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Electrostatic Self-assembly Enabling Integrated Bulk and Interfacial Sodium Storage in 3D Titania-Graphene Hybrid Guiliang Xu, Lisong Xiao, Tian Sheng, Jianzhao Liu, Yi-Xin Hu, Tianyuan Ma, Rachid Amine, Yingying Xie, Xiaoyi Zhang, Yuzi Liu, Yang Ren, Chengjun Sun, Steve M. Heald, Jasmina Kovacevic, Yee Hwa Sehlleier, Christof Schulz, Wenjuan Liu Mattis, Shi-Gang Sun, Hartmut Wiggers, Zonghai Chen, and Khalil Amine Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b04193 • Publication Date (Web): 14 Dec 2017 Downloaded from http://pubs.acs.org on December 14, 2017

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Electrostatic Self-assembly Enabling Integrated Bulk and Interfacial Sodium Storage in 3D Titania-Graphene Hybrid Gui-Liang Xu,1,† Lisong Xiao,2,† Tian Sheng,3,† Jianzhao Liu,1 Yi-Xin Hu,1,4 Tianyuan Ma,1 Rachid Amine,5 Yingying Xie,1 Xiaoyi Zhang,6 Yuzi Liu,7 Yang Ren,6 ChengJun Sun,6 Steve M. Heald,6 Jasmina Kovacevic,2 Yee Hwa Sehlleier,2 Christof Schulz,2 Wenjuan Liu Mattis,8 Shi-Gang Sun,3 Hartmut Wiggers,2,* Zonghai Chen1,* and Khalil Amine1,* 1

Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700 S Cass Ave, Lemont, IL 60439, USA

2

Center for Nanointegration Duisburg-Essen (CENIDE), University of DuisburgEssen, Duisburg 47048, Germany 3

Collaborative Innovation Center of Chemistry for Energy Materials, State Key Laboratory Physical Chemistry of Solid Surfaces, Department of Chemistry, Xiamen University, Xiamen 361005, China 4 Department of Chemistry, University of North Carolina, Chaper Hill, NC 27599, USA 5 Materials Science Division, Argonne National Laboratory, 9700 S Cass Ave, Lemont, IL 60439, USA 6 X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, 9700 S Cass Ave, Lemont, IL 60439, USA 7 Nanoscience and Technology Division, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, IL 60439, USA 8 Microvast Power Solutions, 12603 Southwest Freeway, Stafford, TX 77477, USA

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Abstract Room temperature sodium-ion batteries have attracted increased attention for energy storage due to the natural abundance of sodium. However, it remains a huge challenge to develop versatile electrode materials with favorable properties, which requires smart structure design and good mechanistic understanding. Herein, we reported a general and scalable approach to synthesize 3D titania-graphene hybrid via electrostatic-interaction-induced

self-assembly.

Synchrotron

X-ray

probe,

transmission electron microscopy and computational modeling revealed that the strong interaction between Titania and graphene through comparably strong van-derWaals forces not only facilitates bulk Na+ intercalation but also enhances the interfacial sodium storage. As a result, the titania-graphene hybrid exhibits exceptional long-term cycle stability up to 5000 cycles, and ultrahigh rate capability up to 20 C for sodium storage. Furthermore, density function theory calculation indicated that the interfacial Li+, K+, Mg2+ and Al3+ storage can be enhanced as well. The proposed general strategy opens up new avenues to create versatile materials for advanced battery systems. Keywords: Sodium-ion batteries, Anode, Titania-graphene, Interfacial, and Density function theory


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Introduction Lithium-ion batteries (LIBs) have dominated the markets for powering consumer electronic devices (e.g., laptops) in the past decades and are now making in-roads into transportation because of their superior properties such as high energy density, long cycle life, lack of memory effect and environmental friendliness.1, 2 However, given the limited lithium resources in the Earth’s crust, the increasing demand for largescale electric energy storage such as grid-scale storage has raised concerns about the feasibility of the widespread use of Li-ion batteries.3 As a good and acceptable alternative to Li-ion technology, sodium-ion batteries (SIBs) have recently made a comeback because of the greater abundance and lower cost of sodium.4 In addition, multivalent batteries such as magnesium-ion batteries (MIBs)5 and aluminum-ion batteries (AIBs)6 have also been considered as alternatives to LIBs chemistries owing to their multivalency, low cost, elemental abundance, and environmentally friendly chemistry. However, as usual in any new chemistry, there have been setbacks during the development of high performance battery systems. With rising interest in green battery materials for rechargeable batteries, Ti-based anode materials such as Titania (TiO2), lithium titanate (Li4Ti5O12), Na4Ti5O12 and Na2Ti3O7 have recevied much attention in recent years because of their low cost and minimal environmental impact.7-10 In particular, TiO2 presents a moderate sodium insertion/extraction voltage (~ 0.7 V vs. Na/Na+), which not only efficiently avoids sodium plating on the anode but also provides suitable working voltage when coupled with a cathode material.11 Moreoever, compared to the alloy-based materials such as phosphorus,12 TiO2 undergoes smaller volume expanison during sodiation/desodiation, which could lead to better cycle stability. However, as a semiconductor material, the high rate behavior of Na storage with TiO2 is seriously restricted by its poor electrical conductivity and sluggish ion diffusion kinetics. Recent works have shown that interfacial Li or Na storage may occur when viable solid-liquid or solidsolid interfaces are designed in the battery system by constructing either porous structure or mixed phase composition.13-16 This interfacial Li or Na storage 3 ACS Paragon Plus Environment

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phenomenon could offer excess capacity, improved cycling life and rate capability. Therefore, rational structure design to enable enhanced interfacial sodium storage is an effective way toward high-performance SIBs. On the other hand, the particle size of nanocrystals are also crucial factors for improving cell performance because ultrafine nanoparticles not only endows the required conductivity to individual nanocrystals but also shortens the diffusion length for Na+. Besides, it can help to accommodate the mehcnical strain much better than that of the large particle materials during Na+ insertion/extraction process. All these factors are beneficial for high sodium storage and rate capability.17,

18

Thus,

integrating ultraﬁne nanoparticles into highly conductive carbon matrixes may be a promising approach to simultaneously achieve the outstanding rate capability and long cycle life. However, the main routes used for the fabrication of the TiO2-carbon (TiO2-C) hybrids are usually based on complicated synthetic steps, which may involve a high temperature or long-time hydrothermal/solvothermal process followed by calcination at comparably high temperature to remove the residual organics, leading to increased cost and aggravated scale up.19-22 As a result, there is considerable interest in developing a general and scalable approach to synthesize TiO2-C hybrids with favorable structures for high-performance rechargeable batteries. In addition, the sodium insertion/extraction mechanisms for TiO2-based anode materials for SIBs are not consistent, which will impede further improvement on the electrochemical performance of TiO2-based anode materials. For example, Kang and co-workers reported that the intercalation mechanism of sodium ions into the bulk structure of anatase TiO2 is dominant in the electrochemical Na+ insertion/desertion process of the N-doped open-pore channeled graphene-TiO2 composite.22 Yang and co-workers also found that graphene-TiO2 nanospheres go through a reversible Na+insertion mechanism without involving any conversion reactions.20 On the contrary, Chen and co-workers reported that intercalation pseudo-capacitance dominates the charge storage process of the graphene-coupled TiO2-B sandwich-like hybrid.21 Lu et al. recently reported high-performance Na+ capacitors using mesoporous single4 ACS Paragon Plus Environment

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crystal-like anatase TiO2-graphene nanocomposites, which is also based on pseudocapacitive sodium storage.23 Such debates clearly emphasize the significant benefit of unraveling the correlation between structures and sodium storage mechanisms, which can serve as a good guide to devise better anode materials for SIBs. In this work, we described a general and facile scaling-up approach to directly assemble ultrafine TiO2 nanocrystals (ca. 8 nm) onto conductive host by electrostaticinteraction-induced self-assembly. The as-synthesized TiO2 nanocrystals in our work present positive surface charges, which can be easily anchored onto graphene nanosheets. Such a structure can be expected to improve both the intra-particle electrical conduction and the Na+ diffusion. Furthermore, the ultrafine TiO2 nanocrystals attached to graphene reduces the mechanical stress within the electrode, leading to better structural integrity and thus a longer cycle life. When used as anode material of SIBs, the TiO2-graphene (TiO2-G) hybrid exhibits excellent long-term cycle stability and ultrahigh-rate capability, which can maintain overall reversible capacities of ca. 150 mAh g-1composite at 2 C in 1600 cycles and even 90 mAh g-1composite in 5000 cycles with no sign of capacity fading at extremely high rate of 10 C. High resolution transmission electron microscopy (HRTEM), X-ray absorption near-edge spectroscopy (XANES), in-situ high-energy X-ray diffraction (HEXRD) and density functional

theory

(DFT)

calculations are

used

to

elucidate

the

sodium

insertion/extraction mechanism of the TiO2-G hybrid, providing a foundation for understanding their outstanding performance. It was confirmed that the sodium storage mechanism of the TiO2-G hybrid is based on an integrated bulk and interfacial storage process. The strong interaction between TiO2 NCs and graphene nanosheets through electrostatic-interaction-induced self-assembly not only facilitates the intercalation of sodium into the bulk TiO2 crystal structure but also enhances the interfacial sodium storage. Moreover, DFT calculation indicated that such strong interaction could enhance interfacial Li+, K+, Mg2+ and Al3+ storage as well. It is expected that the scientific findings presented in this work will open up new avenues


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for a series of high performance electrode materials for efficient energy storage systems. Experimental section Synthesis of TiO2 Nanocrystals. TiO2 nanocrystals were synthesized using a sprayflame reactor as shown in Figure S1.24 In brief, titanium (IV) isopropoxide (97%, Sigma-Aldrich) was dissolved in isopropanol to form a 0.5 M precursor solution. The prepared solution was injected (2 mL min-1) by a syringe pump into the reactor through a burner nozzle that was mounted at the bottom of the reactor. Pure oxygen (Air Liquid, 99.95%) was supplied with 9 slm (standard liters per minute) surrounding the nozzle as the dispersion gas, which atomized the injected solution and generated a spray due to the shear forces. The spray was then ignited by a premixed methane/oxygen (1.5 slm/3 slm) pilot flame and supported by an oxygen sheath flow (4 slm). In addition, a nitrogen flow (300 slm) was added downstream to quench the off-gas as well as to suppress the particle growth. The TiO2 nanocrystals were formed during the combustion process and then collected downstream from the reactor on a baghouse filter. The obtained TiO2 nanocrystals were dispersed in deionized water to form a 4 mg ml-1 solution for the next step. Synthesis of graphene oxide (GO). Graphene oxide was synthesized by using natural graphite as the starting material via the modified Hummers’ method.25 Typically, 2.0 g of graphite and 1.0 g of NaNO3 were added into a 1.0 l flask that was seated in an ice bath. Then, 120 ml of concentrated H2SO4 and 5.0 g of KMnO4 were added slowly into the flask under magnetic stirring. The obtained mixture was continuously stirred for 2 h while it was kept in the ice bath. Subsequently, the mixture was heated up to 35 °C and stirred further for another 2 h. Afterwards, 250 ml of warm water (70 °C) was dropped into the mixture, which resulted in a brownish dispersion. At the end, 40 ml of H2O2 aqueous solution (30%) was added, and the color of the dispersion changed to brilliant yellow. The reactant was centrifuged and washed three times with HCl aqueous solution (5 wt.%), then washed another three times with warm water. Finally, the product was dried at 60°C for 24 h under the vacuum. The obtained GO 6 ACS Paragon Plus Environment

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powder was re-dispersed in deionized water to form a 4 mg ml-1 solution. In order to exfoliate the GO layers, strong sonication was applied to the GO suspension for 30 min. Synthesis of TiO2-graphene composites. TiO2-graphene hybrids were prepared via an electrostatic-interaction-induced self-assembly process coupled with a lowtemperature reduction process (reduction of GO to graphene). Zeta potential (ζ) measurements indicated that the ζ values of the TiO2 and GO aqueous suspensions are +40.8 mV and -37.6 mV, respectively. The opposite surface charges of these two materials provided an electrostatic interaction between them, which acted as the driving force for the self-assembling. In the typical synthesis of TiO2-G, aqueous dispersions of TiO2 and GO were separately sonicated for 15 min. After that, the TiO2 dispersion was added into the GO dispersion under magnetic stirring (with TiO2:GO weight ratio of 2:1), and the obtained mixture was further sonicated for another 15 min. Subsequently, ascorbic acid as the reducing agent with a mass ratio of 4:1 to GO was added into the TiO2-GO suspension under stirring. The dispersion was then transferred into glassware and kept in the oil bath at 70 °C for 4 h. After this process, the dark hydrogel that was formed was collected and washed three times with water. The obtained composites were first dried in a freeze dryer (BETA 2-4 LD plus LT, Martin Christ Gefriertrocknungsanlagen GmbH, Germany), and then further dried in a vacuum oven at 120 °C for 12 h. The synthesis of TiO2-G* are similar to the process for TiO2-G except the ratio for TiO2 and GO are 3:1. Structure Characterization. Ex-situ and in-situ high-energy synchrotron XRD experiments were carried out at beamline 11-ID-C and 11-ID-D of the Advanced Photon Source (APS) at Argonne National Laboratory using X-rays with wavelengths of 0.1173 Å and 0.8 Å, respectively. The Ti K-edge measurements for the TiO2-G hybrid electrode at different charge/discharge states were recorded at Beamline 20BM-B of APS. Raman experiments were performed using a Renishaw inVia microscope spectrometer. The morphologies and structures of the materials were characterized by field emission scanning electron microscopy (HITACHI S-4700-II or 7 ACS Paragon Plus Environment

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JEOL JSM-7500) and transmission electron microscopy (JEOL JEM-2100F). Energy dispersive X-ray spectroscopy (EDX) was conducted with an EDX detector in the JEOL JEM-2100F. The Brunauer-Emmett-Teller (BET) measurements of the composite materials were carried out at 77 K on a Quantachrome NOVA2200 analyzer. The mass loading of TiO2 nanocrystals in the TiO2-G hybrid was determined by thermogravimetric (TG) analysis, which was performed on the Netzsch 449 F1 Jupiter (with a heating rate of 10 K min-1 under a synthetic air flow at 25 ml min-1. The zeta potential measurements of TiO2 and GO aqueous dispersions were carried out on a Zetasizer (Malvern Instruments). Electrochemical Characterization. The TiO2, graphene, TiO2-G hybrid, TiO2-G* hybrid and TiO2+G mixture (2/1 in weight ratio) electrodes were prepared by spreading a mixture of 70 wt.% active material, 20 wt.% super-P, and 10 wt.% sodium aliginate (2 wt.% in H2O) onto a copper foil current collector. The as-prepared electrodes were then dried at 100 °C in a vacuum oven for 12 h. The active material mass loading is controlled at about 1.5-1.8 mg cm-2, which is comparable with most of Ti-based anode materials for sodium-ion batteries.7-30 The charge/discharge rate of 1 C equals 200 mA g-1. The electrochemical performance of the as-prepared electrodes were characterized by assembling them as anode in coin cells (type CR2032) in an argon-filled glove box under conditions where the contents of moisture and oxygen were both below 0.5 ppm. The electrode was separated from the sodium counter electrode by a separator (glass fiber, Grade GF/F Glass Microfiber Filter Binder Free, circle, 125 mm). The electrolyte used in the cell was 1 M NaPF6 in propylene carbonate with 2 vol.% fluoroethylene carbonate as additive. The cells were charged and discharged using a MACCOR cycler. Cyclic voltammograms of the TiO2-G hybrid electrode were recorded on a Solartron Analytical 1470 System between 0.01 and 3.0 V (vs. Na/Na+) at different scan rates. Electrochemical impedance spectroscopy of the graphene electrode, TiO2-G hybrid and TiO2 electrode were recorded on a Solartron Analytical 1470 System in a frequency range of 100


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kHz to 0.1 Hz. The electrochemical reaction kinetic current of the TiO2-G hybrid electrode was measured in a home-build high-precision source meter (Keithley 2401). Computational details. The spin-polarized electronic structure calculations were carried out by using the Vienna Ab-initio Simulation Package (VASP) with PerdewBurke-Ernzerhof (PBE) functional of exchange correlation. The projector-augmentedwave (PAW) pseudopotentials were utilized to describe the core electron interaction.26, 27 The on-site Hubbard U term (DFT+U) was added for O 2p orbitals at the value of 6.3 eV and for Ti 3d orbitals at the value of 4.2 eV. The cut-off energy was 450 eV, and a 2×1×1 Monkhorst-Pack k-point sampling was used. The Na coverage was defined as θ = nNa/nTi, where nNa is the number of Na atoms adsorbed on the surface and nTi is the number of Ti atoms in the surface layer, which is sixteen. The adsorption energy was defined as Eb = (ENa/surf –Esurf)/n - ENa, where ENa/surf and Esurf are the total energies of the surface with and without n Na atoms adsorbed, respectively, and ENa is the energy of one Na atom in the bulk position. Results and Discussion The formation mechanism of the TiO2-G hybrid is illustrated in Figure 1a, which can be seen as an “electrostatic-interaction-induced self-assembling” process. The ultrafine TiO2 nanocrystals (NCs) were synthesized in a spray-flame reactor (see Figure S1 and Experimental for details), which directly possess positive surface charges

as-synthesized.

Therefore,

unlike

the

commonly

reported

surface

modification to render nanocrystals a positively charged surface,28-31 our method works without using any additional surfactants or chemical linkers. The graphene oxide (GO) nanosheets were prepared from natural graphite via the well-known Hummer’s method25 with the formation of a high amount of hydroxyl and carboxyl surface function groups, which can act as the anchoring sites for the positively charged TiO2 NCs.32 The zeta (ζ) potentials of TiO2 NCs and the GO nanosheets suspensions in water were measured to be +40.8 and -37.6 mV, respectively (Figure S2), indicating that TiO2 NCs and GO nanosheets were oppositely charged. The electrostatic attraction between these two oppositely charged components provides a 9 ACS Paragon Plus Environment

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driving force for the self-assembling process and is also responsible for the strong interactions between GO nanosheets and TiO2 NCs. Once being in touch with each other, comparably strong van-der-Waals forces induce the self-assembly of TiO2 NCs onto GO nanosheets. Thus, a colloidal suspension was observed directly after mixing the GO and TiO2 dispersions (Figure 1b). After heating at 70 oC for 4 h, the ascorbic acid can reduce GO to graphene and simultaneously crosslink the graphene nanosheets with volumetric shrinkage, giving rise to the formation of threedimensional (3D) TiO2-G hydrogel (Figure 1b). Compared to the previously reported process for the fabrication of TiO2-graphene nanocomposites,19-23 our method is simplified and can be easily advanced to production scale required for battery material generation. Figure 1c further illustrates the proposed sodium storage mechanism of the 3D TiO2-G hybrid. Firstly, the graphene can facilitate the electron transfer and Na+ transportation, which is crucial for high rate charge/discharge.33 Secondly, the strong interaction between TiO2 NCs and graphene nanosheets results in enhanced interfacial sodium storage at the TiO2/G interface, which plays a critical role on its remarkable long-term cycle stability and will be further confirmed by DFT calculation. In addition, the intercalation of Na+ into the bulk TiO2 was also enhanced, as clearly evidenced by HRTEM and synchrotron HEXRD. Therefore, we propose that the TiO2-G hybrid fabricated by the electrostatic self-assembly is based on an integrated bulk and interfacial sodium storage process. HEXRD was employed to determine the crystallographic phases of graphene nanosheets, TiO2 NCs and 3D TiO2-G hybrid. As shown in Figure 2a, graphene nanosheets present amorphous structure, while the HEXRD peaks of TiO2 NCs can be indexed to tetragonal anatase (JCPDF no. 21-1272) and tetragonal rutile (JCPDF no. 21-1276). To quantify the amount of anatase and rutile in the as-prepared TiO2 NCs, Rietveld refinement was conducted under a two-phase model (Figure S3). The results indicate that the majority of TiO2 NCs is anatase, which has a mass percentage of 85%. The HEXRD pattern of TiO2-G hybrid is almost the same as that of TiO2 NCs, indicating that negligible structure changes of TiO2 NCs occurred during the self10 ACS Paragon Plus Environment

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assembly process. The structures of TiO2 NCs and TiO2-G hybrid were further characterized by Raman spectra (Figure S4). The Raman spectrum of TiO2 NCs shows five peaks (at 143 cm-1, 196 cm-1, 396 cm-1, 514 cm-1, and 637 cm-1), which are assigned to the vibrations of the crystalline lattice and match the Raman shifts attributed to the Eg, Eg, B1g, A1g and Eg phonon modes, respectively.34 These peaks, however, cannot be observed in the case of the TiO2-G hybrid, suggesting that TiO2 NCs are very small and homogeneously distributed in the TiO2-G hybrid.35 In addition, two broad peaks are centered at 1323 cm-1 and 1604 cm-1, which are attributed to the D and G bands of the graphene, respectively. The Brunauer-EmmettTeller (BET) results show that the TiO2 NCs, graphene nanosheets and TiO2-G hybrid have a specific surface area of 183, 104, and 203 m2 g-1, respectively (Figure 2b). The increased specific surface area may be due to a combination of TiO2 NCs and graphene nanosheets. Assuming the TiO2 NCs are non-porous and spherical, the average particle size of TiO2 NCs can be estimated to be around 8 nm according to the equation below.36 D=

(Equation 1)

where SSA is the specific surface area measured by BET,

is the density of the

particle (4.26 g cm-3), and D is the particle diameter. The content of TiO2 NCs in the TiO2-G hybrid was further quantified to be around 66.4 wt.% based on thermogravimetric analysis (Figure S5). The microstructures of graphene nanosheets and TiO2-G hybrid were further characterized by scanning electron microscopy (SEM). Figure S6a and S6b show the low- and high-magnification SEM images of graphene, in which well exfoliated nanosheets can be clearly seen. After the attachment of TiO2 NCs, the surface of the TiO2-G hybrid became relatively rough (Figure S6c), indicating that the TiO2 NCs are homogeneously distributed on the graphene nanosheets. Also discovered was that graphene nanosheets connected with each other to form 3D interconnected networks with macropores (Figure S6d). This was attributed to the hydrophobic and π-π stacking interactions of the graphene layers. Closer observations on the morphology 11 ACS Paragon Plus Environment

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and structure of TiO2 NCs and TiO2-G hybrid were obtained by TEM. The TEM image in Figure 2c shows that the TiO2 NCs have an average particle size of ca. 8.1 nm (see particle size histogram in Figure S7), which coincided with the BET result. The inset HRTEM image showed clear lattice fringes with interplanar distance of 0.354 nm, which are in good agreement with the d-spacing value of the (101) plane for anatase TiO2. Figure 2d and 2e show the low and high magnification TEM images of TiO2-G hybrid, which reveal exfoliated graphene nanosheets covered with TiO2 NCs. The scanning transmission electron microscopy (STEM) image and the corresponding elemental mapping images clearly reveal the homogeneous distribution of Ti and C elements in the TiO2-G hybrid (Figure 2f-2h). Figure S8 further shows the XPS spectra of C 1s and O 1s of TiO2-G hybrid. The result revealed that there are only C-C, C-O, C=O, and O-Ti4+ bonds, but without the formation of Ti-C bond, indicating that TiO2-G hybrid was prepared mainly through electrostatic selfassembly due to opposite surface charges. Titania NCs directly assembled on graphene appeared to exhibit strong interactions with the underlying graphene nanosheets since sonication did not lead to their dissociation from the sheets. This strong coupling would lead to significantly improved sodium storage performance as discussed below. The electrochemical performances of the graphene, TiO2, TiO2+G mixture and TiO2-G hybrid were evaluated by assembling them as anodes into coin cells with sodium as reference and counter electrode. Noted that the capacities and charge/discharge rates of the TiO2+G mixture and the TiO2-G hybrid in the present study were not calculated based on the mass of TiO2 only, but based on the overall mass of the TiO2+G mixture and TiO2-G hybrid. Figure 3a compares the cycling performance of graphene, TiO2 NCs, TiO2+G mixture and TiO2-G hybrid at a charge/discharge rate of C/10. As shown, the TiO2 NCs can deliver an initial reversible capacity of only 157.8 mAh g-1 and decreased to 134 mAh g-1 after 50 cycles. These electrochemical capacities are not due to Na+ intercalation but to storage of Na+ near the crystalline surface instead of.37 This is because bulk Na+ 12 ACS Paragon Plus Environment

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diffusion suffers from a high intrinsic activation barrier for the Na+ diffusion through the TiO2 anatase lattice as predicted by nudged elastic band calculations.37 The commonly observed broadening and decreased intensities of the TiO2 reflections in the reported TiO2-graphene nanocomposites during discharge are actually attributed to the partial loss of crystallinity of the outer shells of the TiO2 particles, while the direct formation of NaxTiO2 was not observed in the previously reported TiO2-C nanocomposites.19-23,

38-40

Intriguingly, of the test anodes, the graphene nanosheets

exhibited the highest initial reversible capacity of ca. 250 mAh g-1. The sodium storage mechanism of graphene should be envisaged as the adsorption of sodium on the graphene sheets or active sites, i.e. a surface storage mechanism.41 However, the graphene electrode undergoes a continuous capacity fading, and the capacity retention after 50 cycles was only 60%. This may be attributed to the continuous loss of sodium adsorption active sites during repeated charge/discharge. If there is no synergistic effect between TiO2 NCs and graphene nanosheets, the reversible capacity of the TiO2-G hybrid can be contributed to the overall contribution of individual TiO2 and graphene, i.e. similar to that of TiO2+G mixture. However, it was exciting to find the TiO2-G hybrid can still maintain a reversible capacity of around 200 mAh g-1composite after 50 cycles, showing much higher reversible capacities and better cycle stability than the TiO2+G mixture as well as individual TiO2 and graphene. This finding indicates that the improved electrochemical performance of the TiO2-G hybrid should arises from the synergistic interaction between TiO2 NCs and the underlying graphene nanosheets. More strikingly, with respect to the graphene, TiO2 NCs and TiO2+G mixture, the specific capacities of the TiO2-G hybrid are substantially increased at all investigated charge/discharge rates from C/10 to 20 C, as demonstrated in Figure 3b. The TiO2-G hybrid exhibits average overall reversible capacities of 237, 215, 196, 184, 175, 168, 148, 121, and 98 mAh g-1composite at charge/discharge rates of 0.1, 0.2, 0.4, 0.6, 0.8, 1, 2, 5, and 10 C, respectively. Even at an extremely high rate of 20 C, the overall reversible capacity is still retained at 72 mAh g-1composite. For comparison, the TiO2 13 ACS Paragon Plus Environment

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NCs electrode presented poorer rate performance with capacities of only 150, 129, 113, 106, 100, 97, 86, 69, 53 and 31 mA h g-1 at 0.1, 0.2, 0.4, 0.6, 0.8, 1, 2, 5, 10, and 20 C, respectively. The rate performance of the graphene electrode is unsatisfactory as well, demonstrated by capacities of 219, 175, 151, 138, 125, 110, 90, 75, and 55 mA h g-1 at 0.1, 0.2, 0.4, 0.6, 0.8, 1, 2, 5, 10, and 20 C, respectively. The finding that the rate performance of the TiO2+G mixture is obviously inferior to that of the TiO2-G hybrid, further confirms the synergistic interaction between TiO2 and graphene. The excellent high-rate performance of the TiO2-G hybrid is also partially attributed to the substantial decrease in charge-transfer resistance due to the unique structure of the TiO2-G hybrid. Figure S9 compares the Nyquist plots of the TiO2-G hybrid and TiO2 NCs electrodes. Apparently, the TiO2-G hybrid electrode shows a much lower electrochemical charge transfer resistance than the TiO2 NCs electrode (174.6 Ω vs. 741.8 Ω) due to the high electronic conductivity of graphene nanosheets (112.7 Ω). However, once amount of TiO2 NCs in the hybrid was too high (>70 wt.%) and form agglomeration, the reversibility and rate performance would be decreased (Figure S10). We should also note that the initial Coulombic efficiency of the TiO2-G hybrid at low rate (e.g. C/10) is not very high (Figure S11), which may be ascribed to the solid-electrolyte interphase (SEI) formation on the surface of the electrode material during the first cycle and the irreversible trapping of sodium ions at active sites of the graphene matrix as well as sluggish Na+ intercalation kinetics of TiO2.42-45 Strategies such as pre-sodiation,46 advanced binder,21 or lithium ethoxide treatment47 have been reported to efficiently mitigate the irreversible capacity. Another impressive feature of the TiO2-G hybrid is its exceptional long-term cycle stability. The TiO2-G hybrid can deliver reversible capacity of ca. 150 mAh g-1composite in 1600 cycles at a high charge/discharge rate of 2 C (Figure 3c). Unlikely the previous reported TiO2-based anode materials with floating reversible capacities during cycling,7, 21, 22, 48, 49 the ultra-stable capacity of TiO2-G hybrid makes it easy to couple with the cathode materials in the full cell due to the relatively constant negative/positive ratio. Also note that the Coulombic efficiency increased to almost 14 ACS Paragon Plus Environment

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100% after the first cycle, indicating a highly reversible sodiation/de-sodiation process of the TiO2-G hybrid in the subsequent cycles. For comparison, the TiO2+G mixture can only deliver ca. 90 mAh g-1 after 700 cycles at the same rate, confirming the synergistic sodium storage between TiO2 NCs and graphene nanosheets in TiO2-G hybrid. Figure 3d shows the charge/discharge voltage curves of the TiO2-G hybrid for different cycles at 2 C, which possesses an average discharge and charge voltage plateau of 0.65 V and 0.84 V, respectively (see dQ/dV in Figure S12). The stable voltage profiles and the small voltage hysteresis of only 190 mV upon cycling at 2 C reveal the excellent structure stability and ultrafast sodiation/de-sodiation kinetics of the TiO2-G hybrid. Figure 3e further illustrates the cycle stability of TiO2-G hybrid under ultrafast charge/discharge. As shown, it can maintain an ultra-stable capacity of ca. 90 mAh g-1composite over 5000 cycles at an extremely high rate of 10 C, demonstrating extraordinary long-term cycle stability. In general, the TiO2-G hybrid reported in this work demonstrates better sodium storage performance than most state-of-the-art TiO2-based anode materials for SIBs reported so far (see Table S1). The ultrahigh rate capability, together with ultralong cycle stability and the moderate working voltage of the TiO2-G hybrid makes it a very promising anode material for high-power SIBs. The sodium storage mechanism of TiO2-based anode materials is not well understood and still under debate.44,

50

In order to understand the outstanding

electrochemical performance of the TiO2-G hybrid and clearly unravel its sodium storage mechanism to devise better TiO2 anodes, we carried out in-situ and ex-situ characterizations. First of all, a home-built system for high-precision leakage current measurement was used to track the rate of the parasitic reaction between the sodiated electrode and the electrolyte (Figure S13a). The measured leakage current (i) is proportional to the reaction rate of parasitic reactions between the sodiated electrode and the electrolyte.12,

51-53

Hence, the static leakage current can be used as a

quantitative indicator for the reactions between different intermediates and the electrolytes. Figure S13b shows a typical current relaxation curve. Theoretically, the 15 ACS Paragon Plus Environment

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expected dependence of the leakage current on the potential will follow the Tafel equation (Eq. 2), which is an exponential decay function predicting a monotonically increasing leakage current with decreasing potential. (Equation 2) In the above equation, constant potential E,

represents the static leakage current measured at a

is the exchange current density, which is a kinetic constant for

a specific electrochemical reaction, and is the characteristic redox potential of the electrochemical reaction on the active electrochemical surface.12 Therefore, if there is only one kind of parasitic reaction between the sodiated electrode and the electrolyte, the value of Ln |i| should linearly increase along with decreasing potential. Figure 4a shows the Ln |i|-voltage plot for the sodiated TiO2-G hybrid electrode. As shown, it presents three segments (1.5-1.0 V, 1.0-0.4 V & 0.4-0.01 V) with two critical potentials centered at about 1.0 and 0.4 V. The data for each segment can be linearly fit, implying that at least two intermediates were involved during the sodiation of the TiO2-G hybrid. Cyclic voltammetry was further used to study the electrochemical behaviour of the TiO2-G hybrid (Figure S14a). The relationship between the peak current density ( ) and the sweep rate ( ) is shown in Eq. 3, (Equation 3) in which the b-value parameter is the slope of the log(i)-log(v) plot. A slope of 0.5 (b=0.5) implies an electrochemically diffusion-controlled process, i.e. battery behaviour, while an increase of the slope to 1.0 (b=1.0) suggests electrochemical surface Faradic redox reactions, i.e. capacitance behaviour.54, 55 The log(i)-log(v) plot for the TiO2-G hybrid electrode is depicted in Figure S14b, giving a slope of 0.642 (anodic) and 0.737 (cathodic) over the scan rate 0.2-100 mV s-1, indicating that the performance of sodium storage in the TiO2-G hybrid electrode can be deemed as overall contributions from both battery and capacitance behaviour,45 i.e. an integrated bulk and interfacial sodium storage process.


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To obtain the evidence for the intercalation of Na+ into the bulk TiO2 crystal structure (battery behaviour), we performed HRTEM on a discharged electrode. Figure 4b shows the HRTEM image of a deep-discharged TiO2-G hybrid electrode (0.01 V). Unlike the previously reported TiO2-graphene anodes,20,

22, 23

direct

intercalation of Na+ into the bulk structure of anatase TiO2 was observed for the TiO2G hybrid. As shown, the d-spacing value marked by the while circle was measured to be ca. 0.635 nm, which is much larger than the largest d-spacing value (0.3-0.4 nm) for anatase and rutile TiO2. This may correspond to the formation of a Na-rich phase—NaxTiO2, which is similar to that recently reported by Komaba and coworkers.50 In order to further unravel the structure evolution of the TiO2-G hybrid electrode during cycling, X-ray absorption spectroscopy was used to monitor the oxidation state variation of Ti during charge/discharge. Figure 4c and 4d show the Ti K-edge XANES spectra of TiO2-G hybrid electrode at different charge/discharge states. As shown, the pristine Ti presents a typical XANES spectrum of anatase TiO2, in which Ti is tetravalent. When the cell was discharged to 0.7 V, negligible structure changes (e.g. intensity and pre-edge position) were observed, indicating that the capacities gained from open-circuit voltage to 0.7 V are not due to the electrochemical reduction of TiO2, which may mainly come from the formation of the SEI layer and surface Na storage. When the cell was further discharged to 0.01 V, an energy shift towards lower energies was clearly seen, and the pre-edge feature became more pronounced (Figure 4d). This effect is related to the reduction of Ti4+ to Ti3+ during the insertion of Na+. However, it is totally different from that of Ti foil, demonstrating that no conversion reaction (Ti4+ to Ti0) is involved during the sodiation of the TiO2-G hybrid electrode. When the cell was recharged to 3.0 V, the Ti K-edge shifted to higher energies but did not resemble that of the pristine electrode, suggesting that Ti3+ is not fully oxidized back to Ti4+. To study the formation of NaxTiO2, we further carried out in situ synchrotron HEXRD to track the phase transformation of the TiO2-G hybrid during 17 ACS Paragon Plus Environment

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charge/discharge. Figure 4e shows selected in situ synchrotron HEXRD patterns of TiO2-G hybrid in the first charge/discharge cycle. Figure 4f shows the corresponding integrated area intensity for the anatase (101) reflection at different charge/discharge states. As shown, there are no obvious structure changes or integrated area intensity changes from Point A (1.43 V) to Point C (0.93 V), indicating that most of the gained capacities should come from the decomposition of the electrolytes to form the SEI layer. From Point C (0.93 V) to point E (0.40 V), the integrated area intensity was slightly decreased but no new peaks emerged. This behaviour corresponds to the interfacial Na storage process of the TiO2-G hybrid, leading to partial loss of crystallinity of the outer shells of the TiO2 particles and thus slightly decreased diffraction intensities. Starting at point E (0.40 V), a clear diffraction peak marked by diamond emerged accompanied by a significantly decreased integrated area intensity of anatase (101) reflection, which probably correspond to the intercalation of Na+ into bulk TiO2 to form a Na-rich phase—NaxTiO2. The critical potential is in agreement with the result indicated from the leakage current measurement. During the charge process, the diffraction peak of NaxTiO2 shifted to a lower angle and the integrated area of anatase (101) reflection gradually increased, which is attributed to the extraction of Na+ from NaxTiO2. However, the integrated area and the diffraction pattern did not recover to the pristine electrode because Na-poor phase such as Na0.23TiO2 may dominate the charged product.44 Despite this problem, owing to the strong interaction between TiO2 NCs and the underlying graphene nanosheets, a direct intercalation of Na+ into the bulk TiO2 is evident in the TiO2-G hybrid, which is different from the previously reported TiO2-graphene nanocomposites.19-23 As indicated by the electrochemical results, the TiO2-G hybrid exhibits much better cycle stability and rate capability than individual TiO2 NCs and graphene nanosheets; a synergistic interaction between graphene nanosheets and TiO2 NCs at the TiO2/graphene interface is believed to occur. In order to understand this effect on the enhanced sodium storage performance of TiO2-G hybrid, we further performed DFT calculations to investigate the Na adsorption process on the TiO2 (101) surface 18 ACS Paragon Plus Environment

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and at the TiO2/graphene interface at atomic levels. Figure 5a-5c show the models of graphene, TiO2 (101) phase, and TiO2/graphene interface, respectively. The graphene was modeled by 60 carbon atoms. The TiO2 anatase (101) surface was modeled by p(2×3) periodic slab supercell with three TiO6 layers, including 48 Ti and 96 O atoms. The bottom TiO6 layer was fixed, and other atoms were relaxed during geometry optimization. One graphene layer with 60 C atoms was epitaxially placed above the surface to model the TiO2/graphene interface. The lattice of graphene is changed by 1.4-2.1% to match the surface lattice of TiO2 (101). Figure 5d and 5e show side views of the most stable models for the TiO2 (101) surface and TiO2/graphene interface with an increase of Na coverage from 0.125 to 1.0 monolayer (ML). The adsorption energy is defined as Eb = (ENa/surf –Esurf)/n - ENa, where ENa/surf and Esurf are the total energies of the surface with and without n Na atoms adsorbed, respectively, and ENa is the energy of one Na atom in the bulk position. A more negative Eb indicates a more stable adsorption. As shown in Figure 5f, the binding energy decreases with an increase of Na coverage, and in general, the TiO2/graphene interface is able to provide higher binding energies than the bare TiO2 surface. For one Na atom where the most stable adsorption site is between two O atoms on the surface, the binding energies are -1.31 eV on TiO2 and -1.45 eV at the TiO2/graphene interface, respectively. For one monolayer coverage, the binding energy per Na atom is decreased to -0.19 eV on TiO2 and -0.48 eV at the TiO2/graphene interface (see Table S2 for detailed variation of binding energy with increase of Na coverage). The values of binding energies get closer to 0 eV, indicating that it is approaching the maximum Na accommodation. Furthermore, Bader charge analysis was carried out to quantify the charge distribution on the surface and identify the oxidation state of absorbed Na atoms.56 It can be seen from Figure 5g that at low coverage, Na atoms present ion states (Na+) with an average positive charge of ~0.9 e on the bare TiO2 surface and at the TiO2/graphene interface, that is, the charge separation occurs simultaneously when Na atoms were introduced on the surface. When the coverage increases to 1 ML, Na atoms at the TiO2/graphene interface retain cations well, with a positive charge of 0.85 e. However, 19 ACS Paragon Plus Environment

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Na atoms on bare TiO2 start to present somewhat metallic states, especially when the coverage reaches 0.75 ML. Finally, the positive charge per Na atom for bare TiO2 is only 0.62 e, much lower than the value of 0.85 e for the TiO2/graphene interface, indicating lower accommodation of Na+ during sodiation (see Table S2 for details). Moreover, the graphene layer can act as a second electron acceptor besides the TiO2 surface for accepting the negative charges that originally stored in Na atoms. Figure 5h shows that, at the 0.75 ML coverage, TiO2 obtain little negative charge from Na atoms. However, in the presence of a graphene layer, about ~28% extra negative charges could be effectively transferred to graphene (see Table S3& S4). As a result, the strong interaction between TiO2 and graphene can readily maintain the Na atoms as Na+, leading to a more favored Na adsorption during sodiation/de-sodiation and thus enhanced sodium storage performance. The strong interaction between TiO2 NCs and the underlying graphene nanosheets has led to enhanced interfacial sodium storage through electrostaticinteraction-induced self-assembly. Recently, Dambournet et al. have recently reported that by introducing a large number of titanium vacancies through aliovalent doping in TiO2, it can achieve a reversible Mg2+ and Al3+ insertion in anatase TiO2.57 The above finding has encouraged us to further explore its capability to store other monovalent and multivalent ions for potential use in other rechargeable batteries. In this work, DFT calculation was further used to study the binding energies of Li+, K+, Mg2+, and Al3+ on the TiO2 (101) surface and at the TiO2/graphene interface at atomic levels. As shown in Figure 6a, the bare TiO2 (101) surface gives binding energies of -1.29 eV (Li+), -0.77 eV (K+), -1.05 eV (Mg2+), and -0.22 eV (Al3+). With the presence of the graphene layer (Figure 6b), the binding energies were increased to -1.43 eV (Li+), 1.31 eV (K+), -1.38 eV (Mg2+), and -0.51 eV (Al3+). The enhanced binding strength for different monovalent and multivalent ions at the TiO2/graphene interface is believed to improve their electrochemical performance. Furthermore, the synthetic method reported in this work can be extended to synthesize oxide (e.g. V2O5 and ZnO), sulfide (e.g. FeS2), and phosphate (e.g. FePO4) nanocrystals (Figure S15) as 20 ACS Paragon Plus Environment

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well as other promising energy materials with tunable surface charges (see their corresponding zeta potential profile in Figure S16). As a result, these nanocrystals can be easily assembled on the conductive host with opposite charges. Conclusions In summary, a general and scalable approach based on electrostatic-interactioninduced self-assembly has been developed to synthesize 3D TiO2-graphene hybrid with homogeneous distribution. As a potential anode material for SIBs, the typical TiO2-graphene hybrid demonstrates exceptional long-term cycle stability up to 5000 cycles and ultrahigh rate capability up to 20 C. We further employed high-resolution transmission electron microscopy, operando X-ray diffraction and X-ray absorption spectroscopy to study the structure evolution and oxidation state changes during electrochemical processes. The results confirmed the intercalation of Na+ into the bulk TiO2 crystal structure and the interfacial sodium storage are both enhanced through the strong interaction between TiO2 nanocrystals and the underlying graphene nanosheets. Furthermore, DFT calculations indicated that the unique TiO2/graphene interface formed through electrostatic self-assembly can enhance other monovalent and multivalent ions storage such as Li+, K+, Mg2+ and Al3+. The proposed general strategy would open up a promising avenue for designing nanostructures for versatile materials and for broad applications in efficient energy storage systems.

Appendix A. Supporting materials Table for summary on the electrochemical performance of TiO2-based anode materials for sodium-ion batteries. Tables for DFT calculation details. Scheme for the synthesis of TiO2 nanocrystals. Scheme for the crystal structures of V2O5, ZnO, FeS2 and FePO4. Zeta potential of oxide, sulfide, and phosphates nanocrystals synthesized in the spary-flame reactor. Structure characterization on TiO2 nanocrystals, graphene and TiO2-graphene hybrid including XRD, SEM, TEM. Raman and XPS. Electrochemical characterization on TiO2 nanocrystals, graphene and TiO2-graphene hybrid including EIS, CV, charge/discharge and leakage current test. The Supporting 21 ACS Paragon Plus Environment

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Information is available free of charge on the ACS Publications website at DOI: xxxxx.

AUTHOR INFORMATION † These authors contributed to this work equally. Corresponding Author *To whom correspondence should be addressed. Hartmut Wiggers: Email: [email protected] (H.W.) Zonghai Chen: Email: [email protected] (Z.C.) Khalil Amine: Email: [email protected] (K.A.) Notes The authors declare no competing financial interest. Acknowledgements We gratefully acknowledge support from the U. S. Department of Energy (DOE), Vehicle Technologies Office. Argonne National Laboratory is operated for DOE Office of Science by UChicago Argonne, LLC, under contract number DE-AC0206CH11357. Research at State Key Lab of Xiamen University was funded by National Natural Science Foundation of China (Grants No. 21321062). Part of this work has been funded from the European Union’s Horizon 2020 Research and Innovation Programme, under Grant Agreement No. 646121 (NanoDome). L.S.X. acknowledges the financial support of “Programm zur Förderung des exzellenten wissenschaftlichen Nachwuchses” from University of Duisburg-Essen (UDE). The authors thank Dr. M. Heidelmann from Interdisciplinary Center for Analytics on the Nanoscale (ICAN), UDE for the HR-TEM characterization, Prof. Dr. S. Barcikowski and Prof. Dr. M. Winterer from UDE for the access to the zeta-potential measurements. The authos also thank the support from Clean Vehicles - US-China Clean Energy Research Center (CERC-CVC2). Reference 1. 2.

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Figure 1 a) Schematic illustration and b) digital images for the formation process of the TiO2-G hybrid. c) Schematic illustration for the proposed sodium storage mechanism of TiO2-G hybrid, where the reduced graphene oxide, O, Na+ and TiO2 NCs are in brown, red, yellow, and blue colors, respectively.


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Figure 2 a) HEXRD patterns of Graphene, TiO2 NCs and TiO2-G hybrid, b) N2 isothermal adsorption/de-sorption curve of graphene, TiO2 and TiO2-G hybrid, c) TEM image of TiO2 NCs, d) low and e) high magnification TEM images of TiO2-G hybrid, f) dark field HAADF TEM image and g) Ti and h) C elemental mapping of TiO2-G hybrid. Insets in panels c) and e) are high resolution TEM image of individual TiO2 NCs.


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Figure 3 a) Cycle performance at C/10 and b) rate capability of graphene, TiO2 NCs, TiO2+G mixture, and TiO2-G hybrid; c) cycle performance of TiO2-G hybrid and TiO2+G mixture at 2C; d) charge/discharge curves of TiO2-G hybrid at 2C; and e) cycle performance of TiO2-G hybrid at 10C of charge/discharge. The charge/discharge rates and capacities of TiO2+G mixture and TiO2-G hybrid are calculated based on the overall mass of TiO2+G mixture and TiO2-G hybrid.


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Figure 4 a) Static leakage current as a function of potential for TiO2-G electrode; b) HRTEM images of the deep discharged TiO2-G electrode (0.01V); c) and d) ex-situ Ti K-edge XANES spectra of TiO2-G hybrid electrode at different charge/discharge states; e) selected in-situ synchrotron HEXRD patterns; and f) integrated area of the anatase (101) reflection of the TiO2-G hybrid electrode during the first charge/discharge.


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Figure 5 Models of a) the graphene, b) TiO2 (101) surface, and c) TiO2/graphene interface; side views of the models with an increase of Na coverage from 0.125 monolayer (ML) to 1 ML on d) the TiO2 (101) surface and e) at the TiO2/graphene interface. Red: O; Cyan: Ti; brown: C; yellow: Na. f) Binding energies per Na atom with an increase of Na coverage on the TiO2 (101) surface and at the TiO2/graphene interface, respectively; g) positive charges per Na atom with an increase of Na coverage on the TiO2 (101) surface and at the TiO2/graphene interface, respectively; and h) negative charges in TiO2 and graphene with an increase of Na coverage.


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Figure 6 Side views of the most stable models with Li+, K+, Mg2+, and Al3+ coverage of 0.5 ML on a) the TiO2 (101) surface and b) at the TiO2/graphene interface.


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