Research ArticleMATERIALS SCIENCE

Transformation of oxide ceramic textiles from insulation to conduction at room temperature

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Science Advances  07 Feb 2020:
Vol. 6, no. 6, eaay8538
DOI: 10.1126/sciadv.aay8538

Abstract

Oxide ceramics are considered to be nonconductive brittle materials, which limits their applications in emerging fields such as conductive textiles. Here, we show a facile domino-cascade reduction method that enables rapid transformation of ceramic nanofiber textiles from insulation to conduction at room temperature. After putting dimethylacetamide-wetted textiles, including TiO2, SnO2, BaTiO3, and Li0.33La0.56TiO3, on lithium plates, the self-driven chemical reactions induce defects in oxides. These defects initiate an interfacial insulation-to-conductive phase transition, which triggers the domino-cascade reduction from the interface to the whole textile. Correspondingly, the conductivity of the textile sharply increased from 0 to 40 S/m over a period of 1 min. The modified oxide textiles exhibit enhanced electrochemical performance when substituting the metallic current collectors of lithium batteries. This room temperature reduction method can protect the nanostructures while inducing defects in oxide ceramic textiles, appealing for numerous applications.

INTRODUCTION

Oxide ceramics underpin a wide variety of modern technologies including electronics, energy, catalysis, and so on (1, 2). However, the materials are commonly hard and brittle. Moreover, they exhibit poor electronic conductivity due to the large bandgap (35). Conductive ceramic textiles that could be reversibly bent, polymer-like characteristic, would be extremely attractive for use in myriad engineering applications such as wearable electronics and energies.

One-dimensional structures such as fibers promise to solve the brittleness problem of oxide ceramics. Especially when the diameter of the fiber decreases from micrometer to nanometer, the size refinement will endow nanofibers (NFs) unique mechanical, thermal, and electrical properties (68). Recently, we developed a polymer template synthesis strategy to fabricate soft oxide NF textiles via an electrospinning method followed by calcination, which exhibited advantages of versatility, scalability, and controllability (913). With this approach, a wide variety of oxide textiles, including TiO2, SnO2, BaTiO3, and Li0.33La0.56TiO3, were fabricated (911).

On the other hand, the conductivity of oxides greatly relies on their surface oxygen defects (1417). Tremendous efforts have been focused on high-temperature reduction to create defects (1824). For example, Zhi et al. synthesized black TiO2 by placing the pristine TiO2 and Al metals separately in a two-zone tube furnace and then heated them for 6 hours at 500° and 800°C, respectively (19). However, the high-temperature reduction led to crystal growth and might destroy the nanocrystal structures. In addition, these methods were typically energy and time-consuming. Facile synthesis of conductive oxides at room temperature remains challenging.

Here, we report a rapid transition of oxide NF textiles from insulation to conduction at room temperature with a facile domino-cascade reduction method. Take TiO2 as an example; a soft TiO2 textile was first covered on a lithium (Li) plate without exerting stimuli. After dipping some drops of dimethylacetamide (DMAc) onto the textile, due to the spontaneous chemical reaction, defects were created in the interfacial TiO2 layer, which initiated an insulation-to-conduction phase transition and triggered the domino-cascade reduction from the interface to the whole textile. As a result, the color of the textile quickly changed from white to black in 1 min, corresponding to a sharp conductivity increment from 0 to 40 S/m. Furthermore, this strategy was universal and could be used for the fabrication of conductive Li0.33La0.56TiO3, BaTiO3, and SnO2 NF textiles. Freestanding electrodes based on these conductive textiles exhibited high capacity and stable cycling performance.

RESULTS

Mechanisms of the room temperature domino-cascade reduction

According to the Ellingham diagrams, Li metal can reduce a series of metal oxides (table S1) (25). The reaction equation between Li and TiO2 is as follows4Li+TiO2Ti+2Li2O (ΔG=233.6 kJ mol1at 298 K)

According to the Gibbs function, the reaction occurs spontaneously when ΔG < 0, which means that TiO2 can be reduced by Li metal at room temperature. However, the spontaneous reaction only happens when they contact. Figure 1A shows a general process of using the room temperature reduction strategy to fabricate soft conductive TiO2 textiles. Here, we selected DMAc as a diffusion solvent because of its high wettability with the oxide films (fig. S1A), which increased the contact areas between TiO2 and Li and facilitated the transfer of Li nanoparticles (NPs) or Li+ ions in the film.

Fig. 1 Mechanism illustration of the reduction process.

(A) General picture of using the room temperature domino-cascade reduction strategy to reduce TiO2 textiles. (B) Schematic diagram of the contact corrosion between Li and TiO2. (C) Snapshots of color changes of the reduced TiO2 textiles, which could be used as an electric wire to light a bulb (photo credit: Y. Zhang, Donghua University). E0, vacuum level; WM, work function of Li metal; EFM, electric fermi level of Li metal; WS, work function of TiO2; EFS, electric fermi level of TiO2; CB, conduction band; VB, valence band.

It is well known that electron transfer will occur on the contact surface due to their different electric Fermi levels (EF) (26). According to the work function (W = E0EF), Li has lower W than TiO2 due to its higher EF. Therefore, Li will donate electrons to TiO2 (Fig. 1B), and TiO2 was first reduced at the interface. These self-driven chemical reactions will continue as long as there is a Fermi level difference, and the contact corrosion creates lots of Li+ ions or Li NPs. The conductive interfacial layer works as a bridge and delivers electrons to the adjacent nonconductive TiO2 layers continuously, which create oxygen (O) vacancies in TiO2. At the same time, the Li ions or NPs, which transfer quickly from the interface to the whole film due to the microelectrical fields and the strong siphon effect of the nanofibrous structures, are ready to react with the deprived O atoms (Fig. 1A). The smooth electron flow and Li diffusion led to a layer-by-layer transformation of the textile from insulation to conduction (fig. S1B). We call this process as the “domino-cascade reduction.” A rapid color change from initial white to intermediate blue and then, finally, to black was observed within 1 min (Fig. 1C and movie S1). The conductivity of the film reached 40 S/m, and this film could be used as a wire to light a bulb.

To realize the rapid reduction process, the most critical point was the contact areas between Li and TiO2. As shown in Fig. 2A, there was no color change if they did not contact, even both of them were immersed in the DMAc solvent. But if the Li plate contacted one point with the TiO2 film, then the contacted point became black. When increasing the contact area, the contact part quickly turned black in 10 S, and the noncontact part could also turn black in a longer time (Fig. 2B). Notably, Li metal could react with DMAc, as indicated by the color change of DMAc when Li was immersed into the solvent. However, after taking out the Li metal and immersing the TiO2 film into the Li+-containing solvent, there was no color change in the film (fig. S2A). These results verified our proposed domino-cascade reduction process of the transfer of ions, NPs, and electrons from the contact surface to the whole film. We did a series of experiments to further illustrate this mechanism, as shown in fig. S3. When dipping DMAc from the corner (or center) of the film, the color of the film started to change from the corner (or center), while when directly putting the presoaked film (with DMAc) onto the Li plate, the whole film quickly changed at the same time.

Fig. 2 Material characterization.

(A and B) Captured photographs of the self-driven chemical reactions between Li and TiO2 in DMAc. (C) Cross-sectional image of a typical TiO2 film. (D) Single TiO2 NF. Surface morphology of (E) the pristine TiO2, (F) the reduced TiO2, and (G) the reduced TiO2 after being washed by HCl. Relationships between conductivity and (H) reduction time, (I) Li sizes and solvent amounts, and (J) type of solvents (photo credit: Y. Zhang, Donghua University).

In addition, other factors were also investigated. The importance of the porous structure of the NF film was first verified by pressing TiO2 NPs into a compact film. As shown in fig. S2B, after dipping DMAc onto the film until 6 hours, the color of the film started to change, and after 12 hours, the color was still blue. On the other hand, without using the Li plate, DMAc could not reduce the film (fig. S4A). Similarly, without using DMAc, the Li plate could not reduce the film if it did not apply an external force (fig. S4B). Although the film could be reduced by Li plate after pressing them with high pressure (e.g., 1000 lb) for 12 hours (fig. S4, C and D), the final color was blue not black, indicating that DMAc accelerated the reduction speeds and degrees. From these results, we can conclude that the reduction was closely related to the solvent diffusion in the film, which coincided with the proposed diffusion-reduction mechanism.

Material characterization and conductivity measurements

The thickness of a TiO2 textile film was 42.5 μm (Fig. 2C). The textile demonstrated softness without fragility after various deformations (fig. S5, A to C) and exhibited a much smaller bending rigidity of 10.72 mN compared with the nonwoven fabrics and tissues (fig. S5D). From the microscopic point of view, the NFs kept their structures well without breaking under bending states, further verifying the good flexibility of the oxide films (fig. S6, A and B). Both the cross-sectional images (fig. S6, C and D) and the surface morphology (Fig. 2D) show that the film had an intact structure, in which the NFs entangled and offered a high degree of interconnectivity. The average diameter of the NFs was calculated as 322 μm (fig. S6, E and F). The NFs had smooth surfaces, on which there were well-combined tiny TiO2 grains, as shown by transmission electron microscopy (TEM; Fig. 2E). The NFs had multiple meso- and micropore structures with a large specific surface area of 5.1 m2/g, as determined by the N2 adsorption-desorption measurement (fig. S7A). In addition, the tensile strength of the TiO2 NF films was 0.52 MPa (fig. S7B).

The pristine TiO2 had a high crystalline degree, as checked by x-ray diffraction (XRD; fig. S7C), in which the diffraction pattern of the sample shows the characteristic peaks belonging to the anatase TiO2 (PDF#00-065-5714). Although the characteristic peaks decreased a lot after the reduction, the reduced TiO2 exhibited an identical crystalline structure with the anatase phase, indicating that the Li reduction strategy did not change the intrinsic crystal structure of the TiO2 NFs. In addition, the diameters of the TiO2 NFs did not change much after the reduction (fig. S8, A and B). However, the average pore size of the film decreased greatly from 1 to 0.23 μm, as determined by a porosimeter (Porous Materials Inc.; fig. S7D). To find the reason, we characterized the surface (Fig. 2F) and cross-sectional (fig. S8C) morphology of the reduced TiO2. There were extra dense layers that blocked the pores. We attributed the dense layers to Li2O. As shown in Fig. 2G and fig. S8D, the layers were easily removed by dilute hydrochloric acid (HCl). As a comparison, after immersing the pristine TiO2 NFs into DMAc for 24 hours, there were no such layers (fig. S8E), indicating that the dense layer was caused by the reaction of Li and TiO2.

Since the domino-cascade reaction was influenced by the reduction time, contact areas, Li diffusion rates, and solvent types, we investigated the effects of these factors on conductivity. First, the effect of the reduction time is shown in Fig. 2H. The conductivity of the pristine TiO2 NFs was almost 0, which rapidly increased to greater than 10 S/m in 1 min. During this period, because of the formation of O vacancies in TiO2, the color of the film gradually changed from white to blue and black. After 1 min, the conductivity increased with a much slower rate. In addition, the effects of Li sizes and solvent amounts were investigated, as shown in Fig. 2I. Here, the area of the TiO2 film was 4 cm2, and the test time was 1 min. The areas of the applied Li plates (nos. 1 to 5) were 0.25, 1, 4, 5.3, and 12.3 cm2, respectively. The larger Li areas led to higher conductivity. However, once the Li plate was larger than the film, there were no obvious conductivity changes. On the other hand, the quantities of DMAc used in experimental nos. 1 to 5 were 20, 30, 40, 50, and 60 μl, respectively. The larger DMAc amounts always resulted in higher conductivity. These results further validated our proposed theory of “Li-diffusion reduction of TiO2.”

The rapid domino-cascade reduction greatly relied on the quantities of interfacial Li ions or NPs. According to the theory of contact corrosion, salts could enhance the contact corrosion and created more Li ions or NPs at the interface in a short time (27). With this in mind, we verified the effects of Li salts in the diffusion solvent on the conductivity. Here, 0.1 M concentration of LiPF6 in DMAc was used, and the test time was set as 1 min. As shown in Fig. 2J, the conductivity of the reduced films varied greatly, in which the LiPF6 contained in DMAc resulted in higher conductivity, while the N-methyl pyrrolidone (NMP) had the smallest one. Notably, after immersing Li plates in these three kinds of solvents for 3 hours, there were color changes as shown in the inset of Fig. 2J. The LiPF6 contained in the DMAc demonstrated the largest conductivity increment, while the NMP exhibited the smallest increment (fig. S8F), indicating that the conductivity of the reduced films was related to the reaction intensity between Li and the solvents.

Surface defect characterization of the pristine and the reduced TiO2 NFs

It is well known that surface oxygen defect is one of the most important factors to improve the conductivity of metal oxides. To elucidate the defects, crystal structures of TiO2 before and after reduction were examined by high-resolution TEM. The pristine TiO2 was highly crystallized with high-resolved lattice features (Fig. 3A). The lattice space was ~0.35 nm, corresponding to the (101) plane of anatase TiO2. This was in agreement with the XRD analysis. However, during the reduction process, the surface crystalline structure was gradually destroyed, as shown in Fig. 3 (B and C). TEM images that demonstrated the surface-phase changes at high magnifications are shown in fig. S9 (A to C). After 10 S, the black TiO2 showed a disordered surface layer with a thickness of ~3 nm. The disordered layer was ascribed to the deprivation of surface oxygen by Li during the corrosion process, which resulted in oxygen vacancy defects. To investigate the dynamic changes of the crystal structures during the Li reduction process, a series of XRD spectra of the reduced TiO2 with different reduction times was recorded (Fig. 3D). The diffraction angle moved to the left with the increase in reduction time, indicating the higher reduction degrees of TiO2.

Fig. 3 Defect characterization of the pristine and reduced TiO2.

(A to C) High-resolution TEM of the pristine, blue, and black TiO2 NFs. (D) Dynamic XRD analysis of the TiO2 NFs from white (at 0 S) to black (at 60 S). (E to G) XPS spectra of the pristine and black TiO2. (H and I) Raman spectra of the pristine and black TiO2 NFs. a.u., arbitrary unit.

To analyze the valence and structural changes caused by oxygen vacancies, we tested x-ray photoelectron spectroscopy (XPS) of the pristine and black TiO2 (Fig. 3E), in which O 1s and Ti 2p were provided. As shown in Fig. 3F, the three typical XPS peaks at ~529.6, 531.9, and 533.5 eV could be ascribed to lattice oxygen, oxygen defects, and oxygen species, respectively (28). The content of oxygen defects in the black TiO2 exhibited a notable increase compared with the pristine TiO2, indicating that lots of oxygen vacancies were created. Furthermore, the characteristic peaks of Ti 2p for the black TiO2 shifted to lower binding energies, suggesting the formation of Ti3+ from the generation of oxygen vacancies (Fig. 3G). By calculating the integral peak areas of Ti3+, the defect content in black TiO2 was ~18%. Further evaluation of the oxygen defects and structural changes by Raman spectroscopy is shown in Fig. 3 (H and I). The characteristic peaks at 144.7, 199.0, 399.5, 518.6, and 638.5 cm−1 corresponded to the six typical Raman-active modes of the anatase TiO2 phase (3Eg + 2B1g + A1g) (29, 30). Compared with the pristine one, the black TiO2 exhibited a blue shift and a broaden peak (the full width at half maximum) at the strongest mode of Eg, suggesting the formation of disordered surface domains, oxygen vacancy defects, and Ti3+ interstitial defects (31). The bandgaps of TiO2 before and after reduction were tested, as shown in fig. S9D. The black TiO2 exhibited a lower bandgap of 2.47 eV, in comparison with 3.15 eV of the pristine TiO2. Lower bandgap always resulted in higher conductivity. All these results were consistent with the TEM analysis. After Li reduction, disordered structures were formed on the surface of the TiO2 lattice.

Large-scale production and potential applications of conductive TiO2 textiles

The method for the fabrication of conductive TiO2 textiles broadens the applications of TiO2 in wearable electromagnetic interferrence shielding and sensors, flexible energy devices, facile catalytic and water splitting, and high-temperature military equipment (Fig. 4A). To advance these textiles into applications, it is necessary to increase their production while decreasing their cost substantially. With electrospinning, the morphology of fibers and the thickness of textiles could be easily adjusted to achieve the desired functionalities (32, 33). While a low temperature of 700° to 800°C and a short calcination time of 2 hours were applied for fabricating the metal oxide ceramic NFs, greatly reducing power energy consumption and time cost in comparison with the traditional solid-phase sintering techniques that required a high calcination temperature of >1000°C and a long pyrolysis time of >12 hours (34, 35).

Fig. 4 Application prospect of the reduced TiO2.

(A) Multifunctions of the conductive oxide ceramic textiles. (B) Large-scale production of the conductive oxide NF textiles and their potential applications as freestanding battery electrodes (photo credit: Y. Zhang, Donghua University).

On the other hand, the Li reduction strategy shows advantages including all-room-temperature processing, time efficiency, controllability, versatility, and scalability. As shown in Fig. 4B and fig. S10 (A to D), large-scale TiO2 textiles were fabricated, and the textiles could be reduced quickly with the room temperature reduction strategy. More attractively, the conductive TiO2 textiles maintained their good flexibility without any ruptures after deformations. Both the interconnected NF structures and the arrangement of TiO2 NPs in a single NF contributed to the high mechanical stability and softness of TiO2 films. As shown in fig. S10E, the single TiO2 NF had a structure similar with the classical microstructure of “brick-and-mortar,” in which the crystal TiO2 NPs were brick, and their grain boundaries were mortar, which could absorb or scatter the applied force when bending the fiber via microscale movements, thus avoiding the NF breaking under a small external force. When a bigger external stress was applied, the slippery between fibers will disperse stress concentration and thus block cracks.

The weight density of the TiO2 textile was 1.78 mg/cm2, which is much smaller than copper (Cu; 9.48 mg/cm2) and aluminum (Al; 3.16 mg/cm2) current collectors that are used in Li batteries (fig. S11, A to C). As a potential application, we substituted the easily corroded Cu current collector with the conductive TiO2 textile and directly adopted it as a freestanding electrode. The electrochemical performances of the pristine and black TiO2 textiles were compared in fig. S11 (D and E). At a charge/discharge rate of 0.5 C (1 C = 336 mA h/g), the capacity of the pristine TiO2 decreased rapidly from the initial 153.1 to 83.8 mA h/g over 100 cycles. By contrast, the capacity of the black TiO2 decreased gradually from the initial 182.1 to 120.7 mA h/g over 100 cycles. The black TiO2 delivered an ultrahigh initial Coulombic efficiency (CE) of 99.7%, while for the pristine TiO2, the value was as small as 49.3%. The initial irreversible capacity was mainly from the formation of solid electrolyte interface (SEI); the high CE of the black TiO2 electrodes was ascribed to the formation of SEI before cycling due to the Li reduction with LiPF6 contained in DMAc as the solvent. In addition, the capacity of the black TiO2 was always higher than that of the pristine TiO2. After 700 cycles, the capacity of the black TiO2 remained at 99.8 mA h/g. This enhanced electrode performance could be attributed to the oxygen defects, which facilitated the charge transfer on the interface. This inference was also verified by adopting the conductive TiO2 textiles in Li metal anode protection (fig. S11F).

DISCUSSION

In this work we demonstrated a facile and rapid room temperature Li reduction method to fabricate soft conductive TiO2 textiles. The low-temperature reduction protected the nanocrystal structures while inducing defects during the reactions. The mechanism of the domino-cascade reduction was explored, and the possible factors that influenced the rapid reduction processes were investigated. First, the difference in Fermi levels between Li and TiO2 led to self-driven contact corrosion between the two, which created a large of ions and NPs. At the same time, electrical fields and microbatteries were formed in the system. These microbatteries can transfer electrons and ions from the interface (as anode) to the nonconductive TiO2 layers (as cathode).

Second, the high porosity of the NF structures was also beneficial for the rapid reduction. The strong syphon effect of the NF structures facilitated Li diffusion from the interfacial layer to the whole film, while the ceramic NFs, to a large extent, could provide a large exposure surface to these active Li+ ions, Li NPs, and electrons, and thus enhanced the reduction process. In addition, the size effect and surface effect brought by diameter refinement will give these NFs unique thermal and electrical properties, which also enhanced the reactions between Li and TiO2. The preliminary results showed that the implanted oxygen vacancies and surface disorder layers played key roles in enhancing the electrochemical performance of TiO2.

More practically, this room temperature Li reduction strategy was scalable, and it could be used to obtain a wide variety of conductive oxide NF textiles such as SnO2, BaTiO3, and Li Li0.35La0.52TiO3 (LLTO). As an intuitive observation, after reduction, the color of SnO2 quickly changed from white to yellow, the BaTiO3 changed from white to blue, while the LLTO changed from white to dark black (fig. S12 and movie S2). The conductivities of SnO2, BTO, and LLTO after the reduction were measured as 0.34, 0.21, and 0.32 S/m, respectively. Similarly, there were dense layers on these reduced films, which could also be easily removed by dilute HCl. The surface NF morphology of these three textiles before and after reduction, as well as after being washed with HCl, is shown in fig. S12. In addition, the valence and structural changes of these textiles before and after the reduction were checked by XPS and XRD, as shown in fig. S13. Both the XPS and the XRD analyses for the reduced films showed the decreased valences of Ti and Sn. These conductive ceramic textiles have numerous potential applications. For example, SnO2 is a famous semiconductor and can be used as a channel material for flexible display and organic light-emitting diode. Li BaTiO3 (BTO) is a dielectric compound material that can be used in electronic ceramics such as thermistors and capacitors, while LLTO is an electrolyte material for solid batteries. The improved conductivity of these materials is expected to expand their applications.

In summary, we have successfully demonstrated a sustainable approach for the fabrication of conductive metal oxide ceramic textiles at ambient conditions in an environment-friendly manner by completely avoiding the use of catalysts and high-energy consumption. We consider that the energy- and time-saving reduction method promises to fabricate conductive oxide ceramic textiles that can be applied in emerging fields such as wearable conductive textiles. More practically, the facile NF fabrication technique and the room temperature domino-cascade reduction strategy enable the design of a range of hybrid materials that fall between the boundaries of what have typically been regarded as the limitations of soft and hard, conductive and nonconductive matters.

MATERIALS AND METHODS

Synthesis of TiO2 NF textiles

The synthesis contained two steps; the first step was to use a sol-gel electrospinning method followed by calcination to fabricate soft TiO2 NF textiles. First, polyvinylpyrrolidone (Mw = 1,300,000) powder was dissolved in a mixed solution of ethanol and acetic acid, and then the solution was stirred for 2 hours. Second, we added an appropriate amount of titanium (IV) isopropoxide into the above solution to prepare a transparent and homogeneous spinning solution. All the chemicals were purchased from Aladdin Chemical Co. Ltd. Third, electrospinning was performed on a DXES-4 equipment with a feed rate of 1.5 ml/hour, a voltage of 15 kV, and a rotation speed of 50 r/m. The humidity and temperature for the electrospinning process were 50 ± 2% and 25 ± 2°C, respectively. Last, the as-spun NF films were annealed at 700°C in air for 2 hours with a heating rate of 2°C/min. Similarly, SnO2, BaTiO3, and LLTO NF textiles were also obtained by electrospinning followed by calcination. The detailed synthesis conditions are shown in the Supplementary Materials. The second step was to use the room temperature diffusion-reduction method to prepare conductive ceramic textiles. These experiments were conducted in an argon-filled glovebox. The detailed processes are described in Fig. 1A.

Characterization techniques

Morphology of the oxide ceramic NF textiles was characterized by scanning electron microscopy (Hitachi S-4800) and high-resolution TEM (JEM-2100F). XRD (D/Max-2550 PC) with Cu Kα radiation (λ = 1.5406 Å) was used to test the crystal structures of the samples. XPS spectra of the samples were collected using ESCALAB 250Xi (Thermo Fisher Scientific, USA) with Al Kα source. Raman spectra were recorded on a Nicolet Nexus 670 + Raman Module (Thermo Fisher Scientific) with 532-nm excitation source. The specific surface area and pore structures were measured by Brunauer-Emmett-Teller (BET; ASAP 2020) through nitrogen (N2) adsorption-desorption isotherms. The mechanical properties of the TiO2 NFs and the tensile strength of the silms (5 × 0.3 cm2 and thickness of 50 ± 5 μm) were measured on a tensile tester (XQ-1C) according to the standard (ISO 1798:2008). The absorption spectra of the materials were recorded using an ultraviolet-visible spectrometer (U-3900) with an integrating sphere. The pore sizes of the NF films were analyzed by a capillary flow porometer (Porous Materials Inc., CFP-1100AI). To evaluate the flexibility of the resultant samples (with size of 10 × 10 cm2), the bending rigidity instrument (RRY-1000, Hangzhou Qingtong & Boke Automation Technology Co. Ltd., China) was used according to the standard (ASTM D 2923-95). The electronic conductivity was measured by an ST-2258C multifunction digital four-probe tester. Before testing the conductivity, the generated Li2O of the reduced textiles were removed by dissolving them with diluted HCl and then were dried. All of the experiments were conducted at room temperature unless described otherwise.

Battery assembly and electrochemical measurements

CR2032-type coin cells were assembled in an argon-filled glovebox. TiO2 ceramic NF membranes before and after reduction were used as the electrode, and an Li plate with a thickness of 20 μm and a diameter of 1.65 cm was used as a counter electrode for the battery assembly. The electrolyte used here consisted of 1 M LiPF6 in ethylene carbonate (EC), dimethyl carbonate (DMC), and ethyl methyl carbonate (EMC), whose volume ratio is 1:1:1. Cellgard 2400 microporous membranes were used as separators. A CHI760e electrochemical workstation was used for cyclic voltammetry measurements with a potential range between 0 and 2.5 V at a scan rate of 0.5 mV/s. The electrochemical cycling was carried out in galvanostatic mode on a Land CT2001A tester (Shanghai, China) at room temperature. The applied cycling rate was 0.5 C, whereas 1 C was equal to 336 mA h/g.

SUPPLEMENTARY MATERIALS

Supplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/6/6/eaay8538/DC1

Supplementary Materials and Methods

Fig. S1. Process of domino-cascade reduction for soft TiO2 NF textiles.

Fig. S2. Effects of the Li+-containing DMAc and NF structures on TiO2 reduction.

Fig. S3. Experimental verification of the domino-cascade reduction mechanism.

Fig. S4. Verification of the effects of solvents and Li plates on TiO2 reduction.

Fig. S5. Optical photographs of the fabricated TiO2 textiles and comparison of flexibility of the TiO2 textiles with nonwoven fabrics and tissue papers.

Fig. S6. The cross-sectional and surface images of TiO2 NFs and the diameter histogram.

Fig. S7. Surface area, mechanical, crystal structure, and pore size characterizations of TiO2 NFs.

Fig. S8. Morphology and conductivity characterizations.

Fig. S9. High-resolution TEM of the pristine, blue, and black TiO2 NFs at high magnifications, and Tauc plots of pristine and reduced TiO2 NFs.

Fig. S10. Large-scale production of the pristine and reduced ceramic NF textiles.

Fig. S11. Potential applications of the reduced TiO2 textiles as battery current collectors.

Fig. S12. Surface morphology of the pristine and reduced SnO2, BTO, and LLTO textiles.

Fig. S13. XPS and XRD spectra of the pristine and reduced SnO2, BTO, and LLTO.

Table S1. Comparison of electrode potentials of Li, K, Na, Mg, Zn, Ti, and Sn.

Movie S1. The processes of using the room temperature Li reduction strategy to reduce TiO2 textiles.

Movie S2. The processes of using the room temperature Li reduction strategy to reduce LLTO textiles.

This is an open-access article distributed under the terms of the Creative Commons Attribution-NonCommercial license, which permits use, distribution, and reproduction in any medium, so long as the resultant use is not for commercial advantage and provided the original work is properly cited.

REFERENCES AND NOTES

Acknowledgments: We thank others for any contributions. Funding: This work was supported by the National Natural Science Foundation of China (no. 51702047 and no. 51873029), the National Key R&D Program of China (no. SQ2018YFC200227), the State Administration of Science, Technology and Industry for National Defence, PRC (JCKY2018203c035), the Young Elite Scientists Sponsorship Program by CAST (no. 2018QNRC001), the Innovation Program of Shanghai Municipal Education Commission (no. 2017-01-07-00-03-E00024, and the Program of Shanghai Academic Research Leader (no. 18XD1400200). Author contributions: J.Yan and B.D. conceived and designed the experiments. Y.Zhang synthesized the materials and did characterizations. Y.Zhang, S.L., S.X., J.S., and Y.Zhao conducted battery tests and analysis. J.Yan and Y.Zhang wrote the paper, and all authors contributed to the discussions and revision of the paper. Competing interests: The authors declare that they have no competing interests. Data and materials availability: All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials. Additional data related to this paper may be requested from the authors.
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