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Energy Material Advances / 2022 / Article

Research Article | Open Access

Volume 2022 |Article ID 9876319 | https://doi.org/10.34133/2022/9876319

Fujie Li, Chao Wang, Xiu Song Zhao, "Sodium-Ion Storage Properties of Thermally Stable Anatase", Energy Material Advances, vol. 2022, Article ID 9876319, 12 pages, 2022. https://doi.org/10.34133/2022/9876319

Sodium-Ion Storage Properties of Thermally Stable Anatase

Received15 Jun 2022
Accepted26 Aug 2022
Published06 Oct 2022

Abstract

Anatase titanium dioxide (TiO2) is a potential anode material for sodium-ion batteries (NIBs). However, the low electronic conductivity and sluggish ion diffusion kinetics at high rate hamper its practical applications. Herein, we demonstrate a sol-gel approach to the synthesis of thermally stable anatase nanoparticles with a carbon shell as anode materials for NIBs. A sample calcined at 750 °C (designated as H-750TiO2@C) exhibits high-rate capability and excellent stability against cycling with no capacity loss after 2000 cycles at 1 A g-1. In situ X-ray diffraction and Raman spectroscopy characterization results reveal a nearly zero-strain characteristic of the anatase phase during charge/discharge processes. In situ transmission electron microscopy, ex situ X-ray photoelectron spectroscopy, and scanning electron microscope characterization results of samples collected at different charged and discharged states suggest that the anatase phase undergoes an irreversible sodiation-activation during the initial discharge process to form a sodiated-TiO2 phase. A full cell assembled with H-750TiO2@C as the anode and Na3V2(PO4)3 as the cathode delivers an energy density of 220 Wh kg-1, demonstrating H-750TiO2@C is a potential anode material for NIBs.

1. Introduction

Sodium-ion batteries (NIBs) for large-scale energy storage applications attract increasing attention due to naturally abundant sodium resources [13]. However, the larger radius and heavier molar weight of sodium ion (Na+) than lithium ion (Li+) lead to fundamentally different requirements for electrode materials [4]. For example, graphite, the most popular anode for Li+ batteries (LIBs), has limited capacity in storing Na+ [5]. Therefore, the development of high-performance anode materials is one of the keys to realizing the NIB technology [6, 7].

Titanium dioxide (TiO2) is a promising anode material for NIBs owing to its non-toxicity, low cost, and the abundance of titanium in nature [8, 9]. There are four TiO2 polymorphs, including rutile, anatase, bronze, and amorphous TiO2 that have been studied as anode materials for NIBs [1013]. Among them, anatase is more conductive in storing sodium ions due to its special stacking of TiO6 octahedra with two-dimensional (2D) channels for Na+ transport [14]. However, the electron conductivity and ion diffusion of anatase are poor, significantly limiting its rate capability and long-term cycling performance.

To overcome these challenges, strategies such as coating conductive materials, nano-structuring, and constructing porous structure have been reported [1517]. Zhou and co-workers demonstrated a carbon-modified TiO2 fiber material witch enhanced electron conductivity and improved specific capacity [18]. Wang et al. [19] observed high-rate performance of three-dimensional anatase nanofibers with Na+ storage capacity of about 173 mAh g-1 at 15 A g-1. Li et al. [15] reported that porous anatase TiO2 provides channels for Na+ transport, thus enhancing the rate performance with a specific capacity of 116 mAh g-1 at 20 °C.

In this paper, we demonstrate a sol-gel approach to the synthesis of thermally stable anatase (up to 750 °C) coated with carbon and acid-etched with 36.5 wt% HCl (hereinafter designated as H-750TiO2@C) as anode for NIBs. Acid etching can create channels for electrolyte ion transport and increase interfacial surface area for charge storage [20]. The H-750TiO2@C electrode exhibits a reversible specific capacity of 228 mAh g-1 at a current density of 0.05 A g-1 with 100% capacity retention after 2000 cycles at 1 A g-1. Both in situ X-ray diffraction (XRD) and Raman spectroscopy results reveal a nearly zero-strain characteristic of H-750TiO2@C during charge/discharge processes. In situ transmission electron microscopy (TEM), ex situ X-ray photoelectron spectroscopy (XPS), and scanning electron microscopy (SEM) results suggest an irreversible sodiation-activation process to form a sodiated-TiO2 phase during the initial discharge process. A full coin cell assembled with H-750TiO2@C as the anode and Na3V2(PO4)3 as the cathode delivers an energy density of 220 Wh kg-1.

2. Experimental

2.1. Chemicals

All chemicals, including tetrahydrofuran (THF, 99%, Macklin Biochemical Corporation), pluronic F127 (PEO106PPO70PEO106, , Sigma-Aldrich Corporation), HCl (36 wt.%, Sinpharm Chemical Reagent Corporation), ethylene glycol (EG, 98%, Macklin Biochemical Corporation), titanium tetrabutoxide (TBOT, ≥99%, Aladdin Industrial Corporation), ethanol (95%, Aladdin Industrial Corporation), polyvinylidene fluoride (PVDF, Solvay Corporation), Super P (Timcal Super C65), and N-methyl pyrrolidone (NMP, Aladdin Industrial Corporation) were used as received.

2.2. Materials Synthesis

The carbon coated H-750TiO2@C samples were synthesized in an acidic solution containing THF, F127, H2O, HCl, EG, and TBOT by sol-gel method. For a typical synthesis, 1.6 g F127, 0.2 g deionized water, 2.4 g EG, and 2.4 g concentrated HCl were dissolved in 30 mL THF. The mixture was stirred vigorously for 30 min to form a milky suspension. Sequentially, 3.4 g TBOT was dropwise added under vigorous stirring for 30 min to form a clear solution with a golden yellow color. Then, the solution was transferred into a volumetric flask immersed in an oil bath at 40 °C for 20 h, followed at 150 °C for 12 h. The part of collected samples were calcined at X °C (, 550, 650, 750, 800, and 900) under argon atmosphere for 2 h to obtain a sample designated as XTiO2@C. And the other part of collected samples was calcined at 750 °C for 2 h under an air atmosphere to obtain a sample designated as 750TiO2.

To obtain the channel structure, part of the XTiO2@C was etched off using HCl [20]. 1 g XTiO2@C was hydrothermally treated in 20 mL of 2 M HCl solution at 160 °C for 1 h. The solid was collected by centrifugation (8000 rpm, 30 min), followed by several washes with deionized water and ethanol, and finally freeze-dried at − 65 °C to obtain a sample, which is designated as H-XTiO2@C.

To verify the capacity contribution of the carbon in H-750TiO2@C, the H-750TiO2@C washed using 10 M H2SO4 (30 mL) at 80 °C for 6 h to remove the TiO2 component, leaving behind residual carbon (designated as H@C). The TGA curve in Figure S6b conformed all TiO2 were removed to leave behind pure residual carbon.

A NIB cathode material, Na3V2(PO4)3 (NVP), was prepared using the method as described previously [21].

2.3. Characterization

X-ray diffraction (XRD) patterns were collected on an Ultima IV (Rigaku) with Cu Kα radiation at a scan rate of 10° min-1. Field-emission scanning electron microscopy (FESEM, JEOL JSM-7800F) and transmission electron microscopy (TEM, JEOL JSM-2100Plus) were employed to characterize particle morphology and size. The Raman spectra were obtained at the laser wavelength of 532 nm on a Raman spectroscopy system (Renishaw). Thermogravimetric analysis (TGA, Mettler Toledo) was performed in air at a heating rate of 10 °C min-1. Nitrogen adsorption/desorption isotherms were measured on a Quantachrome Autosorb iQ3. The specific surface area and pore size distributions were computed using the Brunauer-Emmett-Teller (BET) and Density-Function-Theory (DFT) methods, respectively. Surface analysis was performed on an X-ray photoelectron spectroscopy (XPS, PHI5000 Versaprobe III) with an X-ray source of Al Kα (All samples tested by ex situ XPS were etched with argon ions to 200 nm). The hydrothermally treated samples were separated using centrifugation (Thermo Sorvall ST16R).

2.4. Electrochemical Measurements

The electrodes (H-XTiO2@C, XTiO2@C and H@C) were fabricated by mixing the as-prepared active material (80 wt%) with Super P (10 wt%) and PVDF (10 wt.%) in NMP. The obtained slurry was painted on the copper foil and dried in a vacuum oven at 110 °C for 12 h. Then the foil was cut into discs with a loading of ∼1.0 mg cm-2 of active material. CR2032 coin-type half cells using glass microfibers (Whatman) as separator and sodium metal as counter electrode were assembled in an argon-filled glove box (Mikrouna, H2O, O2<0.01 ppm). The electrolyte was a solution of 1 M NaClO4 dissolved in polycarbonate/ethylene carbonate (DMC/EC, ) and 5 vol.% fluoroethylene carbonate (FEC). The full cell was assembled with the same procedures except for the replacement of sodium metal by NVP. The active material mass loading of anode and cathode were 0.77 mg cm−2 and 1.71 mg cm−2, respectively. The N/P ratio was about 1 : 1.28. The electrolyte addition amount was 60 μL for both half and full cells. For TiO2-based electrode materials, an aqueous binder is of a better choice than PVDF [22, 23]. However, for comparison purposes with other TiO2 electrode materials reported in the literature, PVDF was used as binder in this work.

For the in situ measurement, the assembly of the in situ test device was identical to the half cell except that the in situ XRD electrode collector was a beryllium sheet, the in situ Raman electrode without a collector and the laser (532 nm) scanning through a 0.5 cm diameter quartz window. The BTS-400 battery test system was used to perform GCD tests on in situ devices with a current density of 0.1 A g-1 and a voltage range of 0.01-2.5 V. And each in situ XRD pattern had a two-theta scan range of 20-55° with a scan rate of 10° min-1. The sweep range of each in situ Raman spectrum was between 50 and 800 cm-1 with a scan interval of 5 min. The structural and morphological evolution of H-750TiO2@C was observed in real time by in situ TEM characterization performed on an X-mech® electrochemical holder [24]. The two electrodes of the in situ TEM half cell were H-750TiO2@C particles on a silver rod and sodium metal on a tungsten rod, respectively. The naturally occurring sodium oxide (Na2O) coating on the surface of the sodium metal was used as a solid electrolyte. When the two electrodes were in contact, a negative voltage of − 2.5 V was applied to achieve the process of sodiation (sodium metal → Na2O → H-750TiO2@C).

Cyclic voltammetry (CV) curves were collected by an electrochemical workstation (CHI760D, Chenmarine, Shanghai, China) at different scan rates over a voltage range of 0.01 to 2.5 V. The Bio-Logic electrochemical workstation (SP-150, Bio-Logic SAS) was used to measure for electrochemical impedance spectroscopy (EIS) with an AC stimulus of 10 mV and amplitude frequencies in the range of 100 kHz to 0.01 Hz. Galvanostatic charge-discharge (GCD) and galvanostatic intermittent titration technique (GITT) analyses were performed using a Neware instrument (CT-4008) over a voltage range of 0.01 to 2.5 V for half cells and 1 to 3.5 V for full cells. All electrochemical measurements were carried out at room temperature.

3. Results and Discussion

Figure 1(a) and Figure S1 show the XRD patterns of H-XTiO2@C and XTiO2@C (X represents the calcination temperature). All the XRD diffraction peaks are well indexed to anatase phase (PDF # 84-1286) for the samples with calcination temperatures below 800 °C. However, the impurity peaks of rutile (PDF # 04-0551) appear when the calcination temperature reaches 800 °C. Thermally stable anatase can be achieved at 750 °C without the phase transition to Na+-sluggish rutile. Moreover, with the calcination temperature increasing from 450 °C to 750 °C, the resultant samples exhibit shaper diffraction peaks, suggesting higher crystallinity of anatase [14].

The Raman spectra of the H-750TiO2@C and 750TiO2@C samples are shown in Figure 1(b). Two distinct peaks at 1340 cm-1 and 1590 cm-1 for both samples are owing to the D band and G band of the coated carbon [25, 26]. Other peaks from 152 to 625 cm-1 are typical for anatase TiO2 [27]. The Eg peaks at 152 cm-1, 204 cm-1, and 625 cm-1 are caused by symmetric stretching vibrations of Ti-O bond on the (101) plane [27]. B1g peak (398 cm-1) and A1g peak (508 cm-1) are ascribed to the symmetric bending vibrations and antisymmetric bending vibrations of Ti-O bond on the (001) plane, respectively [28]. It is also seen that the peak intensity of sample H-750TiO2@C increased significantly, indicating the increased exposure of (101) and (001) planes after hydrothermal etching. The increase of exposure planes leads to the increase of Na+ diffusion channels, improving the Na+ storage capacity.

The thermogravimetric analysis (TGA) curves of H-750TiO2@C and 750TiO2@C are shown in Figure 1(c) and Figure S2, respectively. The weight loss below 100 °C is mainly due to the elimination of moisture from the sample. The weight loss of H-750TiO2@C between 100 and 710 °C was about 24 wt%, which is due to the burning off of carbon in air [29]. Sample 750TiO2@C contained about 20 wt% of carbon, slightly less than that of H-750TiO2@C, indicating the loss of the polymer template during the HCl etching process.

Figures 1(d) and S3 show the nitrogen (N2) physisorption isotherms for H-750TiO2@C and 750TiO2@C. The sharp N2 uptake at low relative pressure suggests the presence of micropores [30]. The noticeable hysteresis loops in the middle relative pressure indicates the existence of mesopores [31]. The specific surface areas calculated using the Brunauer-Emmett-Teller (BET) method are 225.6 and 121.1 m2 g-1 for H-750TiO2@C and 750TiO2@C, respectively. The insets of Figure 1(d) and Figure S3 show the density functional theory (DFT) pore size distribution curves, confirming the presence of micropores and mesopores with total pore volumes of 0.23 and 0.13 cm3 g-1 for H-750TiO2@C and 750TiO2@C, respectively. The results of N2 physisorption suggest that the acid-etching increased the specific surface area and pore volume. It should be noted that high specific surface area and large pore volume can increase the formation of SEI, resulting in a low ICE, especially for micropores. However, a hierarchical porous structure predominately with mesopores [32] could effectively facilitate Na+ transport while minimize the effect on ICE [15].

Figures 1(e) and 1(f) show the field-emission scanning electron microscopy (FESEM) images H-750TiO2@C and 750TiO2@C. Apparently, HCl etching effectively produced porous channels in H-750TiO2@C, which would facilitate Na+ transport. The transmission electron microscopy (TEM) images in Figures 1(g) and 1(h) show that the sponge-like H-750TiO2@C microstructure consists of TiO2 nanoparticles of about 10 nm in diameter. It was reported that anatase nanoparticles of sizes less than 11 nm are thermally stable without phase transformation [33]. On the other hand, the coated carbon shell also plays a role in stabilizing the anatase nanoparticles by absorbing the strain during thermal treatment [34]. The lattice spacing is about 0.35 nm, corresponding to the (101) plane of anatase TiO2 [35, 36]. In addition, a carbon layer as highlighted by the yellow arrows can be seen on the surface of the H-750TiO2@C nanoparticle. The selected area electron diffraction (SAED) patterns in Figure 1(h) show the anatase has a high crystallinity [29].

Figure S4 show the energy dispersive X-ray (EDX) mapping images, confirming the carbon coating on the H-750TiO2@C nanoparticles. The coated carbon walls have an important impact on the nanocrystal morphology of H-750TiO2@C and retard the transformation of anatase to rutile [37].

The XPS spectrum of the H-750TiO2@C sample shown in Figure S5a demonstrates the presence of elements Ti, O, and C [38]. The Ti 2p spectrum of H-750TiO2@C shown in Figure S5b displays two peaks at 459.5 eV and 465.5 eV, corresponding to Ti4+ 2p3/2 and Ti4+ 2p1/2, respectively [10]. The C 1s spectrum of H-750TiO2@C shown in Figure S5(c) is split as three peaks of 282.8 eV, 286.1 eV, and 288.8 eV, belonging to C-C, C-O and O-C=O, respectively [39]. An H-750TiO2@C high-resolution O 1s spectrum was divided into two peaks at 530.9 eV and 532.6 eV (Figure S5(d)), responding to the C=O bond of coated carbon and the Ti-O bond of anatase TiO2, respectively [40].

Figure 2(a) shows the first five cyclic voltammetry (CV) profiles of electrode H-750TiO2@C at a scan rate of 0.1 mV s-1 over a voltage range between 0.01 and 2.5 V. In the first cathodic scan, there is a broad peak at 1.2 V related to the irreversible formation of solid-electrolyte interphase (SEI) [40]. This peak disappeared in the subsequent cycles, indicating the formation of the SEI is completed in the initial cycle. In addition, the weak peak at 0.6 V and the sharp peak at 0.01 V are associated with the insertion of Na+ into anatase TiO2 and the side reaction between anatase TiO2 and electrolyte. The gradual reduction of the peak intensity from the 1st to the 5th cycle indicates the gradual weakening of the side reaction [41]. In the subsequent cycles, the CV curves almost overlap with two main peaks at 0.55 V and 0.8 V, which are characteristic redox peaks of anatase TiO2 [42], indicating a good electrochemical reversibility of electrode H-750TiO2@C. The small peak at 0.6 V is ascribed to the reversible sodium release from Super P [43].

Figure 2(b) shows the GCD curves of H-750TiO2@C at a current density of 0.05 A g-1. The first discharge curve consists of three voltage regions (the reasons will be described later), between the open circuit voltage (OCV) and 1.5 V (Region I), 1.5-0.6 V (Region II) and 0.6-0.01 V (Region III), corresponding to the pseudocapacitance process as well as the formation of sodiated-TiO2 process, the formation of irreversible SEI layer, and the sodiation process, respectively [2]. The H-750TiO2@C electrode displays a low initial Coulombic efficiency of 44%, which is ascribed to the formation of SEI films and the consumption of Na+ during the sodiation-activation process to form sodiated-TiO2 [43]. It is worth noting that the charging specific capacity gradually increases to a reversible specific capacity of 228 mAh g-1 at 0.05 A g-1 in the subsequent cycles as shown in Figure S6a. The carbon coated H-750TiO2@C electrode needs several cycles to be completely activated [44].

To investigate the capacity contribution of the carbon component to the observed total capacity of the H-750TiO2@C electrode, the TiO2 was washed away using 10 M H2SO4 (30 mL) at 80 °C for 6 h (designated as H@C). TGA curve of H@C is shown in Figure S6(b); the carbon was completely burned out, indicating that the TiO2 in the H-750TiO2@C sample was completely washed away. The GCD profiles of the H@C are shown in Figure S6(c). It is seen that the carbon had little Na+ storage capacity with only about 7.0 mAh g-1 at the first cycle and dropped to 2.0 mAh g-1 at the tenth cycle, confirming that the main contribution to the observed specific capacity of electrode H-750TiO2@C came from TiO2, while the 10 wt% Super P added to the electrode indeed contributes to Na+ storage. However, the contribution is small with only about 10 mAh g-1.

The rate capabilities of the H-750TiO2@C, 750TiO2@C, and 750TiO2 electrodes are compared in Figure 2(c). It is seen that the H-750TiO2@C electrode exhibits the highest rate capability at 1 A g-1 with a specific capacity of 94 mAh g-1. Moreover, the specific capacity returned to 205 mAh g-1 when the current density was decreased to 0.1 A g-1, indicating the structural stability of H-750TiO2@C during the sodiation/desodiation processes.

Comparing H-750TiO2@C with 750TiO2@C, acid-etching formed sponge-like porous channels increased the specific surface area and pore volume, which can enhance the electrolyte/electrode contact area and facilitate the Na+ transport, thus improving the rate performance at low current densities. However, at high current densities (e.g., 1 A g-1), electron conductivity is the factor of limiting the rate of Faraday reactions. The low electron conductivity of TiO2 causes severe polarization of the electrode at high current density, leading to a decrease in capacity. Because of this, H-750TiO2@C and 750TiO2@C exhibited a similar specific capacity at 1 A g-1.

Figure S6(d) shows the rate capability of the H-XTiO2@C (X represents calcining temperatures of 450, 550, 650, 750, or 800) electrodes. The rate and cycling performances of H-XTiO2@C improve with increasing calcining temperature from 450 to 750°C. This can be ascribed to the improved crystallinity of anatase. When the calcination temperature was further increased to 800 °C, the sample obtained (H-800TiO2@C) shows low specific capacity and poor rate capability because of the phase transition to rutile. It is interesting to note that the specific capacity of both H-750TiO2@C and H-650TiO2@C electrodes continuously increased in the first ten cycles. This can be ascribed to the continuous activation of the porous channels in the electrodes (formed during the acid washing). These pore channels have a strong resistance to electrolyte diffusion, requiring time to be activated for storing Na+. However, the low-crystallinity for H-550TiO2@C and H-450TiO2@C, and the rutile impurity phase for H-800TiO2@C induced the poor cycling performance, resulting in the different trends with H-750TiO2@C and H-650TiO2@C electrodes.

Figure S7 shows the rate performances of H-750TiO2@C with different mass loading at different current densities. The electrode with a mass loading of 1 mg cm-2 (Figure 2(c)) exhibited the best rate performance among these electrodes, which may be more conducive to illustrating the charge mechanism.

Figure 2(d) compares the cycling performance of the H-750TiO2@C and 750TiO2@C electrodes at a current density of 1 A g-1. After 2000 cycles, electrode H-750TiO2@C exhibited a capacity retention of 100% with a stable coulombic efficiency of about 100%, indicating excellent reversibility. The XRD patterns of the H-750TiO2@C electrodes after 3, 10,100, and 2000 cycles at a current density of 1 A g-1 are shown in Figure S8. It can be seen that the crystal structure of the material does not change much in the first 100 cycles. However, after 2000 cycles, the crystal structure of the material was damaged due to the repeated insertion/extraction of sodium ions in the material, resulting in the transformation of the material from crystalline phase to amorphous phase, which led to the decrease of the sodium storage performance of the material. However, for the 750TiO2@C electrode, only 60% of the initial specific capacity remained after 2000 cycles. The sponge-like porous channels in H-750TiO2@C can buffer the volume changes during Na+ insertion/extraction resulting in the excellent cycling stability.

Figure 2(e) and Figure S9 show the Nyquist plots of H-750TiO2@C, 750TiO2@C, and H-750TiO2@C at 2000th cycle with the equivalent circuit models. The semicircles and sloping lines in both plots refer to the impedance of charge transfer () and the resistance of solid diffusion () at the electrode/electrolyte interface, respectively. The equivalent series resistance, designated as , represents the resistance of the electrode and electrolyte. Table S1 lists the fitted values of , , and using the equivalent circuit model. The small value of electrode H-750TiO2@C suggests a low charge transport resistance for Na+ because of the abundantly porous channels by hydrothermal etching. At 2000th cycle, the H-750TiO2@C electrodes exhibited a smaller slope than that of before cycling, demonstrating the gradual decrease of the diffusion resistance of Na+ in the electrode internal [45, 46].

The galvanostatic intermittent titration technique (GITT) was used to investigate the dynamic behavior of electrode H-750TiO2@C in NIBs. The potential responses with the Na+ diffusion coefficients of the H-750TiO2@C electrode at the second cycle and at 2000th cycle are shown in Figures 3(a) and 3(b). The overpotentials for both electrodes are initially stable but increase at the end of the charging step. The increase of sodium ion diffusion coefficient is related to sodiation to desodiation during discharging/charging. When the sodiation is changed to desodiation, the potential difference is restored to the maximum, resulting in a dramatic increase in the diffusion coefficient of Na ions. The potential difference is high at the beginning of sodiation and desodiation, leading to fast Na ion diffusion. The potential difference gradually decreases with the sodiation and desodiation progressing, resulting in a gradual decrease in Na ion diffusion coefficient. The diffusion coefficients of Na+ (D Na+) are calculated using the following equation [47, 48]: where is the constant current pulse time and , , and refer to the mass, molar volume, and molar mass of the anode materials, and electrode/electrolyte interface area, respectively. is the voltage change in a single-step including the relaxation, while refers to the difference of cell voltage in a constant current pulse step excluding the IR drop.

During the sodiation process, the gradually decreases and then stabilizes, corresponding to the Na+ insertion in H-750TiO2@C. At a high desodiation level, decreases rapidly, resulting from the large polarization observed in the GITT curve. The mean values of of the H-750TiO2@C electrodes at at 2nd, 500th, 1000th, 1500th, and 2000th cycles are summarized in Table S2. With the increasing cycle, the Na+ diffusion coefficient increased to a maximum at the 1500th cycle and then decreased. This phenomenon explains the gradual increase and then falling of the specific capacity during the long cycle in Figure 2(e), which is also in agreement with the EIS results in Figure S9. Notably, electrode H-750TiO2@C exhibits a high among the previously reported values in the literature [4951]. The unique morphology of H-750TiO2@C with sponge-like porous microstructure provides a large amount of channels for Na+ transport, significantly improving the charge transport kinetics.

To estimate the contribution of capacitance and diffusion mechanisms to charge storage, CV tests were performed on the H-750TiO2@C electrode. As shown in Figure 3(c) and Figure S10, the broad redox peaks were observed at sweep rates from 0.2 mV s-1 to 10 mV s-1. The relationship between the peak current () and the scan speed () during the CV cycle is in accordance with the following equation [52, 53]:

A value of 0.5 implies a diffusion-controlled process, while a value of 1 indicates a capacitance-controlled process. Figure 3(d) is a plot, the slope of which is the value. The values at sweep rates from 0.2 mV s-1 to 10 mV s-1 are around 0.8, showing that the sodium storage kinetics are controlled by a combination of pseudocapacitance and diffusion processes.

Figure 3(e) shows a plot of H-750TiO2@C electrode normalized capacity versus v-1/2, with two regions divided according to the slope. The decrease in capacity with increasing scan rate is not significant with scan rates less than 10 mV s-1 (green area), indicating that the charge transfer is independent of the rate. In contrast, the capacity drops sharply at scan rates higher than 10 mV s-1 (yellow region), showing a rate-limiting diffusion process for sodium storage. At a fixed potential, the contribution of the capacitance can be calculated from the following equation [54]: where denotes the peak current, stands for the capacitance control current, and for the diffusion control current at a fixed potential. The contribution of capacitive partial can be calculated by fitting and at a fixed potential. The inset shows the pseudocapacitance contribution at a sweep rate of 10 mV s-1. The capacitance CV curve (yellow area) and the original CV curve (green area) were calculated from the above equation and based on the integral calculation. The capacitive contribution is about 73.5% of the total Na+ storage capacity. As shown in Figure 3(f), the diffusion control capacitance mechanism ratios were also calculated for other scan rates. The ratio of the capacitive charge storage (pseudocapacitance) contribution increases gradually from 25.3% at 0.1 mV-1 to 73.5% at 10 mV-1 with the scan rate increasing. The high ratio of pseudocapacitive behavior promotes fast surface electrochemical reactions as the active material suffers little structural degradation, resulting in high rate performance and long-cycle stability [55].

Figure 4(a) shows the in situ XRD patterns at OCV. During the sodiation and desodiation processes, no addition peaks were observed, indicating high stability of the H-750TiO2@C electrode. In addition, the main diffraction peaks exhibited no shift, demonstrating a nearly zero-strain characteristic and reversible Na+ insertion and extraction processes in H-750TiO2@C. The corresponding surface color mapping in the two-theta range between 23.5 and 26.5° is shown in Figure 4(b). It is seen the peak intensity of the (101) plane () gradually decreases during the initial sodiation process. There is no change in the subsequent cycle. Figure S11 shows the in situ Raman spectra during the initial discharge/charge processes. The peak at 152 cm-1 is assigned to the stretching vibrations of Ti-O bond on the (101) plane in H-750TiO2@C. The intensity of this peak exhibited the same tendency with the in situ XRD results. This phenomenon is ascribed to a sodiation-activation process of H-750TiO2@C to form sodiated-TiO2, which is supported by the following ex situ XPS and SEM results.

Figure 5(a) shows the ex situ XPS spectra of samples collected at different charge/discharge states (Figure 5(b); all samples were treated with argon ion etching to a depth of 200 nm). The pristine Ti 2p spectrum (State 1) shows two characteristic peaks at 465.5 and 459.5 eV, corresponding to the Ti 2p1/2 and Ti 2p3/2 of Ti4+, respectively. With the progress of discharging, characteristic peaks for Ti 2p1/2 and Ti 2p3/2 of Ti3+ appear at 463.5 and 457 eV, respectively, indicating the gradual transformation from Ti4+ to Ti3+ with the insertion of Na+. We focused our attention on the ratio of Ti3+: Ti4+ to explore the amount of Na+ stored in TiO2. By fitting the XPS spectra of discharged to 1.5 V (State 2), the ratio of Ti3+: Ti4+ was found to be 0.27, which corresponds to about 0.21 Na+ inserted in per TiO2 (Na0.21TiO2). When discharged to 0.01 V (State 4), this ratio increased to about 2.33, which corresponds to 0.70 Na+ inserted in per TiO2 (Na0.70TiO2). During the subsequent charging process this ratio decreases to about 0.27, and at the end of the process, there was still about 0.21 Na+ in per TiO2 (Na0.21TiO2), indicating that about 0.5 Na+ can be reversibly stored per TiO2. This means that there was about 0.2 Na+ left because of the irreversible sodiation-activation process during initial discharge process to form a sodiated-TiO2 phase following 0.2 Na+ + TiO2+ 0.2 e- → Na0.2TiO2. This sodiation-activation process explains why the peak intensity of the (101) crystal plane is reduced in the in situ XRD pattern and Raman spectrum. It is also the reason causing a lower specific charge storage capacity than the theoretical specific capacity, as well as a poor initial Coulombic efficiency.

The Ti 2p XPS spectra of the H-750TiO2@C electrode discharged to 0.01 V with different etching depths between 0 and 200 nm are shown in Figure S12(a). It is seen in Figure S12(b) that the Ti3+: Ti4+ ratio does not change with the etching depth, and the thickness of the SEI layer is about 100 nm, confirming a 200-nm etching depth can guarantee to eliminate the effect of SEI on the Ti3+: Ti4+ ratio in the bulk of H-750TiO2@C in Figure 5(a). The existence of Ti3+ in the electrode charged to 2.5 V indicates that some Na+ remained in the electrode for charge balance, confirming the formation of sodiated-TiO2 phase [2].

The ex situ SEM images of H-750TiO2@C at different charge/discharge states are shown in Figures 5(c1)–5(c7). During the sodiation process (Figures 5(c1)–5(c4)), spherical particles with diameters between 100 and 200 nm appeared, and these spherical particles remained during the subsequent desodiation process (Figures 5(c5)–5(c7)). Figure 5(c8) shows the SEM image of electrode H-750TiO2@C at 2000th cycles at 1 A g-1, demonstrating that the spherical particles were kept well after 2000 cycles. This indicates that the electrode H-750TiO2@C has highly stable cycling performance. Figure S13 shows the EDX mapping results of electrode H-750TiO2@C at different discharge (at1.5, 0.6, 0.01 V) and charge (0.8 V) states, confirming the sodiation-activation process to form spherical sodiated-TiO2 during the sodiation/desodiation process.

To investigate the formation of the spherical particles seen from Figure 5(c), in situ TEM characterization of H-750TiO2@C during the sodiation process was performed, and the results are shown in Figure 6, and a video is shown in Supporting Information (available here). When the in situ biasing increases, the Na+ ions originated from the sodium metal pass through the Na2O solid electrolyte and then insert into H-750TiO2@C. The H-750TiO2@C aggregated to a spherical shape with the deepening sodiation, matching the ex situ SEM in Figure 5(c). The diffractogram in the inset of Figure 6(d) shows the appearance of SAED rings, suggesting the formation of amorphous materials including SEI film and sodiated-TiO2 during the initial discharge process. This is responsible for the peak intensity reduction of the (101) plane observed from the in situ XRD and Raman characterization results.

With the above characterization results, the three-stage sodiation process (Figure 2(b)) is interpreted below. The in situ XRD and Raman results show that a very small amount of charge is consumed when discharging to 1.5 V (the voltage drops quickly) and no structural changes are observed. Due to the large specific surface area of carbon coated TiO2, there is a capacitive behavior, which is also confirmed by the calculation of the capacitance contribution by the CV test at different sweep rates. In addition, the appearance of spherical particles can be observed in ex situ SEM, as well as the appearance of Ti3+ in ex situ XPS, with about 0.2 Na ions irreversibly inserted into per TiO2 to form sodiated-TiO2. When discharging to 0.6 V, the in situ XRD, Raman and ex situ XPS does not change much, while a reduction peak appears in the CV test at 0.1 mV s-1, which is the formation of the SEI. When the discharge is continued to 0.01 V, it can be observed that the Ti3+ content gradually increases with the further discharging degree in the ex situ XPS spectra, indicating that this stage is the insertion process of Na+. By calculating the ratio of Ti3+ and Ti4+, there are about 0.7 Na ions are inserted into per TiO2 (including 0.2 Na ions irreversibly inserted and 0.5 Na ions reversibly inserted). The above analysis divides the storage mechanism of Na+ in TiO2 into three stages: between the open circuit voltage (OCV) and 1.5 V (Region I), 1.5 and 0.6 V (Region II), and 0.6 and 0.01 V (Region III), corresponding to the pseudocapacitance process as well as the formation of sodiated-TiO2 process, the formation of irreversible SEI layer, and the sodiation process, respectively.

To demonstrate the practical application of H-750TiO2@C, a full cell (NVP//H-750TiO2@C) was fabricated with H-750TiO2@C as the anode and Na3V2(PO4)3 (NVP) as the cathode. As shown in Figure S14, the NVP electrode exhibited a reversible specific capacity of 98 mAh g-1 at 0.1 A g-1 in a half cell. Figure S15a shows the GCD profiles of the NVP//H-750TiO2@C full cell at 0.1 A g-1 in the voltage range between 1 and 3.5 V. This full cell delivered a reversible discharge capacity of 114 mAh g-1 (on the basis of active material in the anode). Figure S15b shows the rate performance of the NVP//H-750TiO2@C cell, exhibiting good rate capability with discharge specific capacities of 114, 81, and 58 mAh g-1 at current densities of 0.1, 0.2, and 0.5 A g-1, respectively. The energy density of NVP//H-750TiO2@C was calculated to 220 Wh kg-1 (on the basis of the active material in the anode), which is higher than that of other NIBs reported in Table S3. These results demonstrate the great promise of H-750TiO2@C for the practical application of NIBs.

4. Conclusions

In summary, carbon-coated anatase nanoparticles with high thermal stability were synthesized from the sol-gel approach in the presence of surfactant F127. The anatase-based composite material exhibits excellent electrochemical performance for sodium ion storage because of the high crystallinity, the presence of a carbon layer on the particle surface, and the sponge-like microstructure with porous channels. Importantly, the material is very stable against cycling with 100% capacity retention after 2000 cycles at 1 A g-1. In situ and ex situ characterization results reveal that a sodiation-activation process occurs during the initial sodiation, resulting in a low initial Coulombic efficiency. A full cell assembled with the carbon-coated anatase as anode and Na3V2(PO4)3 as cathode delivers an energy density of 220 Wh kg-1. This work provides an approach to the synthesis of high-performance titanium dioxide-based anode materials with excellence cycling stability for sodium-ion storage.

Data Availability

All data presented in the paper and the supporting information are available from the corresponding author upon reasonable request.

Conflicts of Interest

The authors declare that there is no conflict of interest regarding the publication of this article.

Authors’ Contributions

X.S.Z. and C.W. proposed and designed the project. F.J.L. performed the experiments. All the authors contributed to the write-up of the article.

Acknowledgments

This work was supported by the start-up funding of Qingdao University (No. DC2000005025).

Supplementary Materials

Figure S1: XRD patterns of XTiO2@C (X represents the calcination temperature). Figure S2: TGA curve of 750TiO2@C. Figure S3: Nitrogen sorption isotherms and DFT pore size distribution curve (the inset) of 750TiO2@C. Figure S4: EDX mappings of H-750TiO2@C. Figure S5: XPS survey spectrum (a) and high-resolution XPS spectra of Ti 2p (b), C 1s (c), and O 1s (d) of H-750TiO2@C. Figure S6: (a) GCD profiles of H-750TiO2@C at 0.05 A g-1. (b) TGA curve and (c) GCD profiles of at 0.05 A g-1 of H@C. (d) Rate performances of H-XTiO2@C at different current densities. Figure S7: Rate performance of H-750TiO2@C with mass loadings of (a) 1.61 mg cm-2 and (b) 1.86 mg cm-2 at different current densities. Figure S8: XRD patterns of the H-750TiO2@C electrodes after 3, 10, 100, and 2000 cycles. Figure S9: Nyquist plots of H-750TiO2@C before and at 2000th cycle at room temperature. Figure S10: CV curves of H-750TiO2@C measured at sweep rates between 1 and 10 mV S-1. Figure S11: In situ Raman spectra of H-750TiO2@C in the 50.0-800.0 cm-1 and the surface color mapping between 100.0 and 200.0 cm-1. Figure S12: (a) Ti 2p XPS spectra at different depths by argon ion sputtering of the electrode discharged to 0.01 V (the purple area represents the binding energy of some elements in the SEI layer at 450 eV) and (b) Ti3+: Ti4+ ratio as a function of argon ion sputtering etching depths. Figure S13: FESEM images and EDX mappings of H-750TiO2@C samples collected during the first GCD cycle discharged to 1.5 V (a), 0.6 V (b), 0.01 V (c), and charged to 0.8 V (d). Figure S14: GCD profiles tested in a half cell at 0.1 A g-1 of Na3V2 (PO4)3 (NVP). Figure S15: GCD profiles measured at 0.1 A g-1 (a) and rate performance (b) of a full cell with H-750TiO2@C as the anode and NVP as the cathode. Table S1: Impedance parameters obtained from equivalent circuit fittings. Table S2: Diffusion coefficient of Na +  at different cycles. Table S3: Comparison of the specific capacity and power density of full cell NVP//H-750TiO2@C with other full cells reported in the literature. Video of the in situ TEM test. The video shows the morphological changes of the H-750TiO2@C electrodes during the initial discharging process. (Supplementary Materials)

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