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C with significantly enhanced Na storage capacity for Na-ion batteries
Accepted Manuscript Title: N-doped rutile TiO2 /C with significantly enhanced Na storage capacity for Na-ion batteries Authors: Hanna He, Haiyan Wang, Dan Sun, Minhua Shao, Xiaobing Huang, Yougen Tang PII: DOI: Reference:
S0013-4686(17)30571-6 http://dx.doi.org/doi:10.1016/j.electacta.2017.03.104 EA 29138
To appear in:
Electrochimica Acta
Received date: Revised date: Accepted date:
10-12-2016 10-3-2017 14-3-2017
Please cite this article as: Hanna He, Haiyan Wang, Dan Sun, Minhua Shao, Xiaobing Huang, Yougen Tang, N-doped rutile TiO2/C with significantly enhanced Na storage capacity for Na-ion batteries, Electrochimica Actahttp://dx.doi.org/10.1016/j.electacta.2017.03.104 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
N-doped rutile TiO2/C with significantly enhanced Na storage capacity for Na-ion batteries
Hanna Hea, Haiyan Wanga,b*, Dan Suna, Minhua Shaob*, Xiaobing Huangc, and Yougen Tanga*
a
College of Chemistry and Chemical Engineering, Central South University, Changsha, 410083, P.R China.
b
Department of Chemical and Biomolecular Engineering, The Hong Kong University of Science and Technology,
Clear Water Bay, Kowloon, Hong Kong c
College of Chemistry and Chemical Engineering, Hunan University of Arts and Science, Changde, 415000, P.R
China
Graphical abstract
Abstract: Rutile TiO2 is seldom studied as anode material for Na-ion battery for its much poorer Na storage performance than anatase phase. Herein, strategies of carbon coating and nitrogen doping are proposed and achieved together by a facile ball milling method followed a high temperature sintering process to improve the electrochemical properties. The feature of this work lies in the dual N doping, not only to the carbon layers but also to the TiO2 lattice, resulting in sufficient oxygen vaccancies and defects in TiO2/C. Although the pristine TiO2 prepared by the similar method shows only 10 mAh g-1 capacity, the Na storage performance of N-doped TiO2/C is significantly enhanced. It demonstrates a high reversible discharge capacity of 211.2 mAh g−1 at 16.8 mA g−1 (0.1 C). Moreover, a capacity retention of 92.3% was achieved after 500 cycles at 168 mA g−1 (1 C), verifying ultrahigh reversible capacity and excellent cycling performance for rutile TiO2. The excellent sodium storage performance of N-doped TiO2/C should be ascribed to the improved electronic and ionic conductivity resulting from dual N-doping strategy and shortened sodium diffusion length due to the particle downsizing. Keywords: sodium ion battery · rutile TiO2 · N doping strategy · cycling performance
1. Introduction Li-ion batteries have attracted enormous attention in large-scale electric storage applications, such as stationary power plants.[1, 2] However, the limitation and unevenly distribution of lithium recourses in the world make the cost a major issue for their future large-scale applications. Na-ion batteries have been pursued as the most attractive alternatives due to their low cost, safety, sustainability and similar intercalation chemistry to Li-ion batteries.[3] For the past five years, enormous efforts have been made to develop advanced materials and the matched electrolytes for Na-ion batteries.[4-8] Graphite, the most popular anode material in Li-ion batteries, could hardly intercalate Na ions owing to the larger Na ion radius and weak substrate binding between Na and graphite.[9] For this reason, developing anode materials beyond graphite has been a big challenge for fabricating Na-ion batteries with comparable performance to that of Li-ion batteries.[10] Nanostructured TiO2 with various polymorphs (e.g., amorphous TiO2,[11, 12] TiO2B,[13] and anatase TiO2[12, 14]) has been studied as potential anode for Na-ion batteries due to their low cost, nontoxicity, safe sodiation potential, and high Na storage capacity.[15, 16] Among these polymorphs, anatase TiO2 is the most studied one due to the relatively higher Na storage capacity.[17] In contrast to the anatase phase, previous studies revealed that the rutile phase possessed higher thermal stability but much lower electrochemical performance owing to its lower electronic conductivity resulting from the higher energy barriers and restricted kinetics because of the highly anisotropic diffusion (10-6 cm2 s-1 along the c direction while only 10-15 cm2 s-1 in abplane).[18] Consequently, only a few studies on the performance of rutile TiO 2 for Na-ion batteries were reported and new strategy needs to be developed to further promote its sodium storage performance.[19-21] According to previous works, two
general strategies have been applied to ameliorate the Li storage capacity for TiO2. The first one is to downsize the particle and tune the morphology to increase the contact of the active material and electrolyte resulting in a shortened ion diffusion length.[22] The other way is to improve its electronic conductivity by adding carbon conductive additives.[23, 24] However, these strategies can only enhance the electron transport on the particle surface or between contiguous particles. It is more important to improve the electronic conductivity of bulk rutile TiO2. Recently, it has been proved that doping TiO2 with heteroatoms such as N,[25, 26] B,[27] Sn,[28] Zn,[29] and Nb,[20] could introduce Ti3+ species and oxygen vacancies thus can greatly improve its intrinsic electronic conductivity. Besides, the diffusion of Na ions can be ameliorated by slight modification of the TiO2 lattice with the dopant atoms.[30] Ti3+ self-doped rutile TiO2 nanorods exhibited remarkable rate performance and excellent cyclability in Li-ion batteries.[31] Highly improved sodium storage performance of rutile TiO2 was achieved by Nb-doping strategy. Note that the capacity of Nb-doped and pristine rutile TiO2 were 160 mAh g-1 and 25 mAh g-1, respectively.[20] Moreover, improved sodium storage performance of other titanium based materials were also achieved by doping strategy. For example, Yun et al. reported improved electrochemical performance of lithium titanate by boron-doping.[32] Among various doping atoms, N-doping strategy has made great achievement in modifying the electrochemical performance of TiO2 materials with all kind of polymorphs. For instance, N-doped TiO2 hollow nanofiber with almost twice higher capacity than that of TiO2 nanoparticles was reported.[33] Li et al. prepared N-doped anatase TiO2, which exhibited enhanced lithium storage performance than the pristine one.[34] Zhang and co-workers developed N-doped TiO2-B nanowires with excellent electrochemical performance.[30] It is believed that dual nitrogen doping strategy not
only to the TiO2 nanoparticles but also to the carbon layers would play a more effective role in promoting the Na ion storage performance of TiO2 phase. However, the relevant research on rutile TiO2 is absent currently. In this work, a simple mechanical ball milling method was employed to prepare Ndoped rutile TiO2/C. Notably, the nitrogen was successfully doped into the bulk TiO2 lattice resulting in the introducing of oxygen vacancies and the formation of a thin layer of TiO2-xNy compound on the surface, which should be conducive to improving its intrinsic conductivity thus enhance the electrochemical performance. At the same time, N-doped carbon was well formed simultaneously. When discharged at 16.8 mA g−1, the N-doped TiO2/C exhibited a high capacity of 211.2 mAh g−1. Moreover, excellent capacity retention during charge/discharge processes was demonstrated.
2. Experimental Section 2.1 Synthesis of N-doped TiO2/C. All raw materials were analytical grade and used without further purification. A simple mechanical ball milling process was designed for synthesizing N-doped TiO2/C. First, 3.4 g of tetrabutyl titanate and 0.08 g of phenolic resin were dissolved into 60 mL ethyl alcohol under vigorous stirring. 0.8 g of tripolycyanamide was then added into the mixed solution with vigorous stirring for 0.5 h. The mixture was transferred to a steel jar and milled for 4 h by a planetary ball mill (QM-3SP2). After that, the mixture was heated at 50°C with continuous stirring to remove the solvent. The obtained yellow powder was heated to 800°C for 8h in an Ar/H2 atmosphere and finally the N-doped TiO2/C composite was obtained. For comparison, pristine TiO2, TiO2/C and N-doped carbon were prepared via the similar method. The pristine TiO2 was prepared without adding phenolic resin and tripolycyanamide, and the N-doped carbon and TiO2/C were prepared without adding tetrabutyl titanate and tripolycyanamide, respectively.
2.2 Characterizations. Powder X-ray diffraction (XRD) pattern were characterized by the Dandong Haoyuan DX-2700 diffractometer using a Cu-Kα1 source. Fourier transform infrared (FT-IR) spectra were examined by the Nicolet 6700 FT-IR spectrometer. The specific surface area was obtained by using the Brunauer-Emmett-Teller (BET) equation (SSA-4200) from the N2 adsorption isotherm. Thermal gravimetry (TG) curves were conducted by a NETZSCH STA 449C differential scanning calorimeter at a ramping rate of 10 °C min-1 in air. Raman spectra were measured with LabRAM Aramis (HORIBA Jobin Yvon) spectrometer. X-ray photoelectron spectroscopy (XPS) measurement was investigated by the K-Alpha1063 spectrometer. Scanning electron microscopy (SEM) was performed on a Nova NanoSEM230. Transmission electron microscopy (TEM) and high resolution TEM (HRTEM) images were conducted by a JEOL JEM-2100F transmission electron microscope. Electron spin resonance (ESR) measurements were conducted by JES-FA200. The DC electrical conductivities of disc-shaped samples were measured by direct four-point probe conductivity measurement (RTS-9, Guangzhou, China). 2.3 Electrochemical Measurements. A mixture composed of the active material, polyvinylidene fluoride (PVDF), and Super P with a weight ratio of 80:10:10 were added into N-methyl-2-pyrrolidone solvent to obtain slurry. The slurry was then cast onto Cu foil using the doctor-blade technique and dried at 110 °C for 12 h in a vacuum oven. The CR2016 coin-type cells were assembled with a sodium metal disk as counter electrode. 1 mol L-1 NaClO4 dissolved in a 95:5 volumetric mixture of propylene carbonate and fluoroethylene carbonate was used as the electrolyte. The loading mass of each electrode was 1.5 to 2.0 mg cm-2. All cells were assembled in an Ar-filled glove box (Ipuer, MIKROUNA). Galvanostatic
charge-discharge test was performed on a Neware battery testing system (CT-3008W) at room temperature. Cyclic voltammetry (CV) was conducted on electrochemical station (Shanghai Chenhua, China) with a scan rate of 0.2 mV s−1. Electrochemical impedance spectroscopy (EIS) was recorded by a Princeton workstation (PARSTAT2273, EG&G, US) in the frequency range from 100 kHz to 10 mHz after charging to 0.8 V and then held for 2 h.
3. Results and discussion The crystal structures of all samples were analyzed by XRD measurements. As displayed in Fig.1a, the main diffraction peaks of pristine TiO2, TiO2/C and N-doped TiO2/C could be well indexed to the rutile TiO2 phase, which has the tetragonal space group with P42/mnm (136) (JCPDS No.21-1276). Several weak peaks marked with asterisks are observed for the pristine TiO2, which are assigned to anatase TiO2. It is worth noting that the anatase phase in TiO2/C is greatly reduced, and completely disappeared in N-doped TiO2/C. As shown in the magnified view of the (110) diffraction peaks of N-doped TiO2/C, TiO2/C and pristine TiO2 samples (Fig. S1), the (110) diffraction peak of N-doped TiO2/C shifts to a lower angle, suggesting the lattice expansion (Table S1). Previous reports suggested that the lattice expansion might be caused by the increase of defects such as oxygen vacancies.[35, 36] In our case, the N doping may introduce oxygen defects, thus resulting in expanded lattice. This hypothesis is also supported by XPS, Raman and ESR results, which are discussed later. In the FT-IR profile of pristine TiO2, the bands at 400–700 cm-1 are attributed to Ti–O stretching vibrations and Ti–O–Ti bridging stretching vibrations. The peaks at 1627 cm−1 and 3427 cm−1 are ascribed to O–H bending vibrations and O–H stretching vibrations, respectively (Fig. 1b).[37] These peaks are also observed in TiO2/C and Ndoped TiO2/C. In addition, a broad peak at 1105 cm−1 for C-OH stretching and two
weak peaks at 2920 cm−1 and 2850 cm−1 related to C-H vibration are observed in TiO2/C and N-doped TiO2/C, indicating the formation of carbon layers.[37, 38] It is worth noting that a small peak at 1460 cm−1 appears in N-doped TiO2/C, which should be related to the group of N–Ti–O, suggesting that N atoms have been successfully doped into the TiO2 lattice although the XRD results did not show visible diffraction peaks.[38] To detect the defects and valence states in metal oxides, EPR measurement of all samples were conducted and the results were presented in Fig. 1c. As seen, the signal at g=2.006 is observed in all samples, which is attributed to the single-electron trapped in oxygen vacancies. Obviously, the signal intensity of N-doped TiO2/C sample is much stronger, demonstrating the existence of more oxygen vacancies.[39] Moreover, an additional signal at g=1.962 appears in N-doped TiO2/C, which may originate from NTi-O species.[14] Raman spectra were employed to further investigate the phase, crystallinity and defects of all samples. As exhibited in Fig. 1d, all the peaks below 1000 cm−1 correspond to the typical Raman modes of rutile TiO2. Interestingly, the Eg mode at 431.7 cm−1 for N-doped TiO2/C shows an obvious red shift compared to pristine TiO2 and TiO2/C (Fig S2). The red-shift of Eg mode should be attributed to the introducing of oxygen vacancy defects in rutile TiO2, which agrees with the EPR results.[35] At the same time, both N-doped TiO2/C and TiO2/C exhibit typical D and G bands of carbon. The D band at 1351.4 cm−1 is ascribed to defects and the disordered portion of carbon, while the G band at 1594.58 cm−1 is indicative of the graphitic phase.[40] The ID/IG value for N-doped TiO2/C and TiO2/C are 0.94 and 0.86, respectively. Larger ID/IG value of the former suggests more defects and a lower degree of graphitization of carbon[41, 42], which is more beneficial to the insertion of sodium ion.[10] All the
above characterization results demonstrate that the nitrogen is doped into the TiO2 lattice, generating more oxygen vacancies in TiO2, which would increase the electrical conductivity and narrow the bandgap energy, consequently affording excellent kinetic properties. [43, 44] The samples were further evaluated by XPS measurement. As seen in Fig. 2a, all the three samples show the similar peaks of Ti and O elements except for an obvious peak of N 1s in N-doped TiO2/C, further supporting the successful doping of nitrogen atoms. The weight content of nitrogen is calculated to be about 4.39 % based on the XPS results. The high resolution Ti 2p region of N-doped TiO2/C can be resolved into three peaks (Fig. 2b). The two peaks at 459.1 and 464.8 eV are correspond to the Ti4+ 2p3/2 and Ti4+ 2p1/2, respectively, while the weak peak at 459.5 eV belongs to the N-TiO linkage.[43] The substitution of oxygen by nitrogen on the oxygen sites in TiO2 could increase the oxygen vacancy concentration to maintain the electroneutrality.[25, 43] The N 1S spectrum (Fig. 2c) can be resolved into five functional groups. The peaks located at 402.8, 401.3, 399.1 and 398.3 eV are ascribed to oxidized pyridinic, graphitic, pyrrolic and pyridinic N in carbon layer, respectively.[45] As we know, the nitrogen source is melamine, which would transform into nitrogen doped carbon during the carbonization process in an inert atmosphere, thus the majority nitrogen is doped into the carbon layers. Interestingly, a weak peak, which is ascribed to the N-Ti-O linkage, is observed at 399.5 eV, further demonstrating the successful doping of nitrogen into TiO2 lattice.[46, 47] For the high resolution C 1s region (Fig. 2d), the most intensive peak at 284.6 eV is assigned to C=C, and the other three peaks are assigned to O-C=O (289.2 eV), C=N or C=O (286.4 eV), and C-O (285.3 eV). The Ti 2p and O 1s spectra of all samples were compared in Fig. 2e and Fig. 2f, respectively. Obvious blue shift of Ti 2p and O 1s peaks are exhibited in TiO2/C and N-doped TiO2/C samples, which is
also found in some N-doped[48, 49] and heteroatoms doped TiO2[35, 50], and may be originated from the introducing of oxygen vacancies according to previous reports. As we know, the divided peaks at about 531.3 eV in Fig. 2f are ascribed to the chemisorbed oxygen species or the defect oxygen.[51] For comparison, the proportion of this peak in N-doped TiO2/C is much higher than other two samples, revealing more oxygendefective sites in the former. The results here demonstrate that N-doping into TiO2/C introduced more oxygen vacancies, which would increase the electrical conductivity of rutile TiO2.[43, 44] Carbon contents in N-doped TiO2/C and TiO2/C samples are estimated to be 6.5 wt% and 8.4 wt% from TG results (Fig. S3). The N2 adsorption–desorption isotherms of Ndoped TiO2/C in Fig. S4 shows a typical Langmuir type IV characteristic of a distinct hysteresis loop, demonstrating a mesoporous structure of carbon. The surface areas (Fig. S4. a, c, e) of pristine TiO2, TiO2/C and N-doped TiO2/C are 7.70, 77.93 and 111.05 m2·g−1, respectively. Much higher surface areas of TiO2/C and N-doped TiO2/C should be mainly caused by the carbon coating. Moreover, the decomposition of melamine would release a large amount of gas during the sintering process, leading to the formation of nanopores in carbon layer. (Fig.S4. b, d, f).[37] The TEM image of N-doped TiO2/C in Fig. 3a shows that TiO2 nanoparticles have a uniform size distribution ranging from 5 nm to 50 nm. The SAED analysis (inset in Fig. 3a) confirms the high crystallinity of TiO2 nanoparticles. The HRTEM images (Fig. 3b and 3c) of the individual nanoparticle display lattice spacing of 0.22 nm, which could be indexed to the (111) plane of the rutile TiO2. It is interesting to note that another lattice spacing of 0.10 nm is also observed in the edge of a TiO2 nanoparticle. For comparison, HRTEM images of pristine TiO2 and TiO2/C are shown in Fig. 4. Notably, there are no such lattice fringes in pristine TiO2 and TiO2/C except for those with lattice
spacing of 0.20 nm. According to the XRD patterns, the lattice space of all the planes of the rutile TiO2 is larger than 0.1 nm, thus the 0.1 nm lattice spacing at the edge of the TiO2 nanoparticles cannot be originated from the TiO2 nanoparticles. We once considered that the 0.1nm lattice spacing should be caused by the TiN compound because it could be well indexed to the (400) plane of TiN. However, in XPS spectrum, we could not find the related information of low valance of Ti. Accordingly, we suspect that the lattice spacing of 0.10 nm results from a new phase causing by N-doping. For convenience, this uncertain phase is denoted as TiO2-xNy. Fig. 3d further demonstrates that the TiO2 nanoparticles are embedded in the N-doped carbon and the composite consists of TiO2, TiO2-xNy, and N-doped carbon layers. The TiO2-xNy and doped carbon layers could improve the electronic conductivity and suppress the aggregation of TiO2 nanoparticles during sodium insertion/extraction. The elemental mapping analysis of N-doped TiO2/C (Fig. S4) further confirms the homogeneous distribution of Ti, O, C and N elements among the composites. It should be mentioned that this is the first report involving the dual N-doping not only for carbon layers but also for rutile TiO2, which can greatly improve the intrinsic conductivity and electrochemical performance of rutile TiO2.[33, 52, 53] The electronic conductivities of N-doped TiO2/C, TiO2/C and pristine TiO2 are 4.87×10−4 S cm−1, 1.32×10−4 S cm−1 and 4.22×10−8 S cm−1, respectively, measured by a four-point probe method. Our results are consistent with references showing that ammonia treated TiO2 exhibited much higher Li storage capacity owe to the generation of a thin titanium nitride layer.[25, 38, 46] The particle size of N-doped TiO2/C is much smaller than those of pristine TiO2 (range from 300 to 700 nm) and TiO2/C (range from 100 to 200 nm) as demonstrated in Fig. 4. Smaller particle size could also contribute to a higher BET area (Fig. S5).
The sodium storage of all as-prepared samples were investigated. It has been reported that rutile could only reversibly uptake 0.5 Na per formula unit, thus we calculated its practical theoretical capacity (the detailed calculation was present in the supporting information) and set 1C=168 mAh g-1 in our work.[54] Fig. 5a demonstrates the galvanostatic charge-discharge curves of all samples at a current density of 84 mA g−1 after 100 cycles. It is found that the pristine TiO2 exhibits ultralow discharge capacity (~10 mAh g−1). After carbon coating, the reversible capacity increases to 58.8 mAh g−1 for TiO2/C. N-doped TiO2/C shows a highest capacity of 166.4 mAh g−1. In fact, poor lithium insertion performance was also reported in bulk rutile TiO2 due to its low intrinsic electronic conductivity and poor ion diffusion in the ab planes (Dab=10−15cm2/s), which may prevent Li ions from reaching the thermodynamically favorable octahedral sites and separates Li in the c channels [18]. Here, the significantly enhanced discharge capacity of N-doped TiO2/C may mainly originate from its improved electronic conductivity and larger surface area. In order to understand the contribution of N-doped carbon to the capacity, its charge-discharge curve was also shown in Fig. 5a. A discharge capacity of 94.8 mAh g−1 is obtained after 100 cycles, much less than that of N-doped TiO2/C. Given that the carbon content in N-doped TiO2/C is only 6.5 wt%, it is clear sure that the dominant contribution to the high capacity in N-doped TiO2/C is not from N-doped carbon layer. Fig. 5b shows the cyclability of all samples at 84 mA g−1. Obviously, N-doped TiO2/C demonstrates a great enhancement in discharge capacity compared to TiO2/C and pristine TiO2. It exhibits a high discharge capacity of 154.6 mAh g−1 after 200 cycles, while those for TiO2/C, pristine TiO2 and N-doped carbon are 65.4, 10 and 92.9 mAh g−1, respectively. As we know, the reproducibility of the battery performance is a critical factor for electrode materials, thus we tested the sodium storage performance of N-doped TiO2/C
with different bathes and its cycling performance at 84 mA g−1 was presented in Fig. S6. Interestingly, the cycling performance of N-doped TiO2/C at 84 mA g-1 of different batches show almost the same tendency and discharge capacity, indicating the excellent reproducibility of the electrode materials. The Li storage performance of all samples were also evaluated and presented in Fig. S7. Clearly, N-doped TiO2/C also shows the significantly improved Li storage properties. It delivers a high capacity of 403.3 mAh g−1 (Fig. S7b) at 168 mA g−1, and 138.6 mAh g−1 at 3360 mA g−1 (Fig. S7c). Moreover, the Li storage performance of the N-doped TiO2/C is superior to most reported pristine and modified rutile TiO2 (Table S2).[27, 55] Notably, the discharge capacities increase in the initial several cycles in both Na-ion and Li-ion batteries is caused by electrochemical activation process, that is, the reversible insertion of sodium ions is kinetically hindered thus repeated charging−discharging process is needed, which has been reported in some previous literature involving TiO2-based anode materials for Na ion batteries. [56] The charge-discharge process of Na in N-doped TiO2/C was further evaluated by cyclic voltammetry (CV) measurements (Fig. 6a). In the first cycle, a large irreversible peak is located at about 0.8 V which should be caused by the formation of solid electrolyte interface (SEI) layers and electrolyte decomposition.[57] From the second cycle on, a pair of cathodic/anodic peaks at 0.62/0.83 V corresponding to the desodiation/sodiation process is observed, in agreement with the previous reports.[56, 58] Notably, the peak area almost maintains the same after the first cycle. The irreversible capacity loss in the initial cycle has been realized as a big issue for Na-ion battery anode.[59] A pre-sodiation process, optimizing the binder and modifying the electrolyte surface have demonstrated as an effectively approach to reduce the initial capacity loss.[57]
Charge-discharge curves of N-doped TiO2/C electrode at various current densities are displayed in Fig. 6b. With the current density increasing from 16.8 mA g−1 to 336 mA g−1, the shapes of charge-discharge profiles are preserved well and the discharge capacities of 211.2, 171.5, 161.2, 122.2, 83.1 mAh g−1 are obtained at 16.8, 33.6, 84, 168 and 336 mA g−1, respectively. At each rate, no visible capacity loss is observed when the current density turns back to 16.8 mA g−1 after 45 cycles, manifesting excellent rate capability and cyclability of N-doped TiO2/C (Fig. 6c). Notably, a remarkably large capacity of 211.2 mAh g−1 is achieved at 16.8 mA g−1, equalling 0.63 Na ions insertion into a formula of rutile TiO2, which is superior to most of previously reports.[19, 20] A discharge capacity of 160.2 mAh g−1 of rutile TiO2 at 16.8 mA g−1 was reported by Zhang et al.,[19] and 160 mAh g−1 at 50 mA g−1 for Nb-doped rutile TiO2 by Usui et al.[20] Long cycling performance at different current densities was tested to further evaluate the sodium storage properties of N-doped TiO2/C electrode. As shown in Fig. 6d, the electrode demonstrates a discharge capacity of 175.3 mAh g−1 after 100 cycles at 33.6 mA g−1 (0.2C). The galvanostatic charge-discharge profiles (Fig. 6e) after different cycles repeat well without visible capacity fading, revealing high reversibility of sodium ions during the insertion/extraction process. Remarkably, a discharge capacity of 121.2 mAh g–1 at 168 mA g−1 can still be retained after 500 cycles, with a Coulombic efficiency near 100% throughout the cycling (Fig. 6f). The superior cycling stability is further confirmed by the galvanostatic discharge profiles after different cycles at 168 mA g−1 (Fig. S8). As exhibited in Table S3, amorphous TiO2 with a discharge capacity of 150 mAh g−1 at 50 mA g−1, TiO2(B) with 102 mAh g–1 at 33.5 mA g−1 and TiO2(H) with 85 mAh g−1 at 41.8 mA g−1 were reported.[11, 60, 61] The reversible discharge capacity of N-doped rutile TiO2/C in the current study is much higher than the above mentioned TiO2 polymorphs. Generally, anatase TiO2 has
been considered as a better anode material for Na-ion battery in comparison with other TiO2 polymorphs and some anatase TiO2 with excellent sodium storage performance were reported.[62-64] It is interesting to note in Table S3 that N-doped rutile TiO2/C here is comparable to some reported anatase TiO2 or TiO2/C with typical electrochemical performance.[57, 65-67] Our results demonstrated that a much better performance could be achieved on rutile TiO2 by a simple and scalable N-doping strategy. It should be noted that the synthesis of TiO2 in this work was very simple and the pristine sample exhibited poor electrochemical properties. The key point of this work is to introduce a fascinating N-doping strategy. We believe that the electrochemical properties of rutile TiO2 can be further improved if we combine the Ndoping strategy with morphology tuning of TiO2. To gain more insight into the influence of N-doping, EIS measurements were measured. As shown in Fig. 7, the Nyquist plots consist of an inclined line in the lowfrequency region and one semicircle in the high-frequency region. The fitting equivalent circuit is also attached. The inclined line represents Warburg impedance corresponding to the sodium ion diffusion resistance (W1) in a solid electrode. The small intercept at the Z′ axis demonstrates the internal resistance of electrolyte (Rs) while the semicircle corresponds to charge-transfer impedance on the electrode– electrolyte interface (R1).[31] CPE1 and C1 represent a constant phase and the doublelayer capacitance, respectively.[68] Obviously, the R1 of N-doped TiO2/C (110.5 Ω·cm−2) is much lower than that of TiO2/C (126.5 Ω·cm−2) and pristine TiO2 (215.4 Ω·cm−2), verifying an enhanced charge transfer on the interface of electrode and electrolyte (Table S4). Increasing of R1 has been demonstrated as an important contributor to the capacity fading of active materials. We expect that an improved electrochemical performance of N-doped TiO2/C could be resulted from the lower
R1.[69] Moreover, as shown in Fig. S9, the Na-ion diffusion coefficient of pristine TiO2, TiO2/C and N-doped TiO2/C are calculated to be ca. 9.9×10−15, 1.58×10−13 and 2.68×10−13 cm2 S−1, respectively. Apparently, the appearance of the TiO2-xNy layer on the TiO2 surface and more defects in both carbon layer and intrinsic rutile TiO2 are beneficial to restrain the charge-transfer resistance and increase the sodium ion diffusion coefficient, resulting in superior Na storage properties. To further understand the insertion/extraction mechanism of Na ions in the modified rutile TiO2, XRD patterns of N-doped TiO2/C electrodes at different voltages (marked in Fig. 8a) are collected in Fig. 8. For comparison, the XRD patterns of electrodes (Fig. 8b) show some additional peaks compared to those of N-doped TiO2/C powder in Fig.1, in agreement with our previous reports.[70] In the first cycle, the main diffraction peaks of (110) and (211) diffraction peaks of rutile TiO2 shifts to higher angles during the discharge process (a-d) and return to the original value during the charge process (e - g). The shifts of the peak position indicate the change of lattice parameters, which are listed in Table S5. The reversible of peak shifting implies that the crystal structure of the TiO2 phase is maintained, i.e., reversible Na ion insertion/extraction from the crystal lattice. The comparison of XRD patterns after different cycles is displayed in Fig. S10. We can see that the intensity of the diffraction peaks reduces a little after a long-term cycling, which may be caused by the repeated de-intercalation and intercalation of Na-ions, resulting in slight crystallinity decrease of the N-doped TiO2/C electrode. However, the main diffraction peaks are well maintained, indicating that there is no obvious amorphization of the TiO2 after the cycling, demonstrating the excellent structure stability of N-doped TiO2/C electrode during the cycling.
The electrochemical performance of N-doped TiO2/C was significantly improved in comparison with pristine TiO2 and TiO2/C. Possible reasons for this improvement may include: (1) the N-doped carbon layer or carbon matrix resulted in much smaller nanoparticles and abundant porosity, leading to a shorter sodium diffusion path; (2) the well-formed carbon layer could suppress the clustering of nanoparticles and alleviate the structural strain resulting from the volume change during Na ions insertion/extraction process; (3) the dual nitrogen doping to carbon layer and TiO2 can introduce more defects and oxygen vacancies, greatly enhancing the electronic and ionic conductivities and facilitating the efficient diffusion of Na ions between the carbon and rutile TiO2. As we know, ammonia treated TiO2 exhibited much higher Li storage ability due to the appearance of thin titanium nitride layer on the TiO 2 surface, which has higher electrical and thermal conductivity than the TiO2. [25, 38, 46]
4. Conclusion In summary, N-doped rutile TiO2/C was successfully synthesized by a facile N doping method. It was found that the N-doping into carbon layers and TiO2 surface was achieved together. This sample exhibited significantly improved Na storage performance as anode. It showed a high discharge capacity of 211.2 mAh g−1 at 16.8 mA g−1, and 175.3 mAh g−1 was obtained at 33.6 mA g−1 after 100 cycles. Excellent long-term cycling life with capacity retention of 92.3% after 500 cycles at 168 mA g-1 was also observed. N-doped carbon layers and TiO2-xNy coating layers on TiO2 surface greatly contributed to the enhanced Na storage performance. This work provides a crucial clue for further improving the Na storage ability of rutile phase TiO2 and the asprepared N-doped TiO2/C here can be employed as a promising anode candidate for Na-ion batteries.
Acknowledgement This work was financially supported by the National Nature Science Foundation of China (No. 21671200 and No.21571189), the Fundamental Research Funds for the Central Universities of Central South University and the Open-End Fund for Valuable and Precision Instruments of Central South University (CSUZC201729). The work at the Hong Kong University of Science and Technology was supported by a start up fund.
References
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Figure captions Fig.1 (a) XRD patterns, (b) FT-IR spectra, (c) ESR spectra and (d) Raman spectra of pristine rutile TiO2, TiO2/C and N-doped TiO2/C samples. Fig. 2 XPS survey spectra of pristine TiO2, TiO2/C and N-doped TiO2/C samples (a). Ti 2p (b), N 1s (c), C 1s (d) spectra of N-doped TiO2/C. Ti 2 p (e) and O1s (f) spectra of pristine TiO2, TiO2/C and N-doped TiO2/C samples. Fig.3 (a) TEM image of N-doped TiO2/C. Inset shows the corresponding SAED patterns. (b, c, d) High-resolution TEM images of N-doped TiO2/C. Fig.4 (a) TEM image of pristine TiO2 (a) and TiO2/C (b). High-resolution TEM images of pristine TiO2 (c) and TiO2/C (d). Fig.5 (a) Charge-discharge curves of N-doped carbon, N-doped TiO2/C, TiO2/C and rutile TiO2 after 100 cycles at 84 mA g−1. (b) Cycling performance of N-C, N-doped TiO2/C, TiO2/C and rutile TiO2 at 84 mA g−1. Fig.6 (a) CV curves of as-prepared N-doped TiO2/C between 0.001 and 3.0 V at 0.4 mV s−1, (b) Reversible charge-discharge curves of N-doped TiO2/C at various current densities from 16.8 to 336 mA g−1, (c) Rate performance of N-doped TiO2/C, (d) Cycling performance of N-doped TiO2/C at 33.6 mA g−1, (e) Charge-discharge profiles of N-doped TiO2/C after different cycles at 33.6 mA g−1, (f) Cycling performance and Coulombic efficiency of N-doped TiO2/C at 168 mA g−1. Fig.7 Electrochemical impedance spectroscopys of pristine TiO2, TiO2/C and N-doped TiO2/C electrodes after 3 cycles at 33.6 mA g−1.
Fig. 8 (a) Charge-discharge curves of N-doped TiO2/C at 168 mA g−1. (b) XRD patterns of N-doped TiO2/C at different voltages: a= 3.0 V, b=1.8 V, c=1.2 V, d= 0.01 V, e= 0.8 V, f=1.3 V, g= 3.0 V at 168 mA g-1. Patterns in (c) and (d) are the partially enlarged portions of the highlighted parts in (b).
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S0013-4686(17)30571-6 http://dx.doi.org/doi:10.1016/j.electacta.2017.03.104 EA 29138
To appear in:
Electrochimica Acta
Received date: Revised date: Accepted date:
10-12-2016 10-3-2017 14-3-2017
Please cite this article as: Hanna He, Haiyan Wang, Dan Sun, Minhua Shao, Xiaobing Huang, Yougen Tang, N-doped rutile TiO2/C with significantly enhanced Na storage capacity for Na-ion batteries, Electrochimica Actahttp://dx.doi.org/10.1016/j.electacta.2017.03.104 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
N-doped rutile TiO2/C with significantly enhanced Na storage capacity for Na-ion batteries
Hanna Hea, Haiyan Wanga,b*, Dan Suna, Minhua Shaob*, Xiaobing Huangc, and Yougen Tanga*
a
College of Chemistry and Chemical Engineering, Central South University, Changsha, 410083, P.R China.
b
Department of Chemical and Biomolecular Engineering, The Hong Kong University of Science and Technology,
Clear Water Bay, Kowloon, Hong Kong c
College of Chemistry and Chemical Engineering, Hunan University of Arts and Science, Changde, 415000, P.R
China
Graphical abstract
Abstract: Rutile TiO2 is seldom studied as anode material for Na-ion battery for its much poorer Na storage performance than anatase phase. Herein, strategies of carbon coating and nitrogen doping are proposed and achieved together by a facile ball milling method followed a high temperature sintering process to improve the electrochemical properties. The feature of this work lies in the dual N doping, not only to the carbon layers but also to the TiO2 lattice, resulting in sufficient oxygen vaccancies and defects in TiO2/C. Although the pristine TiO2 prepared by the similar method shows only 10 mAh g-1 capacity, the Na storage performance of N-doped TiO2/C is significantly enhanced. It demonstrates a high reversible discharge capacity of 211.2 mAh g−1 at 16.8 mA g−1 (0.1 C). Moreover, a capacity retention of 92.3% was achieved after 500 cycles at 168 mA g−1 (1 C), verifying ultrahigh reversible capacity and excellent cycling performance for rutile TiO2. The excellent sodium storage performance of N-doped TiO2/C should be ascribed to the improved electronic and ionic conductivity resulting from dual N-doping strategy and shortened sodium diffusion length due to the particle downsizing. Keywords: sodium ion battery · rutile TiO2 · N doping strategy · cycling performance
1. Introduction Li-ion batteries have attracted enormous attention in large-scale electric storage applications, such as stationary power plants.[1, 2] However, the limitation and unevenly distribution of lithium recourses in the world make the cost a major issue for their future large-scale applications. Na-ion batteries have been pursued as the most attractive alternatives due to their low cost, safety, sustainability and similar intercalation chemistry to Li-ion batteries.[3] For the past five years, enormous efforts have been made to develop advanced materials and the matched electrolytes for Na-ion batteries.[4-8] Graphite, the most popular anode material in Li-ion batteries, could hardly intercalate Na ions owing to the larger Na ion radius and weak substrate binding between Na and graphite.[9] For this reason, developing anode materials beyond graphite has been a big challenge for fabricating Na-ion batteries with comparable performance to that of Li-ion batteries.[10] Nanostructured TiO2 with various polymorphs (e.g., amorphous TiO2,[11, 12] TiO2B,[13] and anatase TiO2[12, 14]) has been studied as potential anode for Na-ion batteries due to their low cost, nontoxicity, safe sodiation potential, and high Na storage capacity.[15, 16] Among these polymorphs, anatase TiO2 is the most studied one due to the relatively higher Na storage capacity.[17] In contrast to the anatase phase, previous studies revealed that the rutile phase possessed higher thermal stability but much lower electrochemical performance owing to its lower electronic conductivity resulting from the higher energy barriers and restricted kinetics because of the highly anisotropic diffusion (10-6 cm2 s-1 along the c direction while only 10-15 cm2 s-1 in abplane).[18] Consequently, only a few studies on the performance of rutile TiO 2 for Na-ion batteries were reported and new strategy needs to be developed to further promote its sodium storage performance.[19-21] According to previous works, two
general strategies have been applied to ameliorate the Li storage capacity for TiO2. The first one is to downsize the particle and tune the morphology to increase the contact of the active material and electrolyte resulting in a shortened ion diffusion length.[22] The other way is to improve its electronic conductivity by adding carbon conductive additives.[23, 24] However, these strategies can only enhance the electron transport on the particle surface or between contiguous particles. It is more important to improve the electronic conductivity of bulk rutile TiO2. Recently, it has been proved that doping TiO2 with heteroatoms such as N,[25, 26] B,[27] Sn,[28] Zn,[29] and Nb,[20] could introduce Ti3+ species and oxygen vacancies thus can greatly improve its intrinsic electronic conductivity. Besides, the diffusion of Na ions can be ameliorated by slight modification of the TiO2 lattice with the dopant atoms.[30] Ti3+ self-doped rutile TiO2 nanorods exhibited remarkable rate performance and excellent cyclability in Li-ion batteries.[31] Highly improved sodium storage performance of rutile TiO2 was achieved by Nb-doping strategy. Note that the capacity of Nb-doped and pristine rutile TiO2 were 160 mAh g-1 and 25 mAh g-1, respectively.[20] Moreover, improved sodium storage performance of other titanium based materials were also achieved by doping strategy. For example, Yun et al. reported improved electrochemical performance of lithium titanate by boron-doping.[32] Among various doping atoms, N-doping strategy has made great achievement in modifying the electrochemical performance of TiO2 materials with all kind of polymorphs. For instance, N-doped TiO2 hollow nanofiber with almost twice higher capacity than that of TiO2 nanoparticles was reported.[33] Li et al. prepared N-doped anatase TiO2, which exhibited enhanced lithium storage performance than the pristine one.[34] Zhang and co-workers developed N-doped TiO2-B nanowires with excellent electrochemical performance.[30] It is believed that dual nitrogen doping strategy not
only to the TiO2 nanoparticles but also to the carbon layers would play a more effective role in promoting the Na ion storage performance of TiO2 phase. However, the relevant research on rutile TiO2 is absent currently. In this work, a simple mechanical ball milling method was employed to prepare Ndoped rutile TiO2/C. Notably, the nitrogen was successfully doped into the bulk TiO2 lattice resulting in the introducing of oxygen vacancies and the formation of a thin layer of TiO2-xNy compound on the surface, which should be conducive to improving its intrinsic conductivity thus enhance the electrochemical performance. At the same time, N-doped carbon was well formed simultaneously. When discharged at 16.8 mA g−1, the N-doped TiO2/C exhibited a high capacity of 211.2 mAh g−1. Moreover, excellent capacity retention during charge/discharge processes was demonstrated.
2. Experimental Section 2.1 Synthesis of N-doped TiO2/C. All raw materials were analytical grade and used without further purification. A simple mechanical ball milling process was designed for synthesizing N-doped TiO2/C. First, 3.4 g of tetrabutyl titanate and 0.08 g of phenolic resin were dissolved into 60 mL ethyl alcohol under vigorous stirring. 0.8 g of tripolycyanamide was then added into the mixed solution with vigorous stirring for 0.5 h. The mixture was transferred to a steel jar and milled for 4 h by a planetary ball mill (QM-3SP2). After that, the mixture was heated at 50°C with continuous stirring to remove the solvent. The obtained yellow powder was heated to 800°C for 8h in an Ar/H2 atmosphere and finally the N-doped TiO2/C composite was obtained. For comparison, pristine TiO2, TiO2/C and N-doped carbon were prepared via the similar method. The pristine TiO2 was prepared without adding phenolic resin and tripolycyanamide, and the N-doped carbon and TiO2/C were prepared without adding tetrabutyl titanate and tripolycyanamide, respectively.
2.2 Characterizations. Powder X-ray diffraction (XRD) pattern were characterized by the Dandong Haoyuan DX-2700 diffractometer using a Cu-Kα1 source. Fourier transform infrared (FT-IR) spectra were examined by the Nicolet 6700 FT-IR spectrometer. The specific surface area was obtained by using the Brunauer-Emmett-Teller (BET) equation (SSA-4200) from the N2 adsorption isotherm. Thermal gravimetry (TG) curves were conducted by a NETZSCH STA 449C differential scanning calorimeter at a ramping rate of 10 °C min-1 in air. Raman spectra were measured with LabRAM Aramis (HORIBA Jobin Yvon) spectrometer. X-ray photoelectron spectroscopy (XPS) measurement was investigated by the K-Alpha1063 spectrometer. Scanning electron microscopy (SEM) was performed on a Nova NanoSEM230. Transmission electron microscopy (TEM) and high resolution TEM (HRTEM) images were conducted by a JEOL JEM-2100F transmission electron microscope. Electron spin resonance (ESR) measurements were conducted by JES-FA200. The DC electrical conductivities of disc-shaped samples were measured by direct four-point probe conductivity measurement (RTS-9, Guangzhou, China). 2.3 Electrochemical Measurements. A mixture composed of the active material, polyvinylidene fluoride (PVDF), and Super P with a weight ratio of 80:10:10 were added into N-methyl-2-pyrrolidone solvent to obtain slurry. The slurry was then cast onto Cu foil using the doctor-blade technique and dried at 110 °C for 12 h in a vacuum oven. The CR2016 coin-type cells were assembled with a sodium metal disk as counter electrode. 1 mol L-1 NaClO4 dissolved in a 95:5 volumetric mixture of propylene carbonate and fluoroethylene carbonate was used as the electrolyte. The loading mass of each electrode was 1.5 to 2.0 mg cm-2. All cells were assembled in an Ar-filled glove box (Ipuer, MIKROUNA). Galvanostatic
charge-discharge test was performed on a Neware battery testing system (CT-3008W) at room temperature. Cyclic voltammetry (CV) was conducted on electrochemical station (Shanghai Chenhua, China) with a scan rate of 0.2 mV s−1. Electrochemical impedance spectroscopy (EIS) was recorded by a Princeton workstation (PARSTAT2273, EG&G, US) in the frequency range from 100 kHz to 10 mHz after charging to 0.8 V and then held for 2 h.
3. Results and discussion The crystal structures of all samples were analyzed by XRD measurements. As displayed in Fig.1a, the main diffraction peaks of pristine TiO2, TiO2/C and N-doped TiO2/C could be well indexed to the rutile TiO2 phase, which has the tetragonal space group with P42/mnm (136) (JCPDS No.21-1276). Several weak peaks marked with asterisks are observed for the pristine TiO2, which are assigned to anatase TiO2. It is worth noting that the anatase phase in TiO2/C is greatly reduced, and completely disappeared in N-doped TiO2/C. As shown in the magnified view of the (110) diffraction peaks of N-doped TiO2/C, TiO2/C and pristine TiO2 samples (Fig. S1), the (110) diffraction peak of N-doped TiO2/C shifts to a lower angle, suggesting the lattice expansion (Table S1). Previous reports suggested that the lattice expansion might be caused by the increase of defects such as oxygen vacancies.[35, 36] In our case, the N doping may introduce oxygen defects, thus resulting in expanded lattice. This hypothesis is also supported by XPS, Raman and ESR results, which are discussed later. In the FT-IR profile of pristine TiO2, the bands at 400–700 cm-1 are attributed to Ti–O stretching vibrations and Ti–O–Ti bridging stretching vibrations. The peaks at 1627 cm−1 and 3427 cm−1 are ascribed to O–H bending vibrations and O–H stretching vibrations, respectively (Fig. 1b).[37] These peaks are also observed in TiO2/C and Ndoped TiO2/C. In addition, a broad peak at 1105 cm−1 for C-OH stretching and two
weak peaks at 2920 cm−1 and 2850 cm−1 related to C-H vibration are observed in TiO2/C and N-doped TiO2/C, indicating the formation of carbon layers.[37, 38] It is worth noting that a small peak at 1460 cm−1 appears in N-doped TiO2/C, which should be related to the group of N–Ti–O, suggesting that N atoms have been successfully doped into the TiO2 lattice although the XRD results did not show visible diffraction peaks.[38] To detect the defects and valence states in metal oxides, EPR measurement of all samples were conducted and the results were presented in Fig. 1c. As seen, the signal at g=2.006 is observed in all samples, which is attributed to the single-electron trapped in oxygen vacancies. Obviously, the signal intensity of N-doped TiO2/C sample is much stronger, demonstrating the existence of more oxygen vacancies.[39] Moreover, an additional signal at g=1.962 appears in N-doped TiO2/C, which may originate from NTi-O species.[14] Raman spectra were employed to further investigate the phase, crystallinity and defects of all samples. As exhibited in Fig. 1d, all the peaks below 1000 cm−1 correspond to the typical Raman modes of rutile TiO2. Interestingly, the Eg mode at 431.7 cm−1 for N-doped TiO2/C shows an obvious red shift compared to pristine TiO2 and TiO2/C (Fig S2). The red-shift of Eg mode should be attributed to the introducing of oxygen vacancy defects in rutile TiO2, which agrees with the EPR results.[35] At the same time, both N-doped TiO2/C and TiO2/C exhibit typical D and G bands of carbon. The D band at 1351.4 cm−1 is ascribed to defects and the disordered portion of carbon, while the G band at 1594.58 cm−1 is indicative of the graphitic phase.[40] The ID/IG value for N-doped TiO2/C and TiO2/C are 0.94 and 0.86, respectively. Larger ID/IG value of the former suggests more defects and a lower degree of graphitization of carbon[41, 42], which is more beneficial to the insertion of sodium ion.[10] All the
above characterization results demonstrate that the nitrogen is doped into the TiO2 lattice, generating more oxygen vacancies in TiO2, which would increase the electrical conductivity and narrow the bandgap energy, consequently affording excellent kinetic properties. [43, 44] The samples were further evaluated by XPS measurement. As seen in Fig. 2a, all the three samples show the similar peaks of Ti and O elements except for an obvious peak of N 1s in N-doped TiO2/C, further supporting the successful doping of nitrogen atoms. The weight content of nitrogen is calculated to be about 4.39 % based on the XPS results. The high resolution Ti 2p region of N-doped TiO2/C can be resolved into three peaks (Fig. 2b). The two peaks at 459.1 and 464.8 eV are correspond to the Ti4+ 2p3/2 and Ti4+ 2p1/2, respectively, while the weak peak at 459.5 eV belongs to the N-TiO linkage.[43] The substitution of oxygen by nitrogen on the oxygen sites in TiO2 could increase the oxygen vacancy concentration to maintain the electroneutrality.[25, 43] The N 1S spectrum (Fig. 2c) can be resolved into five functional groups. The peaks located at 402.8, 401.3, 399.1 and 398.3 eV are ascribed to oxidized pyridinic, graphitic, pyrrolic and pyridinic N in carbon layer, respectively.[45] As we know, the nitrogen source is melamine, which would transform into nitrogen doped carbon during the carbonization process in an inert atmosphere, thus the majority nitrogen is doped into the carbon layers. Interestingly, a weak peak, which is ascribed to the N-Ti-O linkage, is observed at 399.5 eV, further demonstrating the successful doping of nitrogen into TiO2 lattice.[46, 47] For the high resolution C 1s region (Fig. 2d), the most intensive peak at 284.6 eV is assigned to C=C, and the other three peaks are assigned to O-C=O (289.2 eV), C=N or C=O (286.4 eV), and C-O (285.3 eV). The Ti 2p and O 1s spectra of all samples were compared in Fig. 2e and Fig. 2f, respectively. Obvious blue shift of Ti 2p and O 1s peaks are exhibited in TiO2/C and N-doped TiO2/C samples, which is
also found in some N-doped[48, 49] and heteroatoms doped TiO2[35, 50], and may be originated from the introducing of oxygen vacancies according to previous reports. As we know, the divided peaks at about 531.3 eV in Fig. 2f are ascribed to the chemisorbed oxygen species or the defect oxygen.[51] For comparison, the proportion of this peak in N-doped TiO2/C is much higher than other two samples, revealing more oxygendefective sites in the former. The results here demonstrate that N-doping into TiO2/C introduced more oxygen vacancies, which would increase the electrical conductivity of rutile TiO2.[43, 44] Carbon contents in N-doped TiO2/C and TiO2/C samples are estimated to be 6.5 wt% and 8.4 wt% from TG results (Fig. S3). The N2 adsorption–desorption isotherms of Ndoped TiO2/C in Fig. S4 shows a typical Langmuir type IV characteristic of a distinct hysteresis loop, demonstrating a mesoporous structure of carbon. The surface areas (Fig. S4. a, c, e) of pristine TiO2, TiO2/C and N-doped TiO2/C are 7.70, 77.93 and 111.05 m2·g−1, respectively. Much higher surface areas of TiO2/C and N-doped TiO2/C should be mainly caused by the carbon coating. Moreover, the decomposition of melamine would release a large amount of gas during the sintering process, leading to the formation of nanopores in carbon layer. (Fig.S4. b, d, f).[37] The TEM image of N-doped TiO2/C in Fig. 3a shows that TiO2 nanoparticles have a uniform size distribution ranging from 5 nm to 50 nm. The SAED analysis (inset in Fig. 3a) confirms the high crystallinity of TiO2 nanoparticles. The HRTEM images (Fig. 3b and 3c) of the individual nanoparticle display lattice spacing of 0.22 nm, which could be indexed to the (111) plane of the rutile TiO2. It is interesting to note that another lattice spacing of 0.10 nm is also observed in the edge of a TiO2 nanoparticle. For comparison, HRTEM images of pristine TiO2 and TiO2/C are shown in Fig. 4. Notably, there are no such lattice fringes in pristine TiO2 and TiO2/C except for those with lattice
spacing of 0.20 nm. According to the XRD patterns, the lattice space of all the planes of the rutile TiO2 is larger than 0.1 nm, thus the 0.1 nm lattice spacing at the edge of the TiO2 nanoparticles cannot be originated from the TiO2 nanoparticles. We once considered that the 0.1nm lattice spacing should be caused by the TiN compound because it could be well indexed to the (400) plane of TiN. However, in XPS spectrum, we could not find the related information of low valance of Ti. Accordingly, we suspect that the lattice spacing of 0.10 nm results from a new phase causing by N-doping. For convenience, this uncertain phase is denoted as TiO2-xNy. Fig. 3d further demonstrates that the TiO2 nanoparticles are embedded in the N-doped carbon and the composite consists of TiO2, TiO2-xNy, and N-doped carbon layers. The TiO2-xNy and doped carbon layers could improve the electronic conductivity and suppress the aggregation of TiO2 nanoparticles during sodium insertion/extraction. The elemental mapping analysis of N-doped TiO2/C (Fig. S4) further confirms the homogeneous distribution of Ti, O, C and N elements among the composites. It should be mentioned that this is the first report involving the dual N-doping not only for carbon layers but also for rutile TiO2, which can greatly improve the intrinsic conductivity and electrochemical performance of rutile TiO2.[33, 52, 53] The electronic conductivities of N-doped TiO2/C, TiO2/C and pristine TiO2 are 4.87×10−4 S cm−1, 1.32×10−4 S cm−1 and 4.22×10−8 S cm−1, respectively, measured by a four-point probe method. Our results are consistent with references showing that ammonia treated TiO2 exhibited much higher Li storage capacity owe to the generation of a thin titanium nitride layer.[25, 38, 46] The particle size of N-doped TiO2/C is much smaller than those of pristine TiO2 (range from 300 to 700 nm) and TiO2/C (range from 100 to 200 nm) as demonstrated in Fig. 4. Smaller particle size could also contribute to a higher BET area (Fig. S5).
The sodium storage of all as-prepared samples were investigated. It has been reported that rutile could only reversibly uptake 0.5 Na per formula unit, thus we calculated its practical theoretical capacity (the detailed calculation was present in the supporting information) and set 1C=168 mAh g-1 in our work.[54] Fig. 5a demonstrates the galvanostatic charge-discharge curves of all samples at a current density of 84 mA g−1 after 100 cycles. It is found that the pristine TiO2 exhibits ultralow discharge capacity (~10 mAh g−1). After carbon coating, the reversible capacity increases to 58.8 mAh g−1 for TiO2/C. N-doped TiO2/C shows a highest capacity of 166.4 mAh g−1. In fact, poor lithium insertion performance was also reported in bulk rutile TiO2 due to its low intrinsic electronic conductivity and poor ion diffusion in the ab planes (Dab=10−15cm2/s), which may prevent Li ions from reaching the thermodynamically favorable octahedral sites and separates Li in the c channels [18]. Here, the significantly enhanced discharge capacity of N-doped TiO2/C may mainly originate from its improved electronic conductivity and larger surface area. In order to understand the contribution of N-doped carbon to the capacity, its charge-discharge curve was also shown in Fig. 5a. A discharge capacity of 94.8 mAh g−1 is obtained after 100 cycles, much less than that of N-doped TiO2/C. Given that the carbon content in N-doped TiO2/C is only 6.5 wt%, it is clear sure that the dominant contribution to the high capacity in N-doped TiO2/C is not from N-doped carbon layer. Fig. 5b shows the cyclability of all samples at 84 mA g−1. Obviously, N-doped TiO2/C demonstrates a great enhancement in discharge capacity compared to TiO2/C and pristine TiO2. It exhibits a high discharge capacity of 154.6 mAh g−1 after 200 cycles, while those for TiO2/C, pristine TiO2 and N-doped carbon are 65.4, 10 and 92.9 mAh g−1, respectively. As we know, the reproducibility of the battery performance is a critical factor for electrode materials, thus we tested the sodium storage performance of N-doped TiO2/C
with different bathes and its cycling performance at 84 mA g−1 was presented in Fig. S6. Interestingly, the cycling performance of N-doped TiO2/C at 84 mA g-1 of different batches show almost the same tendency and discharge capacity, indicating the excellent reproducibility of the electrode materials. The Li storage performance of all samples were also evaluated and presented in Fig. S7. Clearly, N-doped TiO2/C also shows the significantly improved Li storage properties. It delivers a high capacity of 403.3 mAh g−1 (Fig. S7b) at 168 mA g−1, and 138.6 mAh g−1 at 3360 mA g−1 (Fig. S7c). Moreover, the Li storage performance of the N-doped TiO2/C is superior to most reported pristine and modified rutile TiO2 (Table S2).[27, 55] Notably, the discharge capacities increase in the initial several cycles in both Na-ion and Li-ion batteries is caused by electrochemical activation process, that is, the reversible insertion of sodium ions is kinetically hindered thus repeated charging−discharging process is needed, which has been reported in some previous literature involving TiO2-based anode materials for Na ion batteries. [56] The charge-discharge process of Na in N-doped TiO2/C was further evaluated by cyclic voltammetry (CV) measurements (Fig. 6a). In the first cycle, a large irreversible peak is located at about 0.8 V which should be caused by the formation of solid electrolyte interface (SEI) layers and electrolyte decomposition.[57] From the second cycle on, a pair of cathodic/anodic peaks at 0.62/0.83 V corresponding to the desodiation/sodiation process is observed, in agreement with the previous reports.[56, 58] Notably, the peak area almost maintains the same after the first cycle. The irreversible capacity loss in the initial cycle has been realized as a big issue for Na-ion battery anode.[59] A pre-sodiation process, optimizing the binder and modifying the electrolyte surface have demonstrated as an effectively approach to reduce the initial capacity loss.[57]
Charge-discharge curves of N-doped TiO2/C electrode at various current densities are displayed in Fig. 6b. With the current density increasing from 16.8 mA g−1 to 336 mA g−1, the shapes of charge-discharge profiles are preserved well and the discharge capacities of 211.2, 171.5, 161.2, 122.2, 83.1 mAh g−1 are obtained at 16.8, 33.6, 84, 168 and 336 mA g−1, respectively. At each rate, no visible capacity loss is observed when the current density turns back to 16.8 mA g−1 after 45 cycles, manifesting excellent rate capability and cyclability of N-doped TiO2/C (Fig. 6c). Notably, a remarkably large capacity of 211.2 mAh g−1 is achieved at 16.8 mA g−1, equalling 0.63 Na ions insertion into a formula of rutile TiO2, which is superior to most of previously reports.[19, 20] A discharge capacity of 160.2 mAh g−1 of rutile TiO2 at 16.8 mA g−1 was reported by Zhang et al.,[19] and 160 mAh g−1 at 50 mA g−1 for Nb-doped rutile TiO2 by Usui et al.[20] Long cycling performance at different current densities was tested to further evaluate the sodium storage properties of N-doped TiO2/C electrode. As shown in Fig. 6d, the electrode demonstrates a discharge capacity of 175.3 mAh g−1 after 100 cycles at 33.6 mA g−1 (0.2C). The galvanostatic charge-discharge profiles (Fig. 6e) after different cycles repeat well without visible capacity fading, revealing high reversibility of sodium ions during the insertion/extraction process. Remarkably, a discharge capacity of 121.2 mAh g–1 at 168 mA g−1 can still be retained after 500 cycles, with a Coulombic efficiency near 100% throughout the cycling (Fig. 6f). The superior cycling stability is further confirmed by the galvanostatic discharge profiles after different cycles at 168 mA g−1 (Fig. S8). As exhibited in Table S3, amorphous TiO2 with a discharge capacity of 150 mAh g−1 at 50 mA g−1, TiO2(B) with 102 mAh g–1 at 33.5 mA g−1 and TiO2(H) with 85 mAh g−1 at 41.8 mA g−1 were reported.[11, 60, 61] The reversible discharge capacity of N-doped rutile TiO2/C in the current study is much higher than the above mentioned TiO2 polymorphs. Generally, anatase TiO2 has
been considered as a better anode material for Na-ion battery in comparison with other TiO2 polymorphs and some anatase TiO2 with excellent sodium storage performance were reported.[62-64] It is interesting to note in Table S3 that N-doped rutile TiO2/C here is comparable to some reported anatase TiO2 or TiO2/C with typical electrochemical performance.[57, 65-67] Our results demonstrated that a much better performance could be achieved on rutile TiO2 by a simple and scalable N-doping strategy. It should be noted that the synthesis of TiO2 in this work was very simple and the pristine sample exhibited poor electrochemical properties. The key point of this work is to introduce a fascinating N-doping strategy. We believe that the electrochemical properties of rutile TiO2 can be further improved if we combine the Ndoping strategy with morphology tuning of TiO2. To gain more insight into the influence of N-doping, EIS measurements were measured. As shown in Fig. 7, the Nyquist plots consist of an inclined line in the lowfrequency region and one semicircle in the high-frequency region. The fitting equivalent circuit is also attached. The inclined line represents Warburg impedance corresponding to the sodium ion diffusion resistance (W1) in a solid electrode. The small intercept at the Z′ axis demonstrates the internal resistance of electrolyte (Rs) while the semicircle corresponds to charge-transfer impedance on the electrode– electrolyte interface (R1).[31] CPE1 and C1 represent a constant phase and the doublelayer capacitance, respectively.[68] Obviously, the R1 of N-doped TiO2/C (110.5 Ω·cm−2) is much lower than that of TiO2/C (126.5 Ω·cm−2) and pristine TiO2 (215.4 Ω·cm−2), verifying an enhanced charge transfer on the interface of electrode and electrolyte (Table S4). Increasing of R1 has been demonstrated as an important contributor to the capacity fading of active materials. We expect that an improved electrochemical performance of N-doped TiO2/C could be resulted from the lower
R1.[69] Moreover, as shown in Fig. S9, the Na-ion diffusion coefficient of pristine TiO2, TiO2/C and N-doped TiO2/C are calculated to be ca. 9.9×10−15, 1.58×10−13 and 2.68×10−13 cm2 S−1, respectively. Apparently, the appearance of the TiO2-xNy layer on the TiO2 surface and more defects in both carbon layer and intrinsic rutile TiO2 are beneficial to restrain the charge-transfer resistance and increase the sodium ion diffusion coefficient, resulting in superior Na storage properties. To further understand the insertion/extraction mechanism of Na ions in the modified rutile TiO2, XRD patterns of N-doped TiO2/C electrodes at different voltages (marked in Fig. 8a) are collected in Fig. 8. For comparison, the XRD patterns of electrodes (Fig. 8b) show some additional peaks compared to those of N-doped TiO2/C powder in Fig.1, in agreement with our previous reports.[70] In the first cycle, the main diffraction peaks of (110) and (211) diffraction peaks of rutile TiO2 shifts to higher angles during the discharge process (a-d) and return to the original value during the charge process (e - g). The shifts of the peak position indicate the change of lattice parameters, which are listed in Table S5. The reversible of peak shifting implies that the crystal structure of the TiO2 phase is maintained, i.e., reversible Na ion insertion/extraction from the crystal lattice. The comparison of XRD patterns after different cycles is displayed in Fig. S10. We can see that the intensity of the diffraction peaks reduces a little after a long-term cycling, which may be caused by the repeated de-intercalation and intercalation of Na-ions, resulting in slight crystallinity decrease of the N-doped TiO2/C electrode. However, the main diffraction peaks are well maintained, indicating that there is no obvious amorphization of the TiO2 after the cycling, demonstrating the excellent structure stability of N-doped TiO2/C electrode during the cycling.
The electrochemical performance of N-doped TiO2/C was significantly improved in comparison with pristine TiO2 and TiO2/C. Possible reasons for this improvement may include: (1) the N-doped carbon layer or carbon matrix resulted in much smaller nanoparticles and abundant porosity, leading to a shorter sodium diffusion path; (2) the well-formed carbon layer could suppress the clustering of nanoparticles and alleviate the structural strain resulting from the volume change during Na ions insertion/extraction process; (3) the dual nitrogen doping to carbon layer and TiO2 can introduce more defects and oxygen vacancies, greatly enhancing the electronic and ionic conductivities and facilitating the efficient diffusion of Na ions between the carbon and rutile TiO2. As we know, ammonia treated TiO2 exhibited much higher Li storage ability due to the appearance of thin titanium nitride layer on the TiO 2 surface, which has higher electrical and thermal conductivity than the TiO2. [25, 38, 46]
4. Conclusion In summary, N-doped rutile TiO2/C was successfully synthesized by a facile N doping method. It was found that the N-doping into carbon layers and TiO2 surface was achieved together. This sample exhibited significantly improved Na storage performance as anode. It showed a high discharge capacity of 211.2 mAh g−1 at 16.8 mA g−1, and 175.3 mAh g−1 was obtained at 33.6 mA g−1 after 100 cycles. Excellent long-term cycling life with capacity retention of 92.3% after 500 cycles at 168 mA g-1 was also observed. N-doped carbon layers and TiO2-xNy coating layers on TiO2 surface greatly contributed to the enhanced Na storage performance. This work provides a crucial clue for further improving the Na storage ability of rutile phase TiO2 and the asprepared N-doped TiO2/C here can be employed as a promising anode candidate for Na-ion batteries.
Acknowledgement This work was financially supported by the National Nature Science Foundation of China (No. 21671200 and No.21571189), the Fundamental Research Funds for the Central Universities of Central South University and the Open-End Fund for Valuable and Precision Instruments of Central South University (CSUZC201729). The work at the Hong Kong University of Science and Technology was supported by a start up fund.
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Figure captions Fig.1 (a) XRD patterns, (b) FT-IR spectra, (c) ESR spectra and (d) Raman spectra of pristine rutile TiO2, TiO2/C and N-doped TiO2/C samples. Fig. 2 XPS survey spectra of pristine TiO2, TiO2/C and N-doped TiO2/C samples (a). Ti 2p (b), N 1s (c), C 1s (d) spectra of N-doped TiO2/C. Ti 2 p (e) and O1s (f) spectra of pristine TiO2, TiO2/C and N-doped TiO2/C samples. Fig.3 (a) TEM image of N-doped TiO2/C. Inset shows the corresponding SAED patterns. (b, c, d) High-resolution TEM images of N-doped TiO2/C. Fig.4 (a) TEM image of pristine TiO2 (a) and TiO2/C (b). High-resolution TEM images of pristine TiO2 (c) and TiO2/C (d). Fig.5 (a) Charge-discharge curves of N-doped carbon, N-doped TiO2/C, TiO2/C and rutile TiO2 after 100 cycles at 84 mA g−1. (b) Cycling performance of N-C, N-doped TiO2/C, TiO2/C and rutile TiO2 at 84 mA g−1. Fig.6 (a) CV curves of as-prepared N-doped TiO2/C between 0.001 and 3.0 V at 0.4 mV s−1, (b) Reversible charge-discharge curves of N-doped TiO2/C at various current densities from 16.8 to 336 mA g−1, (c) Rate performance of N-doped TiO2/C, (d) Cycling performance of N-doped TiO2/C at 33.6 mA g−1, (e) Charge-discharge profiles of N-doped TiO2/C after different cycles at 33.6 mA g−1, (f) Cycling performance and Coulombic efficiency of N-doped TiO2/C at 168 mA g−1. Fig.7 Electrochemical impedance spectroscopys of pristine TiO2, TiO2/C and N-doped TiO2/C electrodes after 3 cycles at 33.6 mA g−1.
Fig. 8 (a) Charge-discharge curves of N-doped TiO2/C at 168 mA g−1. (b) XRD patterns of N-doped TiO2/C at different voltages: a= 3.0 V, b=1.8 V, c=1.2 V, d= 0.01 V, e= 0.8 V, f=1.3 V, g= 3.0 V at 168 mA g-1. Patterns in (c) and (d) are the partially enlarged portions of the highlighted parts in (b).
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