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Boron-Doped Anatase TiO as a High-Performance Anode Material for Sodium-Ion Batteries Baofeng Wang, Fei Zhao, Guodong Du, Spencer H. Porter, Yong Liu, Peng Zhang, Zhenxiang Cheng, Huakun Liu, and Zhenguo Huang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b03270 • Publication Date (Web): 03 Jun 2016 Downloaded from http://pubs.acs.org on June 4, 2016

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ACS Applied Materials & Interfaces

Boron-Doped Anatase TiO2 as a High-Performance Anode Material for Sodium-Ion Batteries Baofeng Wang1,2,*, Fei Zhao1, Guodong Du2 , Spencer Porter2, Yong Liu3, Peng Zhang2, Zhenxiang Cheng2, Hua Kun Liu2, Zhenguo Huang2,* 1

Shanghai Key Laboratory of Materials Protection and Advanced Materials in Electric Power, Shanghai

University of Electric Power, Shanghai, 200090, China 2

Institute for Superconductor and Electronic Materials, University of Wollongong, NSW, Australia

3

Laboratory of Nanoscale Biosensing and Bioimaging, School of Ophthalmology and Optometry,

Wenzhou Medical University, Wenzhou, Zhejiang, 325027, China ABSTRACT: Pristine and boron-doped anatase TiO2 were prepared via a facile sol-gel method and the hydrothermal method for application as anode materials in sodium-ion batteries (SIBs). The sol-gel method leads to agglomerated TiO2, whereas the hydrothermal method is conducive to the formation of highly crystalline and discrete nanoparticles. The structure, morphology, and electrochemical properties were studied. The crystal size of TiO2 with boron doping is smaller than that of the non-doped crystals, which indicates that the addition of boron can inhibit the crystal growth. The electrochemical measurements demonstrated that the reversible capacity of the B-doped TiO2 is higher than for the pristine sample. B-doping also effectively enhances the rate performance. The capacity of the B-doped TiO2 could reach 150 mAh/g at the high current rate of 2C and the capacity decay is only about 8 mAh/g over 400 cycles. The remarkable performance could be attributed to the lattice expansion resulting from B doping and the shortened Li+ diffusion distance due to the nanosize. These results indicate that B-doped TiO2 can be a good candidate for SIBs. KEYWORDS: sodium ion batteries, anode, anatase, TiO2, boron doping 1 INTRODUCTION The ever-increasing demand for energy necessitates the exploitation of energy from environmentally friendly and renewable resources such as solar and wind. The intermittent nature of these resources, however, demands an efficient energy storage system to level the energy output1. Lithium-ion batteries (LIBs) have been considered for this purpose due to their high energy densities, low toxicity, and long cycle life. The natural abundance of lithium, however, is quite low, and its distribution is highly confined to several countries, which has caused great concern about using LIBs in large-scale energy storage systems2-3. Sodium-ion batteries (SIBs) have thus come to the attention of researchers since sodium is much more naturally abundant and very well distributed around the globe. Compared with the well-established LIB techniques, however, the performance of SIBs in such aspects as cycle life and energy density needs to be dramatically improved for them to reach wide deployment4-7. One factor limiting the performance of SIBs is the anode, where several types of materials have been recently studied. These include hard carbon based materials8-9, layered compounds such as MoS210-11 and SnS212, and elements that form alloys with Na such as Sn13 and P14-15. While some interesting results are reported, overall, these materials still need improvement in certain aspects such as specific capacity, cycle life, reversibility, and rate capability. For example, Sn13 forms alloys with Na and

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delivers a large capacity of more than 800 mAh/g, but it suffers from large irreversible capacity loss in the initial cycles and poor cycling stability afterwards. Among the anode materials studied for SIBs, anatase TiO2 has recently aroused strong interest and has shown quite impressive performance in terms of both cycling stability and rate capability. TiO2 is nontoxic, inexpensive, and abundant, and is well known as an anode material for Lithium ion batteries. At first glance, the large ionic radius of the Na ion (0.102 nm) compared with that of the Li ion (0.076 nm) would make the intercalation and deintercalation of Na ions into TiO2 more difficult. Recent research on using TiO2 for SIBs has shown interesting results, however. Both commercial and synthesized (by the cellulose-based template method or hydrothermal method) nanoscale anatase TiO2 were studied for SIB anode16-20. Nanocrystalline TiO2 showed a reversible capacity of more than 150 mAh g−1 at low charge/discharge currents (36-50 mAh/g) with rather good cycling performance (up to 100 cycles)16, 18-19 .Conductive additives such as carbon have been shown to greatly improve the rate performance and capacity retention of TiO217,

20-21

. Doping has been employed to improve

electrochemical performance, since doping promotes ionic conductivity by enlarging or providing more open channels for ionic diffusion and charge transfer. Boron doping in TiO2 has been mainly studied in photocatalysis. Boron doping can enhance the photocatalytic activity of TiO2 due to or partly due to the formation of Ti3+ ions induced by oxygen vacancy which can increase the conductivity of TiO222-23. Recent research results have shown that the lithium electroactivity of both anatase and rutile TiO2 can be enhanced by boron doping24-25, while the sodium storage performance of boron-doped anatase TiO2 has not been studied. Herein, B-doped anatase TiO2 was prepared via a facile sol-gel method and the hydrothermal method as anode for SIBs. The impact of the preparation methods and the effects of doping on the electrochemical performance were investigated. The results demonstrate that B-doping effectively enhances both the capacity and the coulombic efficiency of TiO2. This improvement can be ascribed to the lattice expansion and the reduction of crystal size after B doping. B-doped TiO2 is therefore a promising anode candidate for SIBs with high capacity, long cycle life, and low cost. 2 EXPERIMENTS 2.1 Materials preparation B-doped TiO2 was synthesized via a modified sol-gel method using titanyl nitrate TiO(NO3)2 and H3BO3 (99.5%, Sinopharm Chemical Reagent Co., Ltd) as the precursors. In a typical preparation, TiO(NO3)2 was synthesized by dissolving TiO2 in 40 mmol HNO3 (65%, Sinopharm Chemical Reagent Co., Ltd). The TiO2 used for preparing TiO(NO3)2 was first obtained from the hydrolysis of 20 mmol tetrabutyltitanate (TBT, 98%, Sinopharm Chemical Reagent Co., Ltd) in ethanol/water solution. 8.4 g citric acid (99.5%, Sinopharm Chemical Reagent Co., Ltd) and 1 mmol H3BO3 were then added into the TiO(NO3)2 solution under vigorous stirring. After adjusting the pH to 7 using ammonium hydroxide, the solution was dried at 90 oC to obtain a gel, which was further dried at 250 oC for 3 h. The obtained solids were ground and finally calcined at 500 oC for 6 h (heating rate = 5 oC /min). The B-doped TiO2 sample obtained via this sol-gel method was denoted as B-S-TO. TiO2 without B doping (S-TO) was synthesized by the same procedure with no addition of B(OH)3. For comparison, TiO2 and B-doped TiO2 powders were also prepared via a typical hydrothermal procedure. Tetrabutyl titanate (10 mL, 98%), hydrofluoric acid (0.8 mL, 47%) solution, and a certain amount of (0.09 g) H3BO3 were mixed in a 100 mL dry Teflon autoclave. After being kept at 180 °C for 24 h, white powders were collected and washed several times with ethanol followed by distilled water.


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The powders were then calcined at 500 °C for 6 h in air to produce B doped anatase TiO2 nanoparticles. The obtained B-doped TiO2 and the pristine TiO2 synthesized by the same procedure were denoted as B-H-TO and H-TO respectively. 2.2 Physical characterization The crystal structure was characterized by X-ray diffraction measurements (XRD, GBC MMA) with Cu Kα radiation (λ = 0.1541 nm). Rietveld refinements were performed using GSASII26. The morphology of the powder was studied using a field emission scanning electron microscope (FESEM, JSM 7500F, JEOL). Transmission electron microscopy (TEM) studies of the samples were conducted using a JEM 2010 JEOL operated at 120 kV. The specific surface area of the powder was measured by the Brunauer-Emmett-Teller (BET) method using a surface area analyzer (NOVA 2200, Quantachrome). Elemental compositions and chemical states were analyzed using X-ray photoelectron spectroscopy (XPS, PHOIBOS 100 hemispherical analyzer SPECS, GmbH). 2.3 Electrode preparation and electrochemical testing Electrochemical measurements were performed using CR2032 coin-type cells assembled in an argon-filled glove box (M. Braun Co.). The composite cathode was made of the active materials (80 wt. %), acetylene black (10 wt. %), and carboxymethyl cellulose (CMC) binder (10 wt. %). All the materials were homogeneously mixed in water by planetary ball milling for 1 h at the speed of 500 r/min and then coating the slurry uniformly on copper foil. Finally, the electrodes were dried under vacuum at 100 °C for 10 h. Sodium foil was used as a counter electrode. The loading of active material on the cathode was about 2 mg cm-2. The electrolyte was 1.0 mol/L NaClO4 in an ethylene carbonate (EC) - diethyl carbonate (DEC) solution (1:1 v/v), with 5 vol.% addition of fluoroethylene carbonate (FEC). Galvanostatic cycling of the assembled cells was carried out using Land CT2001A testers (Wuhan, China) between 2.5 and 0.01 V at room temperature. An Autolab electrochemical workstation (302N) was used for cyclic voltammetry measurements with potentials between 0 and 2.0 V at a scan rate of 0.2 mV s−1 (all the voltage values given herein are referred to the Na/Na+ redox couple) and for electrochemical impedance spectroscopy (EIS) measurements with 5 mV ac signals and a frequency range from 100 kHz to 0.1 Hz. 3 RESULTS AND DISCUSSION

B-H-TO

Intensity / counts.

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


(101)

(004) (103) (112)

(200) (211) (105) (204)

H-TO B-S-TO

S-TO

20

30

40

50

60

70

2-Theta / degree


80


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

Figure 1 XRD patterns of pristine and B-doped TiO2 prepared by the sol-gel and hydrothermal methods. The XRD patterns of pristineTiO2 and B-doped TiO2 are shown in Figure 1. No diffraction peaks related to B2O3 are detected for the B-doped samples. Similar to pure TiO2, both B doped samples exhibit the characteristic diffraction peaks that can be indexed to the planes of anatase TiO2 (JCPDS: 21-1272). Noticeably, for the B-doped samples, the diffraction peaks become broader with a larger full width at half maximum (FWHM), which means that the crystallinity and crystal size are reduced. The average crystal sizes estimated by the Scherrer formula (from the (101) peaks) are 14.6, 9.6, 23.9, and 13.7 nm for the S-TO, B-S-TO, H-TO, and B-H-TO, respectively. The results show that B-doping can inhibit the crystal growth of TiO227-28. Moreover, the H-TO prepared by the hydrothermal method has better crystallinity and a larger crystal size than the S-TO. Rietveld refinement of the XRD data of S-TO, B-S-TO revealed a slight lattice parameter increase along the c-axis, and the unit cell volume increases from 136.288 to 136.328 Å3 after B doping. Similar trend was also observed for the hydrothermal derived samples. The swelling of the unit cell is probably caused by the interstitial boron29-30.

(a)

B 1s 191.3

192.1

B-H-TO

Ti 2p3/2

(b)

Ti 2p1/2

B-H-TO

H-TO H-TO 191.1

192.2

B-S-TO B-S-TO S-TO

186 188 190 192 194 196

Binding energy / eV

S-TO

452

456 460

464 468

Binding energy / eV

Figure 2 XPS spectra of TiO2 and boron doped TiO2 samples: (a) B 1s and (b) Ti 2p. XPS was used to investigate the elemental compositions of the resultant products. Figure 2(a) shows that the B 1s peak for both the B-S-TO and the B-H-TO samples appears at around 189–194 eV, which can be deconvoluted into two peaks cantered at 191.1 and 192.2 eV. The peak at 191.1eV could be attributed to the interstitial B in the B-Ti-O structure, while the peak at 192.2 eV can be ascribed to the B-O-B bonds in the B2O3 structure25, 28, 30-31. Part of boron ions separates from the TiO2 lattice to form diboron trioxide phase, which retards the crystal growth 27. For the Ti 2p XPS spectra (Figure 2(b)), the binding energies of Ti 2p3/2 and Ti 2p1/2 for all samples are at 458.3 eV and 464.0 eV, respectively, with a gap of 5.7eV, which shows that Ti atoms are in an octahedral skeleton and exist mainly in the form of Ti4+.


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Figure 3 TEM analysis of the TiO2 samples. TEM and HRTEM images of S-TO: (a) and (b); B-S-TO: (c) and (d); H-TO: (e) and (f); B-H-TO: (g) and (h). Insets in (a, c, e, g) are the SAED patterns, and insets in (b, d, f, h) show the lattice fringes. TEM was conducted to analyse the structure and morphology of the samples. As shown in Figure 3a and 3c, the sol-gel method yields particles with strong agglomeration. The diameter of the secondary particle size varies from several hundred nanometers to several micrometers (Figure S1 in the Supporting Information). The high resolution TEM (HRTEM) images (Figure 3b, d) show that the as-prepared S-TO and B-S-TO samples are well crystallized with visible lattice fringes. The hydrothermal method leads to well separated particles with much less agglomeration, and the individual particles of H-TO and B-H-TO (Figure 3e, g) have clear crystal facets, with particle size ranging from 20 to 40 nm and 10 to 20 nm, respectively (measured by Digital Micrograph). The BET measurement results indicate that the surface areas of S-TO, B-S-TO, H-TO and B-H-TO are 37.23, 36.11, 35.72 and 60.69 m2·g-1, respectively. Clearly, the TiO2 with B doping has a smaller crystal size, indicating that the addition of boron can inhibit the crystal growth. Well-ordered concentric diffraction rings in the selected area electron diffraction (SAED) patterns of all the TiO2 samples can be indexed to anatase (Figure 3a, c, e, and g).

250

(a)

2.5

Capacity/mAh/g

S-TO B-S-TO H-TO B-H-TO

1.5 1.0 0.5 0.0

(b)

0.1C

0.1 C

0.2C

200

2.0

Voltage/V

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0.5 C 1C

150

2C

100


50 0

0

100

200

300

400

0

20

Capacity/mA/g


40

Cycle number

60

80


50

(c)

500

(d)

400

-50

-Z" / Ohm

-1

0

Current /mA·g


-100 -150

-250


300 200 100

-200 0.0

0.5

1.0

1.5

2.0

2.5

0

3.0

0

100

Voltage/ V

200

300

400

500

Z' / Ohm

250 TO B-TO N-TO B-N-TO

(e) 200

Capacity/mAh/g

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150 100 50 0 0

60

120

180

240

300

360

Cycle number

Figure 4 Electrochemical performance of the TiO2 samples. (a) First cycle discharge/charge profiles at a current density of 0.1 C (1 C = 330 mA/g) between 0.01-2.5 V, (b) cycling and rate capability, (c) first cycle CV curves, (d) EIS spectra, (e) long-term cycling performance at the rate of 2 C (with the TiO2 first activated at the low current of 0.1 C for 5 cycles). The sodiation/desodiation performance of these TiO2 samples was investigated in the Na/TiO2 batteries. The 1st and 2nd cycle discharge/charge profiles of synthesized TiO2 are shown in Figure 4a and Figure S2, respectively. It can be seen that the first discharge curves exhibit three different plateaus: a steep plateau above 1.3 V, a sloping plateau between 1.3-0.3 V, and a flat one between 0.3-0.01 V. The irreversible peaks at about 1.1 V on the cathodic sweep of the 1st cycle cyclic voltammogram (CV) curves of all samples (Figure 4c) could be caused by side reactions such as those relating to solid electrolyte interphase (SEI) layer formation, or irreversible sites for Na-ion insertion in the crystal lattice defects, or the decomposition of organic compounds32. Profiles of B-H-TO for the first four cycles are shown in Figure S3. From the second cycle on, B-H-TO electrode exhibits a pair of peaks at 0.65 and 0.85 V that can be attributed to the reversible sodiation/desodiation. The 1st cycle reversible capacity of the S-TO, B-S-TO, H-TO, and B-H-TO is 152.3, 207.9, 159.8, and 229.6 mAh/g, respectively. Obviously, simple B doping could greatly enhance the reversible capacity of TiO2 anode. Furthermore, the two methods lead to TiO2 with different electrochemical performance, as the sol-gel method leads to agglomeration (Figure 3a and Figure S1), which is unfavourable for sodiation/desodiation. The B-doped nanosized TiO2 prepared via the hydrothermal method has the highest reversible capacity of 229.6 mAh/g for the first cycle (Table S1). This is among the highest values reported so far for TiO2 anode for sodium ion batteries 33-35. Moreover, the first cycle coulombic efficiency (CE) of pure TiO2 (S-TO) is 59.4%, which is much higher than previously reported16-18. It


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was further improved to 62.5% (B-S-TO) by boron doping (Table S1). The higher CE for the B-doped samples indicates much-improved reversible Na-ion sodiation/desodiation compared with the pure TiO2 electrode. Rate performance is a big challenge for rechargeable SIBs since the insertion and de-insertion of the large Na ions (1.02 Å) cause large volume changes in the host materials. The capacities of all the samples decrease with increasing charge/discharge rates, but they are stable and the original capacity can be recovered when the current is reversed back to 0.1 C (Figure 4b.). All the B doped samples deliver much higher reversible capacity than the non-doped ones at the rates tested, and the B-H-TO sample exhibits the highest capacity. The charge capacity of B-H-TO reaches ~225 mAh/g at the charge rate of 0.1 C and decreases to 147 mAh/g at 2 C, while the H-TO sample reaches ~164.8 mAh/g at the charge rate of 0.1 C and decreases to 116 mAh/g at 2 C. As can be seen from the CV curves (Figure 4c), for the samples prepared by the same method, the current density of the B doped sample is higher than that of the un-doped sample. Furthermore, the TiO2 derived via the hydrothermal method performs better than the samples prepared using the sol-gel method. The charge transfer resistance of the S-TO, B-S-TO, H-TO, and B-H-TO derived from the EIS spectra (Figure 4d) are around 380, 310, 280, and 210 Ω, respectively, which indicates that B doping can effectively reduce charge transfer resistance and consequently improve the rate capability. All the TiO2 shows excellent cycling stability (Figure 4e).When cycled at 2 C, the capacity of the B-H-TO sample drops slightly during the first 30 cycles from 150 mAh/g to 144 mAh/g, but the capacity decay is negligible in the following 300 cycles. As for the S-TO electrode, the capacity is around 100 mAh/g at the same rate. B-doping was found to change the resistance of TiO2 during sodiation and desodiation. The EIS spectra of H-TO and B-H-TO collected at different states of charge are shown in Figure 5. All the spectra are composed of a compressed semicircle at high frequencies and a Warburg diffusion line at low frequencies. For both the pristine and the B-doped TiO2, it is obvious that the charge transfer resistance (Rct) increases during the sodiation process and decreases during the desodiation process. Except at the fully discharged (sodiation) state, the Rct of B-H-TO is lower than that of H-TO, which means that B-doping can effectively enhance the charge transfer. At the fully discharged state, however, B-H-TO has a higher Rct. This is likely associated with the higher capacity, 263 vs. 177 mhA/g, meaning that the large amount of Na+ ions stored in the channels of B-doped TiO2 slows down the charge transfer. The Warburg diffusion part shifts to low frequencies during sodiation and moved backwards during desodiation. In addition, the EIS spectra of H-TO and B-H-TO after 2 and 20 cycles (at full charge) were shown in Figure S4. It illustrates that charge transfer resistances of both two samples are increasing upon cycling, while that of B-H-TO is still lower than H-TO after 20 cycles.

Capacity/mAh·g after activating 1 cycle discharge to 0.6 V discharge to 0.01 V charge to 1.1 V charge to 2.5 V

600

(a)

-1

800

Rct /Ohms

800

400

(b)

600 400 200 0

Voltage/V

-Z''/Ohms

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


200

2.4 1.6 0.8 0.0

0 0

200

400

600

Z'/Ohms

800

-200

-100

0

100

Capacity/mAh·g


-1

200


-1 Capacity/mAh·g after activating 1 cycle discharge to 0.6 V discharge to 0.01 V charge to 1.1 V charge to 2.5 V

800

(c)

600

(d)

600 400 200 0

400

Voltage/V

-Z''/Ohms

800

Rct / Omhs

1000

200

2.4 1.6 0.8 0.0

0 0

200

400

600

800

1000

-200

-100

0

100

200

-1 Capacity/mAh·g

Z'/Ohms

Figure 5 (a) EIS spectra of H-TO at different states of charge during the second sodiation/desodiation; (b) Change in charge transfer resistance (Rct) at the different states in (a); (c) EIS spectra of B-H-TO at different states of charge during the second sodiation/desodiation; (d) Change in charge transfer resistance at the different states in (c).

(a)

100 -1

Capacity/mAg

-200

7

5

•6

0•5

-100

8 After 10 cycles

•7 De so dia ti o n

200

•4 •3

6

4 3 So d

•2

2

on iati

-300

•1 -400 0

1

2

1

24

(101)

26

Voltage/ V

Capacity/mAg

-1

7•

100 0•6 •5

-100

De so dia ti o n

(b) 8 •

200

-200 •3

•2

36 38 40 47 48 49

(211)

54 56

9 After 10 cycles 8 7 6 5

3 2

n

-400 0

(200) (105)

2-Theta /°

atio

-300

(004)

4

•4 i Sod

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1

(004)

(200) (105)

(211)

•1 1

2 24

Voltage/ V

26

36 38 40 47 48 49

54 56

2-Theta / °

Figure 6 Ex-situ XRD patterns of H-TO (a) and B-H-TO (b) at different states of charge (SOCs) during the first cycle. To elucidate the effects of B-doping on the structural changes during sodiation and desodiation, ex-situ XRD patterns of H-TO and B-H-TO were collected at different states of charge during the first cycle (Figure 6). A closer look at the evolution of the XRD patterns during the sodiation process (from XRD


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patterns 1 to 5 in Figure 6a) reveals a slight shift of the (200) and (101) reflections to lower 2θ angles, while the (004) and (105) peaks shift to higher angles, accompanied by a slight decrease in intensity. The decrease in intensity indicates a decrease in crystallinity. The shift to lower angles of the two reflections indicates a distortion of the anatase lattice, i.e., lattice expansion along the (200) and (101) directions, and the shift to higher angles of the two reflections indicates lattice shrinkage along the (004) and (105) directions, which is presumably caused by the insertion of sodium ions. The intensity of the (200) and (101) reflections partially recovers after desodiation (from XRD patterns 5 to 7 in Figure 6a), which is similar to the results reported by Usui et al.36. They found that the diffraction peaks of rutile TiO2 shifted toward lower angles and the diffraction intensities weakened significantly during the first sodiation process, while both the peak positions and the intensities could be recovered after desodiation. After 10 cycles, almost all the anatase reflections completely vanished, which indicates that the TiO2 becomes amorphous after cycling. However, the morphology of H-TO and B-H-TO remains nearly the same after cycling (Figure S5). The disappearance of the anatase XRD reflection peaks in the discharged state is in contrast to the results reported by Kim et al.17. There is no obvious difference in these ex-situ XRD patterns between H-TO and B-H-TO. After amorphization, both H-TO and B-H-TO still maintain their capacity over 300 cycles (Figure 4), indicating that the B-doping-induced active sites are still effective in accommodating more Na+ ions. Due to the low level of doping, it is hard to probe the distribution of B at the atomic level. More studies are planned where a high level of B-doping will be used with the aim of discovering the microstructural changes with cycling and the correlation between B atomic positions and the electrochemical performance.

CONCLUSION Pristine and boron-doped anatase TiO2 particles were synthesized via a facile sol-gel method and the hydrothermal method for use as anode in SIBs. The sol-gel method leads to agglomerated TiO2, whereas the hydrothermal method is conducive to the formation of highly crystalline and discrete nanoparticles. The electrochemical measurements demonstrated that the reversible capacity of the S-TO, B-S-TO, H-TO, and B-H-TO samples is 152.3, 207.9, 159.8, and 229.6 mAh/g, respectively. The agglomeration generated by the sol-gel method is not beneficial to the sodiation, since large areas of TiO2 are not exposed to the electrolyte and the Na+ ion diffusion path is long. B-doped TiO2 delivers higher capacity, and this could be ascribed to the smaller crystal size and enlarged lattice volume. After amorphization, B-doping-induced active sites are still effective in accommodating more Na+ ions compared with pristine TiO2. The boron-doped samples exhibit excellent rate performance. The charge capacity of B-H-TO reaches 229 mAh/g at the current rate of 0.1 C and 141 mAh/g at 2 C over 400 cycles, which is likely due to the reduced resistance during sodiation and desodiation. These results indicate that the boron-doped TiO2 materials would be attractive candidate anode materials for sodium ion batteries.

AUTHOR INFORMATION * Corresponding author. Fax: +86 2135303544. E-mail address: [email protected] (B. Wang). [email protected] (Z. Huang)

SUPPORTING INFORMATION: SEM images, 2nd cycle charge/discharge profiles, CV curves of B-H-TO, EIS spectra of H-TO and B-H-TO at different cycles, TEM images of H-TO and B-H-TO after cycling, Summary table of the



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electrochemical performance.

ACKNOWLEDGEMENTS This work was supported by the Science and Technology Commission of Shanghai Municipality. (No.: 14DZ2261000). The authors would like to thank Dr. Tania Silver for polishing the manuscript, and ISEM and the Electron Microscopy Centre (EMC) for their support.

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