Hydrogenation Synthesis of Blue TiO2 for High-Performance Lithium

Apr 14, 2014 - Abstract. Abstract Image. Blue hydrogenated rutile TiO2 nanoparticles (blue TiO2) are prepared by treating white rutile via an enhanced...
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Hydrogenation Synthesis of Blue TiO2 for High-Performance LithiumIon Batteries Jingxia Qiu,† Sheng Li,† Evan Gray,‡ Hongwei Liu,§ Qin-Fen Gu,∥ Chenghua Sun,⊥ Chao Lai,† Huijun Zhao,† and Shanqing Zhang*,† †

Centre for Clean Environment and Energy, Environmental Futures Research Institute, and Griffith School of Environment, Gold Coast Campus, Griffith University, Parklands Drive, Southport, QLD 4222, Australia ‡ Queensland Micro- and Nanotechnology Center and School of Biomolecular and Physical Sciences, Griffith University, QLD 4111, Brisbane, Australia § The Australian Center for Microscopy & Microanalysis, The University of Sydney, NSW 2006, Australia ∥ Australian Synchrotron, Clayton, VIC 3168, Australia ⊥ School of Chemistry, Monash University, VIC 3168, Australia S Supporting Information *

ABSTRACT: Blue hydrogenated rutile TiO2 nanoparticles (blue TiO2) are prepared by treating white rutile via an enhanced hydrogenation process (i.e., high pressure and temperature). The materials characterization results demonstrate that the hydrogenation process leads to the increase in the unit cell volume and decrease in the size compared with the untreated white TiO2. The electrochemical impedance spectra analyses and theoretical energy calculations using density functional theory (DFT) suggest that the hydrogenation process not only improves electronic conductivity due to the formation of oxygen vacancy in the hydrogenation process but also dramatically augments lithium-ion mass transport within the crystalline lattice due to the introduction of oxygen vacancy and crystalline dislocation. Because of these characteristics resulting from the hydrogenation process, the blue TiO2 based lithium ion batteries (LIBs) possess significantly higher energy capacity and better rate performance than the white TiO2 based LIBs. In particular, at the rate of 0.1 and 5 C (1 C = 336 mAh g−1), the discharge capacities of the blue rutile are maintained at ca.179.8 and 129.2 mAh g−1, while the capacities of the white TiO2 are just ca. 119.6 and 55.5 mAh g−1, respectively.

1. INTRODUCTION Lithium-ion batteries (LIBs) have been the most extensively used power sources for electronic devices, such as mobile phones, laptops,and video cameras and possess the greatest potential for electric vehicle (EV) and hybrid electric vehicle (HEV) because of their high energy and power density and they are lightweight, have no memory effect and are environmental benignity. Titanium dioxide (TiO2) has been widely studied for LIBs as anode materials, because it is abundant, low cost, and nontoxic and it has a good cycle capability and good compatibility with most common electrolytes due to its suitable working voltage (ca. 1.5 V vs Li/Li+).1,2 Among the common TiO2 polymorphs, i.e., anatase, rutile, TiO2(B), and Brookite, rutile is of particular interest because it is the most commonly available natural form and the most stable form thermodynamically.3 Also, nanosized rutile TiO2 overcame the poor capacity problem of the conventional microsized rutile TiO2.3−7 Furthermore, its special safety feature bestows it a promising anode materials for LIBs.5 However, low rate capacity of the rutile TiO2 anode limits it from being used in LIBs on a large scale mainly due to two © 2014 American Chemical Society

barriers, poor electronic conductivity and inefficient lithium diffusion within the rutile lattice.3,7 To address these drawbacks of the TiO2 battery materials, strategies such as doping with foreign atoms, using different nanostructures, and conductive coatings can be used to enhance the performance of TiO2. The production of (001) faceted TiO2 nanosheets is also an effective way to improve the LIB performance by enhancing Li-ion diffusion along the c-axis in anatase TiO2.8 However, this method involves the use of toxic hydrofluoric acid and needs long hydrothermal reaction time (24h). In recent years, due to the rapid development of synthesis and application of nanostructured TiO2 as well as effectiveness and simplicity of hydrogenation for TiO 2 functionalization, hydrogenation has been attracting enormous attention. Hydrogenation processes can be conducted either under atmospheric pressure hydrogen mixture environment (e.g., 5% H2 and 95% Ar under ca. 1 atm at ca. 450 °C)1 or high Received: February 20, 2014 Revised: April 1, 2014 Published: April 14, 2014 8824

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energy profile is obtained using the Dimer method:17 a series of images are generated between the initial and final states, and for each image, the distance between Li and the oxygen at the edge of the first octahedral unit is fixed during the optimization. Through full relaxation, transition states (TSs) for Li+-diffusion are identified from the total energies of all images. All geometry optimization and energy calculations were carried out under the scheme of spin-polarized density functional theory (DFT), with the employment of generalized-gradient approximation (GGA)18 and the exchangecorrelation functional of Perdew−Burke−Ernzerhof (PBE),19 as implemented in the Vienna ab initio simulation package (VASP).20 The reciprocal space has been spanned by a planewave basis with a cutoff energy of 380 eV, and k-space is sampled by the gamma point. To correct the on-site electron correlation, DFT plus Hubbard model (DFT+U)21 has been employed, with U = 4.0 eV based on our tests and early publications.22 2.4. Electrode Preparation and Electrochemical Test. A slurry was prepared by mixing 80 wt % blue TiO2 powder or white pristine TiO2 powder as anode material, 10 wt % acetylene black as conducting agent, and 10 wt % poly vinylidene fluoride (PVDF) as binder and a certain amount of N-methylpyrrolidinone (NMP) as solvent. The mixed slurry was coated uniformly onto a thin copper foil, dried in air and subsequently dried in vacuum at 120 °C for 12 h. Pure lithium metal foil was used as the counter and reference electrode. A polypropylene microporous film was used as the separator. The electrolyte consisted of 1 M LiPF6 in a 1:1 (w/w) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC). Coin cell batteries (CR2032) were assembled using the above materials in an argon-filled glovebox. The cells were charged and discharged over a voltage range of 1.0−3.0 V vs Li/Li+ at room temperature (23 ± 0.5 °C) at different current density with a LAND-CT2001A battery tester (Wuhan, PRC). Cyclic voltammetry (CV) was performed on a CHI660D electrochemical workstation (CH Instrument, Shanghai, PRC).

pressure pure H2 environment (e.g., 20 bar H2 at ca. 200 °C).1,9−11 It is well-established that the hydrogenation process can produce abundant oxygen vacancy and Ti3+ that facilitate better electrical conductivity.12 Based on this feature, hydrogenation of anatase TiO2 has been used to improve the performance of anatase-based photocatalysis,13 photoelectrochemical sensors,14,15 supercapacitor,10 and LIBs.1,11,16 In contrast, the effect of hydrogenation on electrical conductivity and Li-ion mass transport of rutile TiO2 nanoparticles for LIBs is still not reported for LIB application. In this work, enhanced hydrogenation (high H2 pressure and high reaction temperature) is used as a facile and effective strategy to hydrogenate the rutile nanoparticles. It is expected that the proposed approach is able to significantly improve the LIB performance of the rutile by boosting the electronic conductivity and Li-ion diffusion efficiency within the TiO2 crystalline.

2. EXPERIMENTAL SECTION 2.1. Materials Preparation. Pristine rutile TiO2 powder (99%, 40 nm, Nanostructured & Amorphous Materials Inc. USA) was used as received. Before the hydrogenation reaction, 0.5−1 g TiO2 was first put into a glass vial then into highpressure stainless steel reactor. The pressure of the reactor was vacuumed to 0.039 bar and the temperature was heated up to 200 °C for 30 min to remove moisture and any impurities and activate the rutile sample. Subsequently, the hydrogen pressure was raised to 40 bar at 10 °C/min and the temperature was maintained at 450 °C for 1 h. 2.2. Materials Characterization. The transmission electron microscopy (TEM) images, high-resolution transmission electron microscopy (HRTEM) images, and electronic diffraction patterns were taken on a JEOL-2200FS instrument with a field emission gun, using an accelerating voltage of 200 kV. X-ray photoelectron spectroscopy (XPS, Thermo Escalab 250, monochromatic Al Kα X-ray resource), all binding energies were referenced to the C1s peak (284.8 eV) arising from adventitious carbon. Raman spectra were recorded at room temperature on a Renishaw system 100 Raman fiber spectrometer. The 785 nm excitation laser was used as the excitation source. The effect of the hydrogenation on the crystalline structures was studied on an in situ high-resolution synchrotron powder X-ray diffraction (HRXRD) using Powder Diffraction Beamline (10BM-1) at the Australian Synchrotron. For phase identification and structure determination, samples were loaded into predried 0.7 mm quartz capillary tubes and sealed into an inhouse designed flow cell (Figure S1) for X-ray diffraction measurements. The sample contained in a capillary flow cell was heated under 40 bar high purity H2 gas (99.99%) with a Cyberstar hot air blower. The sample contained in flow cell was flushed with high purity H2 gas 3 times before heat up. The samples were heated from 30 to 450 °C at a heating rate of 50 °C min−1. Data were collected in 10 °C steps for 10 min at each step. 2.3. Computational Method. In our theoretical calculations, TiO2 lattice is modeled using a 2 × 2 × 2 supercell with a molecular formula of Ti12O36. An Li atom is introduced randomly at the octahedral interstitial void. Oxygen vacancy is generated through removing one lattice oxygen between two interstitial voids in the hydrogenation process. Li-diffusion along the ab-plane is simulated through shifting the Li atom from one octahedral interstitial void to its neighboring one. The

3. RESULTS AND DISSUSION 3.1. Materials Characterization. In order to achieve the enhanced hydrogen penetration into the TiO2 lattice and faster hydrogenation reaction rate, in this work, pristine white rutile TiO2 (namely white TiO2) was treated under high-pressure pure H2 (40 bar) and high temperature (450 °C) for 1 h to produce blue hydrogenated TiO2 (namely blue TiO2). In order to in situ observe the physical and crystalline structural changes during the hydrogenation process, in situ synchrotron XRD facility (see Figures 1 and S1) were used. With this facility, we could visually observe the white TiO2 powder turned into blue color after the flame heater is turned on for just 5 min. The blue color became more intense with increased reaction time and subsequently steady after 30 min. In order to ensure the H2 has sufficient time to penetrate deeply into the TiO2 lattice, the hydrogenation reaction time had been maintained for 1 h. The synchrotron XRD patterns of the blue and white TiO2 in Figure 1a indicate that all of the samples showed good crystallinity and that all samples consist of only rutile and the hydrogenation process does not change the rutile framework. Rietveld refinement for all samples was conducted using starting structure model from ICDD database (2012) by TOPAS4.2 (Bruke, Germany). The peak profile was defined with fundamental parameters approach, the nanocrystalline size 8825

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The color change of the TiO2 during the hydrogenation process is commonly ascribed to oxygen vacancy, surface disorder, and oxygen deficiency.1 Interestingly, the blue TiO2 is changed into gray color instead of the original white color when it is exposed to the air. This suggests that oxygen vacancy can be formed at the surface as well as in the TiO2 lattice due to the enhanced hydrogenation experiment conditions. This can be demonstrated by spectroscopic analysis using UV−vis diffraction spectroscopy as shown in Figure 1b. The blue, gray, and white TiO2 have similar spectrum in UV region (200−400 nm) but have dramatic difference in the visible light range (400−800 nm). Both the white and blue TiO2 absorb UV light, which is determined by the inherent bandgap of rutile (3.2 eV). The similarity of UV spectra of the white and blue TiO2 suggests that oxygen vacancy at the surface has very trivial effect on the UV absorption. These TiO2 samples behave very differently in the visible light region. The blue TiO2 (blue curves in Figure 1b) possess the largest absorption in the visible light region among these samples. The gray rutile also has significant visible light absorption (green curve in Figure 1b). The visible light absorption can be attributed to the creation of midgap levels between the conduction band and valence band of rutile.9 Nevertheless, the blue color can be resulted from oxygen vacancy at the surface and in the bulk of the rutile crystalline while the gray color can be attributed to the oxygen vacancy in the bulk. Oxygen vacancy and Ti3+ at the rutile surface are sensitive to atmospheric air, leading to no distinguishable Ti3+ was observed for the blue TiO2 and white TiO2 in XPS spectra (Figure 1c). It is to note that the formation mechanism of Ti3+ and oxygen vacancies in our work is different from recently reported works, in which the Ti3+ or oxygen vacancies are generated under vacuum due to nanoscale thermodynamics.26,27 The blue and white TiO2 samples were systematically investigated using TEM and HRTEM. In Figure 2, A1 and A2 are TEM bright field images showing that the blue and white TiO2 samples have typical nanorod morphology with an average width of ca. 10 nm and length of ca. 40 nm, respectively. Their corresponding electron diffraction patterns (EDPs) are shown in B1 and B2 which could all be indexed with parameters of tetragonal phase rutile TiO2. C1 and C2 are HRTEM images showing lattice fringes of the blue and white TiO2, insets of the corresponding fast Fourier transform (FFT) images show the orientation of the crystal, in particular, the FFT images are indexed as [110] and [111]. Their inverse fast Fourier transform (IFFT) images show good lattice fringes in D1 and D2 and good crystalline state. The HRTEM image of the blue TiO2 nanoparticles gives a grain boundary image. Interestingly, small angle grain boundaries in the blue TiO2 sample as shown in D2 are very common. HRTEM image of C2 was taken under [111] direction. Grains 1 and 2 have the same projection direction of [111] but the closely packed plane (110) has a small rotation angle of about 8.5°. This induced crystalline dislocation is marked in Figure 2 D2. The grain boundary can be indexed as (9−8−1) which means that the coincident site lattices (CLS) have a Σ number of 146. It could be concluded that the white TiO2 exhibits better crystallinity while the blue one contains more crystal defects, such as dislocation lines and grain boundaries. The dislocations at grain boundaries could be an ion segregation site that might play an important role in Li-ion diffusion efficiency in intercalation/ extraction processes.

Figure 1. Synchrontron XRD spectra (a), UV−vis (b), and XPS diffraction spectra (c) of the blue TiO2 and white TiO2.

was derived from whole pattern modeling method.23,24 The refined unit cell lattice parameters and crystalline size results are summarized in Table 1. The changes along a and b directions Table 1. Unit Cell Parameters of Pure and Hydrogenated TiO2 Nanocrystals Derived from Synchrotron XRD Data (Figure 1a) samples white TiO2 blue TiO2

cell parameter a, b/Å (a = b)

cell parameter c/Å

cell volume/ Å3

crystalline size (nm)

4.612 30

2.952 47

62.8088

8.774

4.611 07

2.971 47

63.1793

7.650

were very trivial, but significant expansion along the c direction was observed. This leads to significant expansion of unit cell volume. The phenomenon of crystalline lattice expansion upon the hydrogenation was also observed by Shin et al.1 and Lu et al.11 Interestingly the average size of the smallest dimension (i.e., the width of the nanorod) calculated from the synchrotron XRD became smaller after the hydrogenation process (see Table 1). This is in line with the observation of SEM (Figure S2) and is consistent with the observation in the literature.25 8826

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Figure 3. Crystal structure of the rutile TiO2. (a) TiO6-octorhedron is represented with blue diamonds. The oxygen atoms sites 1 and 3 have two crystallographic equivalent oxygen atoms but overlapped viewed at [001] direction. Oxygen vacancy planes are possibly located at (300) and (030). (b) Raman spectra of the blue and white TiO2; (c) XRD patterns of the as-prepared blue and white TiO2 with Cu Kα as source.

Figure 2. TEM and HRTEM analysis of the white (left column) and blue (right column) TiO2. A1 and A2: TEM bright field images for the white and blue TiO2, respectively; B1 and B2: the corresponding EDPs of the white and blue TiO2, respectively; C1 and C2: HRTEM images and the corresponding FFT images (inset) of the white and blue TiO2, respectively; D1 and D2: IFFT images of the white and blue TiO2, respectively.

process. First, hydrogen atoms or ions penetrate and are trapped into the inner space of TiO2 unit cell but closer to frame oxygen atoms at the positions 2 and 4 which fall into plane {330). (Here, the plane {330) means the plane group includes both (330) and (3−30).) Second, an oxygen atom at the crystallographic equivalent sites 1 and 3 (Figure 3a) reacts with hydrogen, leaving an oxygen vacancy behind. The aforementioned changes can be reflected by the Raman spectra in Figure 3b and XRD in Figure 3c. The red curve in Figure 3b exhibits typical four Raman vibration peaks of the rutile nanocrystalline at wavenumbers of 143 (unknown mode), 230(B1g mode), 440(A1g mode), and 610 cm−1 (B2g mode), which is consistent with the literature.29 After the hydrogenation process, the peaks at 143, 230 cm−1 are maintained while the peaks at 440 and 600 cm−1 decrease dramatically. The latter peaks are resulted from O−Ti−O symmetric stretch

The crystal structure of rutile TiO2 has been determined to be isomorphous with casseterite (SnO2) by Vegard.28 It is based on a tetragonal lattice with the space group D144h and the unit cell contains 2 molecules. In order to address the role of hydrogenation on the crystalline structure of the rutile samples, the crystal structure of the rutile TiO2 with labeled oxygen atoms number is exhibited in Figure 3a. TiO6-octorhedron is represented with semitransparent blue diamond. The oxygen atoms sites 1 and 3 have two crystallographic equivalent oxygen atoms but overlapped viewed at [001] direction. There might be two changes to the rutile unit cell as a result of the hydrogenation 8827

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Figure 4. Typical first cycle charge/discharge profiles (A), galvanostatic tests at 0.05 C for 100 cycles (B), multicurrent density galvanostatic tests (C), and Nyquist plots and the equivalent circuit (D) of the blue and white TiO2. (1 C = 336 mAh g−1) The insertion of D is the modified equivalent circuit model of the system, Cf and Rf are surface layer capacitance and surface layer resistance of SEI film, respectively. Rs represents the electrolyte resistance, Rct, is associated with the charge-transfer resistance, Cdl represents the double-layer capacitance, Zw is associated with the Warburg impedance, reflecting the lithium diffusion process.

capacity and efficiency can be attributed to the aforementioned increase in unit cell volume, decrease in particle sizes and crystalline dislocation induced by the hydrogenation process. In Figure 4b, the blue TiO2 demonstrates a stable cycle performance and the discharge capacity can be retained at 179.8 mAh g−1 after 100 cycles, while the white TiO2 shows a significantly lower value, i.e., only about 118.1 mAh g−1. The high rate capability and cycle stability of the blue white TiO2 can be demonstrated by the variations in discharge capacity at different rates in Figure 4c. With no exception, the blue TiO2 samples demonstrated higher rate capacity than the white TiO2 at all tested rates. Furthermore, the capacity difference of the blue and white TiO2 samples increases with the increase of rates. In other words, the advantages of the blue TiO2 sample over the white TiO2 are more significant at higher rates. For example, the discharge capacity for the blue TiO2 was about 199.9 mAh g−1 after 10th cycle at the low rate of 0.05 C, and retained ca. 129.2 mAh g−1 at 5 C rate. In contrast, the capacity for the white TiO2 was ca. 150 mAh g−1 at the 0.05 C rate and preserved only 55.5 mAh g−1 at the 5 C rate. The difference between them was ca. 50 mAh g−1 at 0.05 C while ca. 75 mAh g−1 at 5 C rate. Most importantly, the blue TiO2 sample can endure more radical changes of current densities to maintain better stability after such changes than the white TiO2 samples. In particular, after the 60th cycle using the successively increasing current densities, the blue TiO2 sample retained a discharge capacity of 192.9 mAh g−1 (i.e., 96.5% retention percentage) when the rate was returned to 0.05 C. This is practically an invaluable merit for the blue TiO2 anode materials in that this bestows the blue TiO2 sample a long cycle life and abuse tolerance of higher power LIBs.

vibration while Ti atom is at rest. In particular, the peak at 440 cm−1 is attributed to the vibration of bond O1−Ti−O3 while the peak at 600 cm−1 is attributed to the vibration of bond O2− Ti−O4. In the symmetric vibration, atoms 1 and 3 move toward each other and atoms 2 and 4 move away from each other, while in the other mode, the two pairs move in or out in phase. The descending vibration intensity of the bond O1−Ti− O3 and bond O2−Ti−O4 suggests that these two bonds have been disturbed by the disappearance of oxygen atom (oxygen atom are removed from its site by the hydrogen). The vocation very likely occurs at bond O1−Ti−O3 but not O2−Ti−O4. In Figure 3c, no significant difference of XRD patterns between the blue and white rutile samples, except that slightly enhanced diffraction peak intensity at 2 theta of 63°, can be observed for the blue TiO2 samples. This enhancement suggests that the oxygen vacancy very likely locates in ordered fashion at the planes (030) and (300) since O1 and O3 are the most possible vacancy. 3.2. Electrochemical Performance. The blue and white TiO2 samples were used as anode materials to fabricate LIBs under identical conditions. Figure 4a shows the initial discharge−charge curves of the samples at the rate of 0.1 C (1 C = 336 mAh g−1). Different discharge behavior can be observed. The discharge and charge capacity for the white TiO2 was just about 223.6 and 164.2 mAh g−1, respectively. In contrast, for the blue TiO2, a significantly higher initial discharge and charge capacity were 269.7 and 222.8 mAh g−1, respectively. Interestingly, an obvious potentials plateau is presented around 1.4 V for the blue TiO2 only and the blue TiO2 sample presents a higher coulombic efficiency than the white TiO2. The improved performance of the blue TiO2 in 8828

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coefficient of Li-ions in the rutile samples. Due to the existence of the oxygen vacancies, the Ti3+ shall locate in the center of the rutile octahedral crystal. Due to the loss of the oxygen and the increased space permission, Li-ions tend to fill the physical vacancies. However, the vacancies are very close to the Ti3+ cores that repel the positive charged Li-ions immediately and pass the Li-ions to other Ti3+ cores and so on. This electrostatic repellent force as a driving force would accelerate the mass transport of Li-ions within the crystalline structure. To investigate the effect of oxygen vacancy on the mass transport of Li-ions within the crystalline structure, the energy profiles for Li-ion diffusion have been obtained by DFT study, for both perfect TiO2 lattice and TiO2 with oxygen vacancies (see Figure 5). The model is shown in Figure 5a, with TiO2

The significant boosted performance of the blue TiO2 LIBs over the white TiO2 one can be attributed to two reasons, the boosted electronic conductivity (i.e., reduced impedance) and augmented TiO2 Li-ion diffusion in the insertion/extraction process. In order to investigate the effect of hydrogenation on electronic conductivity, 0.100 g blue and white TiO2 nanoparticles are compressed into two round pallets under 15 tons pressure and their conductivities are measured using a fourpoint probe conductivity measurement device, respectively. The i−V curve (Figure S3) of the blue TiO2 shows a much larger slope than that of the white TiO2 sample. In fact, the slope of the white TiO2 is close to zero. Because the slope is directly proportional to electronic conductivity, this observation qualitatively suggests that the electronic conductivity of the blue TiO2 is significantly larger that of the white TiO2. The electrochemical impedance spectroscopy (EIS) was used to investigate the ionic and electronic resistance behavior in the battery enclosure. Corresponding Nyquist plots of the blue and white TiO2 samples were obtained under the identical conditions in Figure 4d. The blue TiO2 sample illustrated two depressed semicircles at high frequency and therefore the corresponding equivalent circuit mode is shown in the inset. It is well-established that the Nyquist plot at high and medium frequency (above 0.2 Hz) is corresponding to the overall internal resistance which represents the electrolyte resistance (Rs), and the charge-transfer resistance (Rct),30 that were calculated using the equivalent circuit in Table 2. Due to the

Figure 5. Computational study of Li-ion diffusion in the TiO2 structure with and without oxygen vacancies. (a) Model used in the calculation of energy profiles for Li-ion diffusion; (b) energy profiles for Li-ion diffusion from one octahedral interstitial void to its neighboring one.

Table 2. Physicochemical Properties Measured and Calculated from EIS Spectra (in Figure 4D) samples white TiO2 blue TiO2

DLi (EIS) (cm2·s−1) −15

4.73 × 10 1.20 × 10−14

Rs (Ω)

Rct (Ω)

5.69 4.11

98.02 27.23

framework as stick and Li as purple spheres. Two neighboring octahedral interstitial voids distributed at two different ⟨001⟩ channels are indicated by the dashed green lines, together with the tetrahedral gap indicated by blue lines. Early experimental and theoretical studies have clearly demonstrated that the onedimensional diffusion along the ⟨001⟩ channels is ultrafast with a small barrier.21−23 While along the ab-plane, the diffusion is very slow due to the strong attractive O−Li interaction. Figure 5b shows the calculated barrier for Li-diffusion in the ab-plane from the left octahedral interstitial void to its neighboring one through the path indicated by the arrow with a barrier more than 2.00 eV, being consistent with the measured low diffusion coefficient. However, if oxygen vacancies are presented between the two octahedral interstitial voids (indicated as a black circle in Figure 5a), the diffusion barrier can be reduced to 1.23 eV, which vividly demonstrates that generating oxygen vacancies can promote the Li-diffusion in the ab-plane. Experimentally, if TiO2 samples have been exposed to high-pressure hydrogen, a part of hydrogen (e.g., via being exposed to the air or heat treatment) may diffuse to the surface and even leave from the surface, which can facilitate the formation of oxygen vacancies and thus promote the Li-diffusion.

use of identical electrolyte system, there is small difference in Rs values between two samples. In strong contrast, Rct value of the pristine white TiO2 sample (98.02 Ω) decreased dramatically to 27.23 Ω of the blue TiO2 after the hydrogenation process. All the measurement confirmed that the enhanced hydrogenation process could significantly improve the electron transport conditions of the white TiO2. Furthermore, the linear Warburg regions at low frequency (from 0.2 to 0.01 Hz) in Nyquist plots can be used to represent the Li-ion diffusion behavior within the solid-state electrode materials.30 The diffusion coefficient of Li+ also calculated by the following formula (eq 1): DLi =

2 1 ⎡⎛ Vm ⎞ dE ⎤ ⎢⎜ ⎟ ⎥ 2 ⎢⎣⎝ FAσw ⎠ dx ⎥⎦

(1) 3

where Vm is the molar volume of rutile TiO2 (18.88 cm /mol), F is the Faraday constant (96485 C/mol), A is the total contact surface area between the electrolyte and the electrode, and σw is the Warburg prefactor obtaining from the Warburg region of impedance response. The results in Table 2 show that the Liion diffusion coefficient is ca. 1.20 × 10−14 cm2·s−1 in the blue TiO2 while it is about ca. 4.73 × 10−15 cm2·s−1 in the white TiO2. In other words, the hydrogenation process has enhanced the mass transport within the TiO2 lattice by a factor of 3 times. 3.3. DFT Calculation. DFT calculation was used to explain the mechanism responsible for the enhanced diffusion

4. CONCLUSION An enhanced hydrogenation process was used to prepare the blue rutile TiO2 as anode for LIBs for the first time. The hydrogenation process increases the unit cell volume of the pristine rutile crystalline, reduces the size of the rutile nanoparticles, causes crystalline dislocation, and produces oxygen vacancy throughout the bulk and surface of the crystalline. These changes facilitate fast lithium-ion mass 8829

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transport and electron transfer during the charge/discharge process. The enhanced Li-ion mass transport mechanism was verified by EIS measurement and DFT calculation. Consequently, the as-prepared blue TiO2 possesses significantly enhanced energy capacity, rate performance, and operational stability in comparison with pristine rutile.

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ASSOCIATED CONTENT

S Supporting Information *

Color change of the sample by in situ synchrotron XRD anylysis (Figure S1); SEM images of the blue and whiteTiO2 (Figure S2); i−V current measurement of the blue and white TiO2 pallets (Figure S3). This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*Tel.:61-7-55528155. Fax: 61-7-55528067. E-mail: s.zhang@ griffith.edu.au. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge the financial supports from Australia Research council and Griffith University, and the facilities and the scientific and technical assistance of the Australian Microscopy & Microanalysis Research Facility at the Center for Microscopy and Microanalysis at the University of Queensland.

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REFERENCES

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dx.doi.org/10.1021/jp501819p | J. Phys. Chem. C 2014, 118, 8824−8830