the debated and still unclear energetics of lithium insertion sites of the TiO2 (B) structure.

Experimental Details Synthesis and electrode preparation The TiO2 samples were synthesised via hydrothermal synthesis, a detailed description has been previously published elsewhere. 8 Briefly, TiO2 powder was dispersed in an alkaline aqueous solution and kept in an autoclave for 72 h at 150 ◦ C. After cooling to room temperature, the material was treated with HCl for ion exchange, and then washed with deionised water upon total removal of chlorides. Finally, the samples were calcined in air at 450 ◦ C. 4 ACS Paragon Plus Environment

Page 4 of 29

Page 5 of 29

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


Self-supported electrodes were obtained by mixing TiO2 (B) with 8% Carbon black and 20% PVdF in an acetone based slurry, which was subsequently tape casted with a doctor blade on a smooth clean surface. After air drying, electrodes were cut out and pealed off from the smooth surface.

Methods Operando X-ray absorption spectroscopy and X-ray diffraction were performed using a customized cell with a sturdy stainless steel body and Beryllium window as both closure and current collector. 22 Self-supported electrode films prepared as described above were mounted in half cell configuration against Li metal using EC:PC:3DMC 1M LiPF6 and Celgard as the electrolyte and the separator, respectively. XRD measurements were carried out using a Philips X’pert diffractometer with Cu Kα radiation at room temperature, in the range from 20-50◦ 2 Θ using a step size of 0.05 ◦ and dwell time of 700 s yielding to the acquisition of about 1 diffraction pattern per hour. The computer program Powder Cell 23 as well as Fullprof 24 were used to deduce unit cell parameters. Full pattern matching was performed based on crystal structure published by BenYahia et al. 25 to refine the cell parameters a, b and c, but no atomic position. XAS measurements were carried out on beamline A (HASYLAB @ DESY, Hamburg, Germany). Ti K-edge spectra were recorded in the transmission mode in the range 48006100 eV , recording about three spectra per hour. Fourier transform were performed using k2 weighing and the structural parameters were determined by curve-fitting procedures using the IFEFFIT software suite 26 on the basis of the published TiO2 (B) crystal structure by BenYahia et al. 25 We included in the fit one or two two-leg contributions for 6 nearest oxygen atoms, and three two-leg contributions carrying 7 nearest Ti atoms. No three-leg contributions were considered. For each two leg contribution the position (delR) and the Debye-Waller factor (ss) were fitted individually. Electrochemical cycling was carried out in the galvanostatic mode on a Biologic mul5 ACS Paragon Plus Environment


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

Page 6 of 29

tichannel potentiostat. Measurements were carried out at cycling rate of C/20, where 1C corresponds to the insertion of 1 mol Li per hour into 1 mol TiO2 , and is equivalent to a current of 336 mA/g. Calculations were performed within the framework of Density Functional Theory (DFT) and its DFT+U variant 27 using the rotationally invariant method of Dudarev 28 as implemented in the Vienna Ab-Initio Simulation Package (VASP). 29,30 The generalized gradient approximation of Perdew, Burke, Ernzerhof (PBE) 31 was chosen for exchange and correlation energy. The valence/core interaction was described in the projected augmented wave formalism (PAW) 32 with Ti(2p) and Li(1s) included in the atomic valence shells. Convergence of the calculations was checked with respect to the energy cut-off (up to 800 eV ) and the k-points grid (2×6×4) leading to reliable converged structures. Atomic coordinates and lattice parameters were fully relaxed using conjugate gradient energy minimization with a force tolerance set to 1.10−3 eV/˚ A. Fukui functions were also computed for various Lix TiO2 compositions. Within the conceptual DFT framework, 33 they are defined as the change in the system electron density due to an added charge (electron or hole): ∂ρ(r) ∂N

±

f (r) =

!

(1) v

where N is the amount of added charge and v assumes no structural change between the neutral and charged systems. The f + (r) and f − (r) quantify the electronic response of the system to the addition of a fraction of electron and hole, respectively and are often used to probe the frontier orbitals of the system, namely the highest occupied (HO) and lowest unoccupied (LU) crystal orbitals. In contrast to spin density map plots, Fukui functions include the polarization effects induced by the change in density on all orbitals (not only the frontier orbitals) and are therefore meaningful descriptors for describing electrochemical properties.


Page 7 of 29

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


Results The here presented material is composed of TiO2 (B) rich nanorods with average diameter of 40-60 nm and a specific surface area of 50 m2 /g. A comprehensive characterization of this material can be found in our previous publication. 8

Operando XRD The evolution of the diffraction pattern of TiO2 (B) upon consecutive Li+ insertion (charge) and de-insertion (discharge) is presented in Fig. 1. The only reflection peak that cannot be attributed to TiO2 (B) at 46◦ corresponds to hexagonal Be originating from the Be window of the in situ cell. This feature can be used to calibrate the diffraction patterns.

Figure 1: Operando XRD measurement upon consecutive charge (Li+ insertion) and discharge (Li+ de-insertion) of TiO2 (B) self supported film, showing reversible peak shift of (110),(003),(601) and (020) all attributed to TiO2 (B), while peak at 46◦ results from Be window. A gradual shift of the Bragg peaks of TiO2 (B) is observed upon lithiation, indicative of a single phase transformation which is shown to be perfectly reversible in lithium deinsertion. The main changes are observed for the (020) and (110) reflection at 48.5 ◦ and 24.9 ◦ , respectively, the (60¯1) at 44.5 ◦ undergoes less pronounced position change. 7 ACS Paragon Plus Environment


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

Because of the low crystallinity of the TiO2 (B) samples and the short acquisition time induced by the operando set up, the obtained diffraction patterns are not suitable for a complete Rietveld refinement. A simple pattern matching could nevertheless be performed by following the shift of the main diffraction peaks TiO2 (B) with the computer program Powder cell, which provided a first estimate of the changes in the lattice parameters with Li content. The evolution of the lattice parameter as a function of the amount of inserted lithium is shown in Fig. 2. By accumulating multiple diffraction patterns of the fully lithiated and of the pristine state, we were able to obtain sufficiently good patterns to perform a full pattern matching for these two conditions, using much more performing computer program FullProf. The obtained data points are shown as hollow markers in Fig. 2, and follow well the trend of those obtained using the computer program Powder Cell.

Figure 2: Modification of a, b and c lattice parameters as a function of the amount of lithium inserted in TiO2 (B) calculated from the evolution of the XRD patterns using Powder Cell (filled markers) and full pattern matching with FullProf (hollow markers). After the reaction of ≥0.7 Li, an increase of the lattice parameters a and b of 2.9 and 4.5% is observed, respectively, while c remains virtually unchanged within the experimental error. This anisotropic expansion agrees well with the findings by Armstrong et al. 14 and with the orbital interpretation given below. Applying the full pattern matching method, the expansion in a and b is somewhat reduced to 1.7 and 3.6%, respectively. The discrepancy


Page 8 of 29

Page 9 of 29

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


between the two evolution methods can be explained by the variation of the angle β, which results of almost 1.5 % in the full pattern matching, whereas the angle could not be varied in the analysis with the computer program Powder Cell due to the small number of reflections used for the evaluation of the lattice parameters.

Operando XAS The modification of the Ti K-edge (4966 eV ) upon lithium insertion along with the shift of the Ti K-edge position throughout reduction and subsequent oxidation are presented in Fig. 3 (a) and (b), respectively. The changes in the pre-edge features are attributed to the filling of Ti 3d orbitals and changes in the Ti-coordination sphere (local symmetry of the TiO6 octahedra), while the change in position of the K-edge is directly connected to the oxidation state of the emitting Ti atom. 34–37 Upon lithium insertion, TiIV is reduced to TiIII which is expressed by a shift of the absorption edge to lower energies. It is noteworthy that this process is largely reversible, although at the end of charge, the starting point of the pristine TiO2 is not completely recovered. Such irreversibility has been previously reported for TiO2 based electrodes. 8,35,38–40 Accordingly it could be an indication of parasitic reactions leading to stable lithium compounds, or of partial lithium trapping in irreversible sites during the first Li+ insertion, thus avoiding pure TiIV recovery at the very end of charge.

Figure 3: (a) Evolution of Ti K-edge upon Li+ insertion and (b) reversible shift of Ti K-edge position during reduction and subsequent oxidation of TiO2 (B) rich sample. 9 ACS Paragon Plus Environment


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

The evolution of the EXAFS signal with lithium insertion, and that of the corresponding FT schematically representing the average distribution of the next neighbours to the absorbing Ti-centre, are presented in Fig. 4 (a) and (b), respectively. The black lines resemble the non lithiated state while with growing lithiation the colour fades to light grey. For the sake of clearness only four XAFS curves at different stages of lithiation have been selected in Fig. 4 (a). Based on the good quality of the EXAFS data, the FT was performed using a sinusoidal window in the range 2.5< k

New Insights on the Reversible Lithiation Mechanism of TiO 2 (B)

Recommend Documents