Porous TiO₂ nanoribbons and TiO₂ nanoribbon/carbon dot composites for enhanced Li-ion storage

Abstract

Porous TiO₂ nanoribbons and TiO₂ nanoribbon/carbon dot composites with excellent performances in lithium-ion batteries were prepared via a simple and efficient method.

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TiO₂ has attracted great attention as a promising anode material in Li-ion batteries (LIBs) due to its low cost, excellent rate capability, and superior safety.^1–3 However, the wide application of TiO₂ is still limited by its low electron transport ability, which always results in a poor cycling performance at high discharge–charge rates.^4,5 To solve this problem, preparing TiO₂ on a nanoscale has been proven to be an efficient way to reduce the pathways of the ions and electrons.^6,7 However, the nanomaterials often significantly self-aggregated owing to their high surface energy, which largely reduced the effective contact areas of the conductive additives and electrolytes in battery applications. In this way, keeping the effective contact areas as large as we can and fully realizing the advantages of active materials at the nanometer scale is still a challenge and of great importance. To date, considerable efforts have been made to explore TiO₂ with a variety of different nanostructures (e.g., sphere, porous, tube, wire, arrow, etc.)^8–11 and composite compositions (TiO₂/CNTs, TiO₂/graphene, TiO₂/Fe₂O₃, TiO₂/Sn, etc.), to resolve the problem of poor rate capability and the strong trend of aggregation.^12–15 Among these materials, the use of one dimensional (1D) TiO₂ nanostructures (rod, fiber, wire, tube, etc.) is a particularly efficient way to incorporate the advantageous characteristics of the material while achieving a desirable performance improvement.¹⁶ These 1D nanomaterials with a low degree of aggregation tendency could serve as anodes for LIBs in several aspects, as follows: (i) allowing a short Li-ion insertion–extraction distance; (ii) facilitating strain relaxation upon electrochemical cycling, (iii) providing a large surface-to-volume ratio for contact with the electrolyte, which can improve the capacity and cycle life of LIBs.¹⁷ Therefore, to some extent, such hierarchical 1D nanomaterials are one of the most favorable structures to use as anode or cathode materials for high-performance LIBs.

Moreover, 1D TiO₂ nanomaterials with high porosity, or which have been modified with carbon dots, can ensure a better electrochemical performance. The porous structures are not only conducive to lower current densities at the electrode–electrolyte interface,¹⁸ but are also beneficial for fully utilizing high electrochemical reaction rates per unit volume¹⁹ and enhancing diffusion kinetics, by reducing the diffusion pathway for electronic and ionic transport.²⁰ Alternatively, the doping/modifying/implanting of carbon dots into the matrix of the TiO₂ could result in a series of unique physical and chemical properties due to the intriguing characteristics of carbon dots,²¹ and could also significantly improve the electrochemical performance of the TiO₂, because they can narrow its band gap, reduce the polarization, enhance the conductivity, and increase the capacity stability.^22,23 Furthermore, the introduction of carbon can also induce the formation of Ti³⁺ sites in the TiO₂ crystalline structure, improving the conductivity and structure stability, which can efficiently enhance the rate capability and cycling performance of the TiO₂ electrode.^24,25 Most importantly, unlike the traditional coating of its surface with carbon, or loading it into the matrix of carbon, embedding carbon dots into the structure of TiO₂ could accelerate the transfer ability of the electrons from the internal body to the surface of the TiO₂.

Herein, we report a simple method to prepare 1D porous TiO₂ nanoribbons (P-TNRs) and TiO₂ nanoribbon/carbon dot composites (TNR/C) by just varying the calcination conditions. Compared to bulk TiO₂, they show a largely improved performance as anode materials in LIBs. For instance, P-TNRs, with a porous structure (∼8 nm) and high specific surface area (128 m² g⁻¹), exhibited a high capacity of around 250 mA h g⁻¹ at a current rate of 0.1 C, and a reversible capacity of over 138 mA h g⁻¹ after 100 cycles at 1 C (where 1 C = 168 mA h g⁻¹, TiO₂ + xLi⁺ + xe⁻ → Li_xTiO₂, x = 0.5). Meanwhile, the TNR/C that were embedded with plentiful carbon dots (1–3 nm) successfully delivered a higher capacity of around 300 mA h g⁻¹ at 0.1 C and a reversible capacity of over 225 mA h g⁻¹ after 100 cycles at 1 C. Furthermore, the resulting effects on the electrochemical properties of the TiO₂, caused by the embedding of the carbon dots and the introduction of porosity, were comparatively investigated and discussed.

In a typical synthesis, 4 g glucose and 1 g TiO₂ were completely dissolved in a 10 M NaOH solution, and then the solution was sealed in a Teflon-lined stainless steel autoclave and heated at 190 °C for 48 h. Finally, the as-prepared yellow-white intermediates were treated by HCl (pH = 1.0) and then calcined at 500 °C under a nitrogen atmosphere, giving rise to black TNR/C. P-TNRs could also be readily prepared from the intermediates, under an air atmosphere. According to the X-ray powder diffraction (XRD) patterns, the crystalline structures of the TNR/C and P-TNRs are both ascribed to normal anatase TiO₂ (JCPDS Card no. 89-4921) (Fig. 1). However, they show quite different colors, being black (TNR/C) and white (P-TNRs) respectively, due to the embedding of the carbon dots in the TNR/C.

Fig. 1 (a) XRD patterns and digital images (inset) of P-TNRs and TNR/C.

The intermediates have a uniform ribbon morphology and a good distribution in diameter, ranging from 100 to 300 nm, while their lengths can achieve several micro-meters, as shown in the scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images (Fig. 2a and b). The SEM images of the TNR/C and P-TNRs indicate that they can maintain the intermediates' morphology even after calcination, and also possess the characteristics of nanoribbons with a length of several micro-meters (Fig. S1†). The thickness of TNR/C is about 52 nm, as confirmed by the atomic force microscopy (AFM) image (Fig. S2†). This demonstrates that the nanoribbons have a high ratio of the length to width, which is consistent with the character of 1D nanomaterials. The details of the TNR/C nanostructures were also confirmed by TEM and HRTEM (Fig. 2c and d). The diameter of a single TNR/C nanoribbon is about 200 nm, and amorphous C-dots with a diameter of around 1–3 nm (marked by blue line) can be observed clearly in the structure of the TiO₂. It is well-known that carbon dots can be synthesized directly from glucose by a one-step alkali-assisted ultrasonic treatment, or by hydrothermal treatment. In this method, the carbon dots probably result from two approaches, hydrothermal treatment and the carbonization of residual glucose/intermediates of carbohydrates,²⁶ as confirmed by the Fourier transform infrared (FTIR) image (Fig. S3†). In the HRTEM image (Fig. 2d), the lattice fringe with a spacing of 0.35 nm occupies most of the areas on the TNR/C, and it corresponds well to the spacing of the (101) planes of TiO₂. The uniform distribution of the elements C, Ti, and O in the energy-dispersive X-ray spectrometry (EDS) mapping of single TNR/C unambiguously confirms that the carbon dots were indeed embedded in the structure of the TiO₂ nanoribbons (Fig. S4†). By changing the calcination conditions for the intermediate to an air atmosphere, P-TNRs with a porous structure could be obtained, as shown in Fig. 2e. The diameter of the P-TNRs is about 200 nm and the length is also in a micro-size distribution. The inter-planar distance is 0.35 nm, which belongs to the facet of the (101) planes of TiO₂. Notably, numerous pores with the diameter of 2–15 nm could be observed on the structure of the TNRs, directly confirming the porous characteristics of the P-TNRs.

Fig. 2 (a) SEM and (b) TEM images of the intermediates, (c) TEM and (d) HRTEM images of the TNR/C; (e) TEM and (f) HRTEM images of the P-TNRs.

The carbon content of the as-prepared P-TNRs and TNR/C was determined by thermogravimetric analysis (TGA), as shown in Fig. 3a. The sample of TNR/C shows a significant weight loss at 300 °C, which could be attributed to the combustion of carbon. Up to a temperature of 475 °C, the total weight loss of the TNR/C is 1.5 wt%, while it is 0 wt% for the P-TNR sample. It is clear that there is no residual carbon in the P-TNR sample, while the carbon dots in the TNR/C could be easily burned at a low temperature of 300–475 °C. In addition, the high-resolution Ti2p X-ray Photoelectron Spectroscopy (XPS) results for the TNR/C and P-TNRs are compared in Fig. 3b. Two broad peaks centered near 464.4 and 458.6 eV were well in accordance with the characteristic peaks of Ti2p_1/2 and Ti2p_3/2 of Ti⁴⁺, in both samples.²⁷ A small negative shift can be observed in the peak of Ti2p for the TNR/C sample, which demonstrates a change in the structure caused by the embedding of the carbon dots. By subtracting the normalized Ti2p spectra from those of the TNRs and the P-TNRs, two extra peaks appear at ∼462.0 and ∼457.6 eV, which are attributed to the Ti2p_1/2 and Ti2p_3/2 peaks of Ti³⁺.²⁸ This confirms that Ti³⁺ sites are successfully created in the TNR/C during calcination, under the reduction ability of glucose and the C-dots (high degree calcination at 500 °C introduces the following reaction: CO_xH_y + TiO₂ → TiO_2−n + CO_m). Moreover, the carbon dots formed in the structure of the TiO₂ can react with the TiO₂ to form lots of chemical bonds, such as Ti–O–C and Ti–C–O (Fig. S1†). This can reduce the redox potential of TiO₂ and introduce a lot of active centers, which are good for lithium ion storage.²⁹

Fig. 3 (a) TGA data of the as-prepared TNR/C and P-TNRs; (b) overlay of the normalized XPS high-resolution spectra of the Ti2p regions of the TNR/C and P-TNRs, respectively.

Fig. 4a shows the N₂ adsorption–desorption isotherm of the TNR/C sample. It might be categorized as a type III isotherm without a distinct hysteresis loop. This is conceivable because the conversion of residual glucose into carbon dots could cover and/or fill the porous structure of the TNRs. As a result, a relatively low surface area of 57 m² g⁻¹ was obtained. In contrast, the N₂ adsorption–desorption isotherms curves (Fig. 4a) of P-TNRs show a well-defined adsorption step and exhibit a typical type-IV isotherm, demonstrating the porous characteristics of the P-TNRs. The Brunauer–Emmett–Teller (BET) specific surface area and pore volume of the P-TNRs are 128 m² g⁻¹ and 0.5 cm³ g⁻¹, respectively. In contrast to the sample of non-porous TNR/C, the P-TNRs possess pores with a relatively narrow diameter distribution, ranging from 2–15 nm, as confirmed by the Barrett–Joyner–Halenda (BJH) pore size distribution. The tap densities of the TNR/C and P-TNRs are 0.66 and 0.60 g cm⁻³, respectively, which were evaluated by an Autotap device (Quantachrom).

Fig. 4 (a) Nitrogen adsorption–desorption isotherm and (b) BJH pore size distribution plot for the TNR/C and P-TNRs.

To circumstantiate the formative mechanism of the TNR/C and P-TNRs in the present reaction system, it could be interpreted as follows, based on the experimental route. At first, during the hydrothermal reaction process (190 °C, 24 h), a composite of Na₂Ti₃O₇/C-carbohydrate with a ribbon morphology could be obtained in the hydrothermal reaction of glucose and NaOH.¹⁶ Then, the as-prepared Na₂Ti₃O₇/C nanoribbons are treated by HCl (pH = 1.0) to form the intermediate H_xTiO_2−x/2/C-carbohydrate. With the calcination of these intermediates in an N₂ atmosphere, the H_xTiO_2−x/2 would be dehydrated and transform to TiO₂, at which point the carbon/residual carbohydrate could also be further carbonized to form highly conductive carbon dots. In this way, carbon dots are in situ encapsulated and/or embedded into the structure of the TNRs. Alternatively, upon calcination of the intermediates of H_xTiO_2−x/2/C in an air atmosphere, the composite could dehydrate and transform to TiO₂ in the same manner, but the carbon-carbohydrate in H_xTiO_2−x could also be burned, leaving numerous pores in the structure of the TNRs.

The electrochemical performances of the TNR/C and P-TNRs in LIBs have been evaluated in a half-cell test. The cyclic voltammograms (CV) of the TNR/C and P-TNRs from the 2nd cycle at a scan rate of 0.1 mV s⁻¹ in the voltage range of 0.01–3.0 V are shown in Fig. 5a. The sharp oxidation peak at 1.70 V in the anodic scan, and the reduction peak at about 2.0 V in the cathodic scan, are associated with the Ti⁴⁺/Ti³⁺ redox couple during lithium insertion and extraction.¹³ The anodic peak (2.04 V) of the TNR/C is more negative than that of the P-TNRs (2.09 V), indicating the lower polarization coefficient which the TNR/C possess.¹⁸ Moreover, the current density and the integrated area of the TNR/C are larger than that of P-TNRs, demonstrating that the carbon dots lead to a complete oxidation reaction during anodic cycling. Fig. 5b shows the 2nd rate charge–discharge curves for the TNR/C and P-TNRs electrodes between 1.0 and 3.0 V at 0.05 C. Both curves are in good agreement with the CV analysis. Each discharge (lithium insertion) can be divided into three stages: (i) a fast decrease in voltage from the open-circuit potential to ∼1.7 V; (ii) a fairly long plateau region at ∼1.7 V, reflecting the process of lithium insertion into the channels of the crystal structure; (iii) a long gradual decay of the voltage following the plateau, indicating the insertion of lithium ions into the microstructure of the material. Although the P-TNRs have a larger surface area than that of the TNR/C, a much longer plateau can be observed for the TNR/C, indicating a rich interface (carbon dots/TiO₂) existing in the TNR/C for lithium-ion storage. Due to the high degree of the redox reaction and the possible interfacial lithium storage (pseudo-capacitive effects),³⁰ a high discharge capacity of 318 mA h g⁻¹ can be delivered in the first discharge, followed by a charge capacity of 297 mA h g⁻¹ for the TNR/C sample, but these are only 269 and 264 mA h g⁻¹ for the P-TNRs. The TNR/C electrode also exhibits a smaller separation between charge and discharge voltage plateaus, demonstrating that the advanced structure possesses a lower electrochemical polarization and induces a better reversibility during the discharge–charge processes.³¹ A comparison of rate capabilities is shown in Fig. 5c. The TNR/C demonstrate a better performance than the P-TNRs and bulk TiO₂, delivering discharge capacities over 283, 250, 212, 146, 104 and 77 mA h g⁻¹ at rates of 0.1 C, 0.5 C, 1 C, 2 C, 5 C and 10 C respectively, and finally recovering to around 256 mA h g⁻¹ at 0.1 C. Fig. 5d shows the cyclability of the TNR/C, P-TNRs, and bulk TiO₂ electrodes at 1 and 10 C, respectively. The TNR/C electrode shows an excellent cyclic performance and a high reversible specific capacity of over 225 and 73 mA h g⁻¹ after 100 and 200 cycles, which is much higher than that of the P-TNR (∼138 and 45 mA h g⁻¹) and bulk TiO₂ (∼30 and 19 mA h g⁻¹) electrodes. The SEM images of the TNR/C and P-TNRs electrodes after 20 cycles at a rate of 0.1 C show that their morphology is similar to that before cycling, without any obvious aggregation. This demonstrates the high structural stability of the TNRs during the charge–discharge, owing to the small volume variation of the TiO₂, of less than 4%.^1–3 Simply compared to other TiO₂-based materials reported before, the TNR/C materials demonstrate a higher capacity at a rate of 1 C (Table S1†).

Fig. 5 (a) CV of the TNR/C and P-TNRs electrodes; (b) voltage profiles of the TNR/C and P-TNR electrodes at a rate of 0.05 C; (c) at different rates (0.1–10 C); (d) cycling performances at 1 and 10 C of the TNR/C, P-TNR, and bulk TiO₂ electrodes.

To interpret the results, herein we built several models of electronic transfer routes in a composite of C–TiO₂ to interpret and explain the high performance of the materials, notwithstanding the known advantages of 1D structures. An electrode consisting of active materials, a binder of PVDF, and Super P conductive material, with a mass ratio of 8 : 1 : 1, was cast on copper foil, as illustrated in Fig. 6a. Normal bulk TiO₂ is a well-known semiconductor with a low electronic conductivity (∼10⁻⁹ S cm⁻¹), and therefore the rate of internal electrons transferring from TiO₂ to the Super P&PVDF should be very slow (Fig. 6b). As an alternative, numerous traditional works focused on coating the TiO₂ with a layer of carbon. This could improve the electrochemical performance of TiO₂@C because the electrons could transfer to the walls of carbon in any directions (Fig. 6c). But the enhancement was limited because the transfer rate of electrons from the internal TiO₂ to the carbon walls was still low if the TiO₂ particle was too big. To shorten the transfer route of electrons in the structure of TiO₂, in this work we introduced a composite of carbon dots embedded into the TiO₂; that is, TNR/C. As shown in Fig. 6d, the electrons could transfer easily over a short distance via a network, which could undoubtedly make the internal electrons transfer quickly and conveniently to the Super P&PVDF.³² The increased electronic conductivity of 3.19 × 10⁻³ S cm⁻¹ was well in accordance with speculation. Although the transfer rate of electrons was enhanced, it would be better to further coat the TiO₂/carbon dots with a layer of carbon (Fig. 6e), because in this way the electrons could transfer from any direction to the walls of the carbon layer, which is another promising candidate for an anode. The disadvantage of porous TiO₂ nanoribbons could be easily observed from the transfer route of electrons (Fig. 6f). Numerous pores could inevitably be a barrier for the transfer of electrons, but the porous characteristics of the material could also enhance the contact areas with the Super P&PVDF (Fig. 6f), as well as the electrolyte. So, this is a conflict, and the balance between the porosity and the conductivity should be controlled. However, the electrode conductivity of the P-TNRs still showed a higher electronic conductivity (2.75 × 10⁻⁵ S cm⁻¹) than that of bulk TiO₂ (1.24 × 10⁻⁶ S cm⁻¹), which should be ascribed to the smaller size and specific structure of the nanoribbons. This could be interpreted to mean that the P-TNRs showed a higher capacity than the bulk TiO₂ but a lower capacity than that of the TNR/C. The higher conductivity of the TNR/C compared to the P-TNRs was also proved by the electrochemical impedance spectroscopy (EIS) Nyquist plots (Fig. S6†). Based on the analysis of the results, an ideal structure of ordered porous TiO₂/carbon dots@porous C with a one dimensional structure is presented (Fig. 6g), which should be good for the transferability of electrons via a network consisting of carbon dots, and should also adsorb enough electrolyte for the uptake of lithium ions. Research about such advanced materials deserve to be further undertaken.

Fig. 6 (a) Model of cast electrode on copper foil. Probable electronic transfer route in (b) TiO₂, (c) TiO₂@C, (d) TiO₂/carbon dots, (e) TiO₂/carbon dots@C, (f) porous TiO₂ and porous TiO₂/carbon dots@porous C.

In summary, a one-step strategy to prepare porous TiO₂ nanoribbons and TiO₂ nanoribbon/carbon dot composites with a narrow size distribution and uniform morphology was successfully developed. The probable mechanism of the formation of the porosity and composite structure were discussed. As an anode material for LIBs, P-TNRs with porous structures (∼8 nm) and high specific surface area (128 m² g⁻¹) are capable of a high capacity of around 250 mA h g⁻¹ at a current rate of 0.1 C and a reversible capacity of over 138 mA h g⁻¹ after 100 cycles at 1 C. In addition, TNR/C that are embedded with abundant carbon dots (1–3 nm) can deliver a higher capacity of around 300 mA h g⁻¹ at a rate of 0.1 C and a reversible capacity of over 225 mA h g⁻¹ after 100 cycles at 1 C. This dramatic improvement is attributed to the optimization of material design in terms of the structure and/or compositions, which can advance similar investigations into other kinds of metal oxide (e.g., FeO_x, CoO_x, NiO_x, MnO_x) as anode materials. Moreover, this kind of 1D porous TiO₂ nanoribbon and TiO₂ nanoribbon/carbon dot composite could also be widely applied in other areas of catalysis, electrochemistry, photo-electronics and materials science. The analysis of the introduced models of electronic transfer routes in the metal oxide–carbon composite could be expanded, and also be significant for designing new advanced materials.

Financial support from the Nature Science Foundation of China (no. 20873089, 20975073), the Nature Science Foundation of Jiangsu Province (no. BK2011272), the Industry-Academia Cooperation Innovation Fund Projects of Jiangsu Province (no. BY2011130), the Key Laboratory of Lithium Ion Battery Materials of Jiangsu Province, the China Scholarship Council (File. no. 201306920005) and the Graduate Research and Innovation Projects in Jiangsu Province (CXZZ13_0802) is gratefully acknowledged.

Notes and references

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Footnote

† Electronic supplementary information (ESI) available: Materials and characterization. SEM images of TNR/C, P-TNRs and their electrodes before and after cycling. AFM image of TNR/C. FTIR of TNR/C, P-TNRs and their intermediates. Comparative performances of P-TNRs, TNR/C and other TiO₂-based materials in previous literature. Nyquist plots of TNR/C and P-TNRs electrodes. See DOI: 10.1039/c4ra00133h

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