Reduced Graphene Oxide Hybrid Films for

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Cite This: ACS Appl. Mater. Interfaces 2018, 10, 3634−3643

3D Porous MXene (Ti3C2)/Reduced Graphene Oxide Hybrid Films for Advanced Lithium Storage Zhiying Ma,†,‡ Xufeng Zhou,*,† Wei Deng,†,‡ Da Lei,†,‡ and Zhaoping Liu*,† †

Key Laboratory of Graphene Technologies and Applications of Zhejiang Province and Advanced Li-ion Battery Engineering Lab, Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Zhejiang 315201, P. R. China ‡ University of Chinese Academy of Sciences, 19 A Yuquan Rd, Shijingshan District, Beijing 100049, P. R. China S Supporting Information *

ABSTRACT: MXenes, as a new family of 2D materials, can be used as film electrodes in energy storage devices because of their hydrophilic surface, metallic conductivity, and rich surface chemistries. However, the poor ion transport of MXene film electrodes causes a great loss of surface reactivity, which significantly inhibits the full exploitation of the potential of MXene-based materials. To solve this issue, we report a facile electrolyte-induced self-assembly method to construct a 3D porous structure in the MXene−rGO hybrid film, which effectively facilitates rapid diffusion and transport of electrolyte ions in the film electrode while still maintaining high electrical conductivity. When the hybrid film is employed as electrode materials for lithium-ion batteries, it exhibits high specific capacity of 335.5 mA h g−1 at 0.05 A g−1 and good rate capability of 30% capacitance retention at 4 A g−1. Additionally, the film electrode exhibits excellent cycling stability without capacity decay after 1000 cycles under high rates (1 A g−1) owing to its stable structure. Furthermore, the electrochemical analysis also demonstrates that the novel 3D porous microstructure plays an important role in the fast reaction kinetics and high capacity of the MXene−rGO hybrid film electrode. This work may provide a new strategy to solve the issues related to poor ionic transport in MXene-based film electrodes. KEYWORDS: MXenes, MXene−rGO, film electrode, two-dimensional materials, 3D porous structure, lithium-ion batteries

1. INTRODUCTION MXenes are an emerging family of 2D transition metal carbides and nitrides with a general formula of Mn+1XnTX, where M is an early transition metal, X represents C, and/or N, TX denotes surface functional groups (−O, −OH, −F), and n = 1, 2, or 3. Generally, MXenes are prepared by selective etching of the Agroup (generally, group III A or IV A) element layers from MAX phase precursors (Mn+1AXn) which comprise a >70membered family of layered, hexagonal early-transition-metal carbides and nitrides.1,2 Ti3C2TX,3 as the first reported MXene material, has been widely researched in energy storage applications,4−6 including supercapacitors7−11 and Li-ion12−18 and Na-ion batteries19−21 because of its metallic conductivity,1 high volumetric capacity,8 and hydrophilic surfaces.13 Furthermore, delamination of Ti3C2TX produces nanosheets with a thickness of several nanometers and a lateral size of the order of micrometers.22−24 Therefore, the MXene nanosheets can be used as the basic units to construct film electrodes with different microstructures,5 which also display advanced energy storage performance. However, the MXene nanosheets tend to stack densely in the film electrode because of their two-dimensional layered structure, which seriously hinders the infiltration of the electrolyte, resulting in poor electrochemical performance. To © 2018 American Chemical Society

fully utilize their electrochemical performance, strategies including creating porous structures,25 introducing interlayer spacers,26−37 and integrating two-dimensional (2D) nanosheets into 3D macroscopic structures38 were proposed. For example, Gogotsi and co-workers reported that Li-ion storage capacity was increased by a factor of four after creating mesopores in MXene nanosheets, along with an excellent rate performance.25 Peng’s group demonstrated that porous Ti3C2/CNT composite electrode materials showed better electrochemical performance than the pure Ti3C2 because of the introduction of CNTs which improved electrical conductivity and provided efficient ion transport channels in the binder-free anodes.27 Recently, Gogotsi’s group provided a good solution to enhance the ionic transport in the film electrode by processing 2D Ti3C2TX into 3D macroporous frameworks using sacrificial poly(methyl amethacrylate) spherical templates, and the as-fabricated macroporous Ti3C2TX films were free-standing, flexible, and highly conductive and exhibited outstanding Na-ion storage performance.38 These methods propose good solutions to solve the issue of poor ionic transport of the film electrodes. Received: November 14, 2017 Accepted: January 3, 2018 Published: January 3, 2018 3634

DOI: 10.1021/acsami.7b17386 ACS Appl. Mater. Interfaces 2018, 10, 3634−3643

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ACS Applied Materials & Interfaces

2. EXPERIMENTAL SECTION

However, they have certain disadvantages which cannot be neglected. For example, the process of creating mesopores in MXene nanosheets is complicated and environmental unfriendly because of the use of hydrofluoric acid. The introduction of spacers into MXene sheets plays a limited role in improving the ion transportation in film electrodes because the interlayer distance enlarges in the range of nanoscale,28,37 and some spacers significantly decrease the conductivity of the MXene electrode, such as polymers,33 metal-oxide nanoparticles,36 and large ions.32 Even though the sacrificial template approach can integrate the 2D MXene into 3D macroscopic structures, adequately facilitating the infiltration of the electrolyte in the film electrode, it needs high temperature to remove the template, which may decrease the conductivity of MXene, for MXene is easily oxidized at high temperatures. Therefore, it is necessary to adopt new strategies to construct MXene electrodes with 3D porous structures via moderate approaches. As recently reported, the assembly of 2D materials into 3D architecture is the best strategy to fabricate a binder-free electrode with fast mass transport in the electrode, thereby resulting in high-performance devices.39 In particular, the 3D graphene framework has been synthesized by integrating graphene materials through self-assembly, offering significantly enhanced performance in applications such as energy, environment, sensing, and biological fields. For example, our group has successfully fabricated 3D porous graphene materials with oriented pore structure by electrolyte-assisted self-assembly and vacuum filtration, which showed superior elasticity and anisotropic mechanical, electrical, and thermal properties.40 When this porous graphene network was used in Li−S batteries as the host for sulfur, the cell exhibited excellent rate capability and outstanding cyclic stability because the graphene network provided fast electron-transportation pathways and facilitated rapid diffusion of the electrolyte.41 Similar to graphene oxide (GO), stable delaminated MXene solution can be obtained because of the strong electrostatic repulsion between MXene sheets induced by oxygen/fluorine-contained functional groups.22,24 However, it is difficult to construct 3D-macrostructure MXene frameworks by self-assembly because of the relatively small size and the rigidity of MXene flakes. Inspired by the successful construction of 3D macroporous graphene network and the structural similarity between GO sheets and MXene flakes, we propose a new strategy to fabricate porous MXene-based electrode materials under the assistance of GO. By adding GO into MXene, a stable complex solution is formed owing to the strong electrostatic repulsion between the two components. Then, the 3D porous MXene (Ti3C2)−GO hybrid material was obtained by electrolyteassisted self-assembly and vacuum filtration. Specifically, the addition of GO induced the formation of a 3D porous structure and a MXene−GO hybrid scaffold was produced in the hybrid film because of the addition of the electrolyte, which breaks the electrostatic repulsion between the MXene and GO sheets. After moderate postprocessing for the reduction of GO to reduced GO (rGO), the MXene−rGO hybrid films were obtained. Importantly, when the Ti3C2−rGO film was employed as the anode material for LIBs, it exhibited outstanding rate capability (98.9 mA h g−1 at 4 A g−1) and excellent cycle performance (212.5 mA h g−1 after 1000 cycles at 1 A g−1 without degradation), which was much better than that of pure MXene and rGO electrode.

2.1. Synthesis of MXene (Ti3C2) and GO Solutions. Ti3C2TX was synthesized by etching the Ti3AlC2 phase (400 mesh, Figure S1b) with HF as reported.3 Typically, Ti3AlC2 powder (1 g) was slowly added into HF (40 wt %, 10 mL) solution, and the reaction lasted for 24 h at 40 °C under stirring. The resultant precipitates were repeatedly washed with deionized water and centrifuged at 3500 rpm for 5 min. The obtained sample was dried at 80 °C under vacuum for 10 h. Then, 1.0 g of the dried HF-pretreated sample (Figure S1c) was immersed in aqueous solution of tetramethylammonium hydroxide (TMAOH, 25 wt %, 10 mL) and reacted for 24 h under stirring, after which the products (Figure S1d) with TMA+ intercalation were collected by repeated centrifugation and thorough washing with pure H2O to remove the residual TMA+ and OH−.42 Ti3C2 nanosheets were finally obtained by carefully dispersing the TMA+-intercalated slurry in H2O and sonicating for 10−20 min under the protection of an inert environment at room temperature. GO was synthesized from natural graphite flakes (80 μm) following a modified Hummers method as reported in our previous paper.43 To prepare the GO solution, the obtained graphite oxide was dispersed in water by sonication for 15 min, reaching a concentration of 1 mg mL−1. 2.2. Synthesis of Ti3C2, rGO, Ti3C2−rGO, and Pure Ti3C2 Film. The typical synthesis procedure of the Ti3C2−GO hybrid film is as follows. A 100 mL Ti3C2 solution (0.2 mg mL−1) and a selected amount of GO (1 mg mL−1) were mixed together under continuous stirring. After 1 h, 0.8 g ammonium bicarbonate (0.1 M based on the Ti3C2 solution) was added to the mixed solution and sonicated for 1 min. Then, the composite was separated from the solvent through filtration and directly dehydrated via a freeze-drying process to maintain a 3D porous architecture. Hybrid films of various compositions with Ti3C2−GO weight ratios of 4:1, 3:1, 2:1, 1:1, and 1:2 were prepared by adjusting the volume of the GO solution (1 mg mL−1). For comparison, pure GO and Ti3C2 films were also prepared from their respective colloidal suspension of GO and Ti3C2 sheets with the same procedures. Finally, all films were subjected to hydrazine vapor for 10 h at 90 °C and annealed at 200 °C for 2 h under Ar atmosphere to obtain rGO, Ti3C2, and Ti3C2−rGO films. To detect the primitive natures of the MXene film, the pure Ti3C2 (P-Ti3C2) film was also prepared through filtering MXene solution without the addition of ammonium bicarbonate and reduction by hydrazine vapor and annealing. 2.3. Materials Characterization. Powder X-ray diffraction (XRD) patterns were recorded on a D8 X-ray diffractometer using Cu Kα irradiation (λ = 1.5406 Å) at 40 kV, 20 mA over the 2θ range from 5 to 70°. The microstructure of all samples was characterized by field-emission scanning electron microscopy (SEM, FEI, Sirion 200) and transmission electron microscopy (TEM, JEOL 2100F). Chemical compositions of the samples were analyzed by high-resolution X-ray photoelectron spectroscopy (XPS) recorded with an AXIS ULTARDLD spectrometer equipped with a monochromatic Al Kα X-ray source (hν = 1486.6 eV). Peak fitting was carried out using CasaXPS version 2273. Raman spectra were measured on a Renishaw inVia spectrometer using a laser of 514.5 nm at room temperature. A dimension 3100 atomic force microscope was used to analyze the thickness distribution of MXene sheets. Nitrogen adsorption/ desorption isotherms were recorded at 77 K using a Micrometritics ASAP 2020 analyzer. 2.4. Electrochemical Measurements. For preparing working electrodes, all films (Figure S5a) were cut into disk electrodes (dia. = 1.3 cm, Figure S5b), and each film electrode was about 2−3 mg. The electrochemical performance of the film electrode was evaluated via CR2032 type coin cells with a pure lithium foil as the counter electrode, Celgard 2300 as the separator, and 1 M LiPF6 in a mixture of ethylene carbonate and dimethyl carbonate (v/v = 1:1) as the electrolyte. The cells were assembled in an argon-filled glovebox with concentrations of moisture and oxygen below 1.0 ppm. After assembling, the cells were charged and discharged in a LAND CT 2001A system at various current densities in the voltage range of 0.005−3.0 V. Cyclic voltammograms (CV) were measured on a 3635


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Figure 1. Schematic illustration of the fabrication of the MXene, rGO, and MXene−rGO films.

Figure 2. (a) XRD patterns and (b) Raman and (c) XPS spectra of GO, rGO, P-Ti3C2, Ti3C2, and Ti3C2−rGO films.

films were directly dehydrated via a freeze-drying process. For improving the electronic conductivity, hybrid films were subjected to hydrazine vapor, aiming to reduce GO to rGO. Then, the as-obtained films were annealed at 200 °C for 2 h under Ar atmosphere to remove the residual ammonium bicarbonate, water, and some functional groups, which may have an adverse influence on the electrochemical performance of hybrid films. The hybrid films are denoted as x-y-Ti3C2− rGO, where x and y are the weight ratio of MXene (Ti3C2TX) to GO nanosheets in the composites. For comparison, pure MXene and rGO films were also prepared by the same procedures except that the original solution was of pure GO and pure MXene, respectively. Because of the small size and the rigidity of MXene flakes, MXene sheets are linked together tightly by the NH4+ ions in the solution, resulting in the dense stacking of the Ti3C2 layers in the film. However, for GO sheets, the GO microclusters formed in the solution prevent dense stacking of GO sheets and align together to construct 3D porous microstructure in the GO film. Therefore, GO sheets contribute to the formation of the 3D porous structure in the hybrid film, which prevents the dense stacking of MXene flakes. After adding 0.1 M NH4HCO3 to the MXene colloidal suspension (Figure S2a), an obvious flocculation (Figure S2b) can be found due to the electrostatic interaction between the nanosheets and NH4+ ions. The pure Ti3C2 (P-Ti3C2) and

Solartron 1400 workstation with the testing voltage between 0.005 and 3.0 V at different scan rates. The electrochemical impedance spectroscopy (EIS) measurements of the active material was recorded on an electrochemical workstation (Solartron 1400) using the frequency response analysis. The impedance spectra were obtained by applying a sine wave with an amplitude of 5.0 mV over the frequency range from 100 kHz to 0.01 Hz.

3. RESULTS AND DISCUSSION Figure 1 schematically illustrates the fabrication of the densestacked MXene film and 3D porous rGO and MXene (Ti3C2)− rGO composite films by adopting the electrolyte-assisted selfassembly and vacuum filtration method proposed in our previous reports.40 By adopting the preparation procedure of the Ti3C2−rGO hybrid film, for example, first, GO and MXene solutions were mixed together, forming a stable homogeneous solution (Figure S2e) via vigorous stirring for 1 h. Second, ammonium bicarbonate (NH4HCO3) was added to the mixed solution and the appearance of NH4+ ions broke the electrostatic balance between MXene and GO sheets, resulting in the formation of Ti3C2−GO hybrid microclusters (Figure S2f), in which MXene and GO nanosheets are strongly linked together by NH4+ ions. During the subsequent dehydration process by vacuum filtration, the hybrid microclusters gradually cross-linked and the 3D porous composite film was obtained finally. To maintain the 3D porous architecture, the hybrid 3636


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Figure 3. (a,d) Cross-sectional SEM images of the Ti3C2 film. (b,e) Cross-sectional SEM images of the 1-1-Ti3C2−rGO hybrid film. (c,f) Crosssectional SEM images of the rGO film.

Ti3C2 films can be prepared through filtering the MXene solution (Figure S2a) and MXene solution containing NH4HCO3 (Figure S2b), respectively. Their XRD patterns (Figure 2a) show that these two films display similar and good crystalline structure, which demonstrates the highly ordered stacking of the Ti3C2 flakes in films. Nevertheless, the (002) plane of the Ti3C2 film locates at a higher angle than that of the P-Ti3C2 film because of the removal of the water and NH4+ ions between the layers after hydrazine vapor treatment and annealing at 200 °C for 2 h under Ar atmosphere. GO can be reduced effectively to rGO by this moderate postprocessing, which is evidenced by the appearance of the characteristic peak of the (002) plane of rGO. For the Ti3C2−rGO film, its XRD shows a characteristic peak of rGO at around 26°,44 and no obvious peak can be indexed to Ti3C2, which can be ascribed to the disordered stacking of the Ti3C2 sheets in the 3D hybrid film. To further confirm the existence state of Ti3C2 nanosheets in the hybrid film, Raman spectra for selected samples are presented in Figure 2b. For GO and rGO films, two legible peaks at 1346 cm−1 (D band) and 1589 cm−1 (G band) can be observed, which are typical existence of graphene materials. Compared with GO, a high ID/IG ratio is obtained for the rGO film, indicating that GO is deeply reduced to rGO by the hydrazine vapor reduction method.45 Meanwhile, the P-Ti3C2 and Ti3C2 films show similar Raman spectra, indicating that the postprocessing has no obvious effect on the Ti3C2 film. Specifically, the mode at 200 cm−1 is A1g symmetry out-of-plane vibrations of Ti atoms, whereas the modes at 388 and 591 cm−1 are the Eg group vibrations, including in-plane (shear) modes of Ti, C, and surface functional group atoms. After the incorporation of rGO, two broad bands appear at 1346 and 1589 cm−1 for the Ti3C2−rGO hybrid film, which are the characteristic peaks of the D and G bands of graphitic carbon, confirming the presence of rGO nanosheets in the hybrid films. Compared with the Ti3C2 film, the Eg band of the Ti3C2−rGO film red-shifts by 38.6 to 629.6 cm−1, which is probably caused by the interaction between MXene and rGO sheets.37 To confirm the chemical composition and surface electronic states of the selected films, XPS spectra were tested to detect the component change of various films. As shown in Figure 2c, the C/O ratio of the GO film is 2.2, and it reaches 12 after reduction, indicating an effective reduction effect of the postprocess. High-resolution XPS spectra of C 1s core levels of both GO and rGO films (Figure S3a,b) can be fitted with

four components centered at 284.6, 284.9, 286.8, and 287.9 eV, which can be assigned to C−C/CC, C−N, C−O, and CO bonds, respectively. The dramatically decreased peak intensity of C−O bonds in rGO films suggests the removal of oxygencontaining groups by the reduction with hydrazine. The XPS also indicates that the MXene films are mainly composed of C, Ti, O, and little F, demonstrating that oxygen-containing groups are dominant on the surface of MXene because of the TMAOH intercalation process. After the treatment of hydrazine vapor and annealing, the O content of the Ti3C2 film decreases from 41.01 (at. %) to 29.62 (at. %). Furthermore, the high-resolution C 1s XPS spectra (Figure S3c,d) also displays that the content of C−O in the Ti3C2 film decreases obviously after postprocessing, owing to the reduction process of hydrazine vapor. Obviously, the C 1s spectra (Figure S3e) of the Ti3C2−rGO film contain characteristic signals from both Ti3C2 and rGO, confirming the hybridization of MXene and rGO. The Ti 2p spectrum (Figure S3f) of the Ti3C2−rGO film is similar to that of the Ti3C2 film, and the atomic percentage of Ti−O decreases from 45.1% for Ti3C2 film to 29% for the 1-1-Ti3C2−rGO hybrid film, implying that the Ti3C2 sheets were not oxidized, thereby retaining the high conductivity of MXene flakes. . Figure 3a−c show the cross-section of Ti3C2, Ti3C2−rGO, and rGO film, respectively. It can be seen that the Ti3C2 (Figure 3a) displays a dense thin film morphology with a small thickness of about ∼2 μm, whereas both Ti3C2−rGO and rGO films (Figure 3b,c) show a 3D porous microstructure with much larger thickness of ∼30 and ∼250 μm, respectively. The obvious morphological difference between Ti3C2 and rGO films prepared under identical conditions can be ascribed to the different combined forms of their respective layers in the solution. The Ti3C2 flakes have a smaller size and higher distribution density of hydrophilic functional groups (fluorine/ oxygen-containing groups) than the GO sheets. Therefore, the Ti3C2 nanosheets were linked together tightly by NH4+ ions in the solution, leading to the dense-stacked structure. However, for GO sheets, the functional groups distribute unevenly and enrich around the edges of the nanosheets, resulting in the selective link of GO sheets by NH4+ ions and the formation of porous microclusters. During the subsequent dehydration process by vacuum filtration, the microclusters gradually cross-link, forming a 3D porous film. Nevertheless, the Ti3C2−rGO hybrid film displays a porous structure (Figure 3b), which is different from Ti3C2 but similar to the rGO film, 3637


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Figure 4. TEM images of rGO (a), Ti3C2 (b), and Ti3C2−rGO (c) films. (d) EDS result of the selected area marked in panel (c) with a white-dashed line.

demonstrating the contribution of GO sheets in constructing the 3D porous structure in the hybrid film. The highmagnification SEM image (Figure 3d) of the Ti3C2 film further confirms the layer-by-layer stacking of MXene flakes in the film. However, the Ti3C2−rGO film (Figure 3e) shows macroporous structure, similar to the rGO film (Figure 3f). However, the walls constructing the porous Ti3C2−rGO film possess a multilayer structure (inset, Figure 3e), which is different from the single-layer structure of graphene walls in the rGO film (inset, Figure 3f), indicating the stacking of Ti3C2 and rGO sheets in the hybrid film. To further confirm the distribution of MXene and rGO sheets in the hybrid film, an energy dispersive spectrometer (EDS) was employed. The SEM-EDS mapping (Figure S4) shows that the distribution of Ti, C, O, F, and N atoms are homogeneous on the cross-section of the Ti3C2− rGO film, which means that Ti3C2 and graphene sheets are mixed evenly in the hybrid film. Furthermore, the elementcontent analysis (Figure S5d) confirms that Al atoms are etched to a large extent and the oxygen-containing functional groups become dominant on the surface of the Ti3C2 sheets, which is accordant with the XPS results. A transmission electron microscope was also used to observe the structures of rGO, Ti3C2, and Ti3C2−rGO films. As shown in Figure 4a,b, Ti3C2 sheets have a smaller lateral size than graphene sheets, which is also reflected by SEM characterization (Figure S6a,b). Also, Ti3C2 sheets do not have obvious wrinkles as graphene does, which indicates that the MXene sheets are not as flexible as graphene. Furthermore, the atomic force microscopy results (Figure S6c,d) demonstrate that the average thickness of MXene is ca. 4 nm, corresponding to 4 MXene monolayers. MXene sheets are thicker than graphene sheets; therefore, they have high mass thickness contrast under TEM. As shown in Figure 4c, small MXene sheets are spread on large graphene sheets, forming a hybrid scaffold evidenced by the obvious contrast between two components. The TEM− EDS result (Figure 4d) of the selected area in Figure 4c shows that the atomic content of Ti, C, O, N, and F is 30.37, 47.76,

19.11, 1.78, and 0.96%, respectively, also indicating the coexistence of MXene and graphene sheets in the hybrid film. The electrochemical properties of the film electrodes were evaluated using CR2032 coin-type cells with lithium metal as the counter electrode. The CV curves of the 1-1-Ti3C2−rGO film are presented in Figure 5a. A sweep rate of 0.2 mV s−1 and a voltage range from 0.005 to 3 V were employed to investigate the lithium storage behavior of the electrode. In the first cathodic scan process, two reduction peaks can be observed. The first peak around 1.5 V can be attributed to the formation of a solid electrolyte interphase (SEI) in the working electrode.27 The second peak around 0.75 V may be ascribed to the trapping of Li+ between Ti3C2 and rGO sheets. When the anodic sweep proceeds, a wide oxidation peak around 1.5 V can be ascribed to the extraction of Li+ from Ti3C2 and rGO sheets.18 The well-overlapped CV curves from the second to fifth cycles demonstrate the good reversibility of the hybrid film in the electrochemical reaction process. Figure 5b presents the potential versus capacity profiles of the hybrid film in the initial five cycles at a rate of 0.05 A g−1, which is accordant with the CV results. The 1-1-Ti3C2−rGO film displays high discharge and charge capacities of 940 and 473 mA h g−1, respectively, during the first cycle. The irreversible capacity is generally considered to be derived from some irreversible processes, including the formation of an SEI layer and the decomposition of the electrolyte. In the subsequent cycles, the overlap of the discharge−charge curves indicates good stability of the film electrode. Moreover, the charge−discharge curves (Figure 5c) of the 1-1-Ti3C2−rGO electrode show slow decay of the reversible capacity with the increase of the current, demonstrating its excellent rate performance. Even under a high current density of 4 A g−1, the 1-1-Ti3C2−rGO electrode also shows a decent specific capacity of 98.9 mA h g−1, which has rarely been reported previously. Figure 5d compares the rate performance of Ti3C2−rGO films with different mass ratios of Ti3C2 to rGO along with pure Ti3C2 and pure rGO film electrodes. The 1-1-Ti3C2−rGO 3638


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Figure 5. (a) CV profiles of the 1-1-Ti3C2−rGO film electrode. (b) The charge−discharge profiles of the 1-1-Ti3C2−rGO film electrode at 0.05 A g−1 in the initial five cycles. (c) The charge−discharge profiles of the 1-1-Ti3C2−rGO film electrode at different current densities. (d) Rate capability of different samples. (e) Cycle performance of different samples at 1 A g−1 (the line colors of different samples are identical to those in panel (d)). (f) Histogram showing the capacity retention ratio of different samples at 1 A g−1 after 1000 cycles.

sheets inhibit fast ion transport, also resulting in low capacity at high rates. However, the hybrid film inherits the advantage of graphene in the formation of a porous structure and MXene in high electrical conductivity. As shown in Figure S7, the crosssectional images of various hybrid films are porous, and the thickness increases with the increasing content of rGO in the hybrid film. This is because graphene sheets can help to construct a more open structure, resulting in the thickening of the hybrid film. Also, the nitrogen adsorption/desorption study (Figure S8) also shows that the more rGO content leads to a larger specific surface area of the film, indicating the increased exposed surface area of the material. Therefore, more active sites can be utilized in the film electrode, which can significantly promote the full exploitation of the potential of MXene-based materials. The 3D porous structure can provide efficient ion transport channels in the MXene-based electrode, but the relative poor conductivity of rGO sheets increases the electron

composite shows the highest capacity at all currents. Besides, when the current density returns to 0.05 mA g−1, the 1-1Ti3C2−rGO film electrode recovers a high value, implying its good reversibility. It is obvious that pure Ti3C2 and rGO films show inferior rate performances, whose charge capacities are 33 and 16.7 mA h g−1, respectively, at high current density (4 A g−1) with low capacity retentions of 15.1% (Ti3C2) and 5.2% (rGO). This is because both ion transmission rates and electron conductivity influence the electrochemical performance of the film electrode, especially at high rates. For the rGO film, the 3D porous structure facilitates ion transportation in the electrode, but the relatively low electron conductivity hinders the rapid transfer of electrons, thereby leading to an inferior rate performance. As shown in Table S1, the conductivity of the pure rGO film is 1.64 S/cm, which is much lower than that of the pure Ti3C2 film (2863 S/cm). Though the MXene film possesses metallic conductivity, the dense-stacked MXene 3639


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Figure 6. (a) CV profiles of the 1-1-Ti3C2−rGO film electrode at different scan rates. (b) The relationship between current and scan rates from 2 to 10 mV s−1 for the 1-1-Ti3C2−rGO electrode at 2 V. (c) EIS spectra of all electrodes (inset is the magnification of the high-frequency region). (d) Electron transfer resistance (Rct) of all electrodes. (e) Linear fit showing the relationship between Z′ and ω−1/2 in the low-frequency region. (f) Normalized imaginary capacitances of different film electrodes (inset is the minimal characteristic relaxation time constant τ0).

the poor electrical conductivity of the film (4.53 S/cm). Therefore, an appropriate number of rGO nanosheets are needed in the hybrid film to achieve the optimum synergistic effect of the porous structure and high conductivity of the scaffold. In our work, the optimal mass ratio of Ti3C2 to GO sheets is 1:1. In addition to the rate performance, the cycle performance is also measured to examine the structural stability of all samples during cycling. Figure 5e displays the cycling performance of Ti3C2, rGO, and various Ti3C2−rGO composite electrodes at 1 A g−1. It clearly shows that the capacity of the 1-1-Ti3C2−rGO electrode increases during cycling; however, the capacity of other electrodes tends to decrease. After 1000 cycles, the 1-1Ti3C2−rGO electrode shows the highest charge capacity of 212.5 mA h g−1, much higher than pure Ti3C2 (56.4 mA h g−1), pure rGO (58.4 mA h g−1), 4-1-Ti3C2−rGO (70.8 mA h g−1), 3-1-Ti3C2−rGO (85.8 mA h g−1), 2-1-Ti3C2−rGO (101.8 mA h g−1), and 1-2-Ti3C2−rGO (137.1 mA h g−1) electrodes. As shown in Figure 5f, the 1-1-Ti3C2−rGO electrode not only shows the highest capacity but also displays the highest capacity retention after cycling. The superior cycling performance can

transfer resistance of the hybrid film as displayed in Table S1. Therefore, the porous structure and electrical conductivity of the hybrid film are strongly dependent on the mass ratio between the two components; and the electrochemical performances of the films are also correlated with their compositions as reflected in Figure 5d. Among these hybrid films, the 1-1-Ti3C2−rGO electrode exhibits the highest specific capacitance of 335.5 mA h g−1 at 0.05 A g−1, slightly higher than those of 4-1-Ti3C2−rGO (303.6 mA h g−1), 3-1-Ti3C2− rGO (310.1 mA h g−1), 2-1-Ti3C2−rGO (308.8 mA h g−1), and 1-2-Ti3C2−rGO (311.7 mA h g−1). It is worth noticing that the capacity retention of the 1-1-Ti3C2−rGO electrode (30%) at 4 A g−1 is also much higher than those of 4-1-Ti3C2−rGO (18.7%), 3-1-Ti3C2−rGO (18.3%), 2-1-Ti3C2−rGO (18.7%), and 1-2-Ti3C2−rGO (24%). The high rate capability of 1-1Ti3C2−rGO should be attributed to the 3D porous microstructure and good conductivity of the Ti3C2−rGO film. Even though the 1-2-Ti3C2−rGO hybrid film (Figure S7d) shows more open microstructures than 1-1-Ti3C2−rGO (Figure S5a), the rate capability is worse than that of 1-1-Ti3C2−rGO, which is probably due to the high amount of rGO sheets, leading to 3640


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Figure 7. (a) Cross-sectional and (b) top-view SEM images of the 1-1-Ti3C2−rGO film electrode after 1000 cycles at 1 A g−1. (c) TEM images of the 1-1-Ti3C2−rGO film electrode after 1000 cycles at 1 A g−1. (d) TEM-EDS results of the selected area marked in panel (c) with white-dashed lines.

accordance with the results (Table S1) tested by the four-point probe. Even though the 1-1-Ti3C2−rGO film electrode owns a medium resistance value among all samples, it is still much smaller than the rGO film electrode, indicating that the hybrid Ti3C2−rGO scaffold provides faster electron transfer channels than the pure graphene scaffold. Furthermore, from the relationship between Z′ and ω−1/2 (ω = 2πf) in the lowfrequency region, as shown in Figure 6e, the slopes of the curves decrease when the mass ratio of GO to Ti3C2 increases, demonstrating the fast ion diffusion/transportation kinetics. Overall, the increasing amount of Ti3C2 in the hybrid film brings about higher conductivity but lower ion diffusion/ transportation kinetics. These results are accordant with our hypothesis. Despite the fact that the rGO component in the hybrid film decreases the electron conductivity, it transforms the dense-stacked MXene sheets to a 3D porous microstructure, which is beneficial for the fast mass transport in the interior of the film electrode. Therefore, a hybrid film electrode displays better electrochemical performance than either pure rGO or pure MXene electrode due to the synergetic effect of the ion transmission rates and electrical conductivity. This can be further confirmed by the minimal characteristic relaxation time constant τ0 (the minimum time needed to discharge all the energy with an efficiency of >50%). As shown in Figure 6f, the 1-1-Ti3C2−rGO film shows a minimal τ0 of 1 s, which is smaller than the pure rGO film (6 s), Ti3C2 film (2 s), 4-1Ti3C2−rGO film (1.5 s), 3-1-Ti3C2−rGO film (2 s), 2-1Ti3C2−rGO film (1.8 s), and 1-2-Ti3C2−rGO film (1.8 s). Therefore, the 1-1-Ti3C2−rGO film has the potential to deliver high power, which is accordant with the experimental results. Furthermore, compared with other MXene-based anode materials, the 3D Ti3C2−rGO hybrid film electrode shows a better capacity retention ratio at high rates (4 A g−1, Table S2), which was rarely reported in the previous literature.18,25,27,30,32,46 The cycling performance comparison between 3D Ti 3 C 2 −rGO and other materials is also summarized in Table S3. Apparently, even at a high current

be ascribed to the synergistic effect of its individual component, in which the hybrid scaffold guarantees a high conductivity, and the 3D porous structure provides fast mass transport access. Therefore, the deep active sites in the film can be utilized gradually during cycling, leading to the increased capacity. This further demonstrates that the mass of rGO nanosheets in the hybrid film should be controlled in an appropriate range to achieve the best performance. The CV curves of the 1-1-Ti3C2−rGO film electrode were also tested from 0.1 to 10 mV s−1 to verify its superior electrochemical performance. Generally, the relationship between current (i) and scan rates (v) obeys the power law i = avb

where a and b are appropriate values. The diffusion-controlled process during the electrochemical reaction corresponds to a b value of 0.5, whereas the nondiffusion-controlled behavior corresponds to a value of 1.37 Thus, the nondiffusion-controlled capacitive effect and diffusion-controlled insertion process can be well-distinguished depending on the value of b. To confirm the charge storage mechanism of the 1-1-Ti3C2−rGO electrode, the plot of log(i) versus log(v) from 0.1 to 10 mV s−1 for the anodic and cathodic scans at 2 V is shown in Figure 6b. The b values of the anodic and cathodic processes are both larger than the value of 0.5, indicating that the charge storage process is predominantly nondiffusion-limited, thus providing superior rate capability. To directly prove the contribution of a novel 3D porous microstructure in the fast reaction kinetics and high capacity of the 1-1-Ti3C2−rGO electrode, the EIS (Figure 6c) was tested to quantify the electron-transfer resistances of all film electrodes and evaluate their ion transfer rates in the film electrodes. After fitting the EIS spectra by the equivalent circuit (Figure S9a), the electron transfer resistances of all samples were calculated and are shown in Figure 6d. Obviously, the Ti3C2 film has the least electron transfer resistance, and the resistance rises with the increase of the mass ratio of GO to Ti3C2, which is in 3641


ACS Applied Materials & Interfaces density of 1 A g−1, the hybrid film electrode still exhibits excellent cycling stability. The good electrochemical performance of the 3D Ti3C2−rGO hybrid electrode should be ascribed to the following: (i) the Ti3C2 MXene sheets maintain inherent high conductivity, which guarantees superior electron conductivity of the film electrode and (ii) the 3D porous structure prevents the dense stacking of MXene sheets and also facilitates the infiltration of electrolytes in the film, leading to better ion diffusion/transportation kinetics. Because of the synergistic effect of superior electronic and ionic conductivity, the hybrid film electrode displays excellent electrochemical performance especially at high rates. To verify the structural stability of the hybrid film, the microstructure of the 1-1-Ti3C2−rGO film after 1000 cycles at 1 A g−1 was characterized. After cycling, the 3D porous structure (Figure 7a) of the film is still maintained and the top surface of the film is smooth (Figure 7b), but some small particles can be detected in the high-resolution SEM (inset, Figure 7b) of the film due to the formation of an SEI layer on the sheets. The TEM image (Figure 7c) also shows that the scaffold of the film still maintains the original structure, in which the Ti3C2 and graphene nanosheets distribute evenly on the hybrid film. Furthermore, the TEM−EDS results (Figure 7d) show that the content of Ti and C is 30.37, 47.76% respectively, indicating the co-existence of MXene and rGO sheets in the hybrid film.

ACKNOWLEDGMENTS

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b17386.

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This work was supported by National Natural Science Foundation of China (grant no. 21371176), Key Research Program of the Chinese Academy of Sciences (grant no. KGZD-EW-T08), Ningbo Key Science and Technology Projects (grant no. 2014S10008), and Ningbo Science and Technology Innovation Team.

4. CONCLUSIONS In summary, we have demonstrated a simple method to prepare 3D porous MXene−rGO hybrid films through electrolyteassisted self-assembly, vacuum filtration, and moderate postprocessing. The addition of GO transforms the densestacked MXene sheets to a 3D porous structure, which effectively facilitates the diffusion of electrolyte ions in the film electrode, whereas the MXene sheets endow the hybrid film with high electrical conductivity. Because of the synergistic effect of the fast electronic and ionic conductivity, the hybrid film electrode exhibits outstanding rate capability and excellent cycle performance in LIBs. Therefore, we firmly believe that this work may provide a new strategy to solve the issue of poor ionic transport in MXene-based film electrodes, further inspiring the energy storage potential of MXene materials.

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Research Article

XRD, TGA, additional SEM images, and electrochemical testing data (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (X.Z.). *E-mail: [email protected]. Phone/Fax: +86-574-8668-5096 (Z.L.). ORCID

Xufeng Zhou: 0000-0002-3153-6954 Notes

The authors declare no competing financial interest. 3642


Research Article


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