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Cellulose Nanofiber @ Conductive Metal− Organic Frameworks for High-Performance Flexible Supercapacitors Shengyang Zhou,†,§ Xueying Kong,‡,§ Bing Zheng,‡ Fengwei Huo,‡ Maria Strømme,*,† and Chao Xu*,† Downloaded via BUFFALO STATE on July 17, 2019 at 12:29:17 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.
†
Nanotechnology and Functional Materials, Department of Engineering Sciences, Ångström Laboratory, Uppsala University, 751 21 Uppsala, Sweden ‡ Key Laboratory of Flexible Electronics (KLOFE), Institute of Advanced Materials (IAM), Nanjing Tech University (NanjingTech), 211816 Nanjing, China S Supporting Information *
ABSTRACT: Conductive metal−organic frameworks (c-MOFs) show great potential in electrochemical energy storage thanks to their high electrical conductivity and highly accessible surface areas. However, there are significant challenges in processing c-MOFs for practical applications. Here, we report on the fabrication of c-MOF nanolayers on cellulose nanofibers (CNFs) with formation of nanofibrillar CNF@cMOF by interfacial synthesis, in which CNFs serve as substrates for growth of c-MOF nanolayers. The obtained hybrid nanofibers of CNF@c-MOF can be easily assembled into freestanding nanopapers, demonstrating high electrical conductivity of up to 100 S cm−1, hierarchical micromesoporosity, and excellent mechanical properties. Given these advantages, the nanopapers are tested as electrodes in a flexible and foldable supercapacitor. The high conductivity and hierarchical porous structure of the electrodes endow fast charge transfer and efficient electrolyte transport, respectively. Furthermore, the assembled supercapacitor shows extremely high cycle stability with capacitance retentions of >99% after 10000 continuous charge− discharge cycles. This work provides a pathway to develop flexible energy storage devices based on sustainable cellulose and MOFs. KEYWORDS: conductive metal−organic frameworks, cellulose nanofibers, interfacial synthesis, flexible nanopaper electrodes, flexible energy storage devices
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effective ways to solve these problems that would boost the applications of MOFs in EES. Currently, there are two common approaches to overcome the drawback of insufficient conductivity of traditional MOFs. One is that of pyrolyzing MOFs into porous carbons or carbon− metal/metal oxide composites where MOFs serve both as precursors and templates.14−16 Although such MOF-derived materials have increased conductivity and inherited porosity, the intrinsic functionalities of pristine MOFs are sacrificed. Another approach is adding conductive materials (e.g., conducting polymers, carbon nanotubes, graphene) onto the surface of MOF particles or into their pores.17−20 However, the additives would block the interconnected porous channels and reduce the effective surface area of MOFs. Recently, a number of conductive MOFs (c-MOFs) have been reported by linking fully conjugated organic linkers with d7−9 metal centers (e.g.,
etal−organic frameworks (MOFs) are a class of crystalline and porous materials that link inorganic clusters and organic ligands via coordination bonds.1−3 By judicious selection of inorganic and organic components, MOFs with predetermined architectures and tunable properties can be designed by molecular engineering.4,5 As a result, a variety of MOFs have been synthesized for studies in gas storage and separation, drug delivery, catalysis, etc.6−8 In addition to these well-established applications, many efforts have been devoted into the exploitation of their uses in electrochemical energy storage (EES) due to their highly accessible surface areas, tunable pore sizes, built-in redox-active metal centers, and versatile functionalities.9−13 However, the use of traditional MOFs in EES faces two main challenges: (1) the intrinsic low electrical conductivity of MOFs impedes charge transfer in the framework and thus significantly limits their electrochemical performances (2) the difficulty in processing MOF crystals hampers their practical applications, especially in flexible EES devices. Therefore, it is highly desirable to create © XXXX American Chemical Society
Received: June 14, 2019 Accepted: July 11, 2019 Published: July 11, 2019 A
DOI: 10.1021/acsnano.9b04670 ACS Nano XXXX, XXX, XXX−XXX
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Figure 1. (a) Schematic of synthesis procedure for CNF@c-MOF hybrid nanofibers. (b) XRD patterns of pure CNFs, pure Ni−HHTP and Ni− HITP powder, CNF@Ni−HHTP, and CNF@Ni−HITP hybrid nanofibers. (c) TEM image of a single CNF@Ni−HITP hybrid nanofiber, the inset shows the high resolution TEM image collected from the interface area of CNF@Ni−HITP nanofiber. (d) High-resolution Ni 2p3/2 X-ray photoelectron spectra of CNF@Ni−HITP, pure Ni−HITP, and CNF−Ni (Ni2+ ion-exchanged CNF). (e) High-resolution SEM image of CNF@ Ni−HITP hybrid nanofibers. (f) SEM images of CNF@Ni−HITP nanopaper at different magnifications. (g) Photograph of an origami folded by CNF@Ni−HITP nanopaper.
Co(II), Ni(II), Cu(II)).21−27 The highly π-conjugated organic ligands and the strong metal−ligand interactions lead to full charge delocalization in the 2D plane, while the stacked πconjugated nanostructures facilitate charge hopping between the layers, which result in high electrical conductivity of up to 1580 S cm−1 for the c-MOFs.28−31 Therefore, pristine c-MOFs with high conductivity, high surface areas, and redox-active centers are ideal electrode materials for supercapacitors and batteries, providing high volumetric and gravimetric capacitance or capacity.23−26 It is noteworthy that the bulk conductivity of cMOFs is largely dependent on their physical forms. For example, the reported c-MOF Ni−HITP (Ni3(HITP)2, HITP = 2,3,6,7,10,11-hexaiminotriphenylene) has much lower bulk conductivity in the form of a pressed pellet compared to its thin film.22,32 The reduced conductivity can be attributed to the rich grain boundaries, large contact resistance, and random orientation of MOF microcrystals within the pellet. In this context, formation of c-MOFs with controlled morphology and alignment in the form of nanofibers or nanosheets is of great interest for achieving a significant increase of their bulk
conductivity and thus enhancement of their electrochemical performances. However, the brittle and insoluble nature of MOF crystals obtained from conventional synthesis makes it difficult to process MOFs into desired nanomaterials with controlled morphology and alignment. To date, a few studies report the fabrication of MOF nanowires or nanosheets on certain substrates (e.g., graphene oxide, MoS2 nanosheet, Cu(OH)2, track-etched polycarbonate).33−36 Nonetheless, these approaches are largely dependent on the surface and structure of the substrates and not suitable for the fabrication of all types of MOFs. Meanwhile, the low processability of MOFs complicates the formation of MOF-based flexible materials. Although depositing MOF nanoparticles on flexible substrates (e.g., carbon cloths, polymers, textiles) could endow a certain flexibility for the composites, the weak interfacial compatibility between MOFs and the substrates would cause aggregation of MOF nanoparticles and thus reduce their accessible porosity.17,37,38 Very recently, we have developed a sustainable and scalable approach to assemble a range of MOFs with the B
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Figure 2. (a) N2 adsorption and desorption isotherms and pore size distributions of pure CNF, CNF@Ni−HHTP, and CNF@Ni−HITP nanopapers. (b) Comparison of electrical conductivity of CNF@c-MOF nanopapers, pure c-MOFs pellets, and CNF-c-MOFs papers. (c) Normalized resistance of CNF@Ni−HITP nanopaper at different bending states. (d) Schematic diagram showing the charge transfer and electrolyte ion transport in the nanofibrous and conductive networks of CNF@c-MOF nanopapers.
The composition of CNF@c-MOF was determined by powder X-ray diffraction (XRD). The diffraction patterns of CNF@c-MOFs are compared with pure CNF and c-MOF powders, as illustrated in Figure 1b. The diffraction peaks at 2θ ≈ 13.1, 15.2, 20.2° indicate the (100), (010), and (110) reflections of CNF, respectively.39 The typical peaks at 2θ ≈ 4.7 and 9.5° correspond to the (100) and (200) planes of the cMOFs (Ni−HITP and Ni−HHTP), respectively.21,22 Additionally, the broader peaks at 2θ ≈ 27.6°, attributed to the (001) reflections of c-MOFs, suggest a long-range order along the c direction. It is noteworthy that the diffraction peaks at 2θ ≈ 27.3° for Ni−HHTP and Ni−HITP powders slightly shift to 27.4 and 27.6° as observed in CNF@Ni−HHTP and CNF@ Ni−HITP, respectively, indicating a reduced d-spacing of (001) planes for the c-MOFs grown on CNFs. It could be speculated that the c-MOF nanolayers on CNFs have stronger π−π stacking interactions and probably higher electrical conductivities than those of c-MOF powders. Figure 1c shows the transmission electron microscope (TEM) image of a CNF@Ni−HITP hybrid nanofiber. In contrast to the TEM images of pure CNFs (Figure S1), a typical core−shell structure is clearly observed, in which continuous nanolayers of Ni−HITP compactly wraps the CNFs. A clear contrast is observed at the interface between the core and the shell as seen from the high-resolution TEM image (inset in Figure 1c). The honeycomb structure of the porous framework of Ni−HITP along the plane (001) reveals a hexagonal unit cell agreeing well with the results observed for pure Ni−HITP.32 The energydispersive X-ray (EDX) spectroscopy mapping of the nanofiber confirmed the uniform distribution of N and Ni in the Ni−HITP nanolayer (Figure S4). Furthermore, we performed X-ray photoelectron spectroscopy (XPS) analysis to investigate the binding mode between the CNF and the Ni−HITP nanolayer. The binding energy (BE) of Ni 2p3/2 in CNF−Ni (Ni2+ ionexchanged CNF) was 856.4 eV, similar to the value reported for
assistance of cellulose nanofibers (CNFs) into freestanding, flexible nanopapers, in which continuous MOF nanolayers were grown on the surface of CNFs by interfacial synthesis.39 With this strategy in mind, we envisioned that fabrication of continuous nucleated c-MOF nanolayers on CNFs would not only increase the electrical conductivity of the nanocomposites by reducing the grain boundaries and contact resistance of the cMOFs but also warrant high flexibility of the nanocomposites. Herein, we report the nanoengineering of c-MOFs on the surface of carboxylated CNFs with the formation of highly conductive, hierarchical porous, flexible nanopapers of CNF@cMOF as freestanding and additive-free electrodes for highperformance supercapacitors.
RESULTS AND DISCUSSION Figure 1a provides a schematic of the interfacial synthesis procedures for CNF@c-MOF hybrid nanofibers. CNFs extracted from Cladophora algae were first treated by TEMPO oxidation (TEMPO = 2,2,6,6-tetramethylpiperidin-1-yloxyl) to introduce carboxyls on the surface (Figure S1). Thereafter, the carboxylated CNFs were thoroughly ion-exchanged with Ni2+ ions. Subsequently, an aqueous solution containing Ni(NO3)2· 6H2O and organic ligand HHTP (HHTP = 2,3,6,7,10,11hexahydroxytriphenylene) or HITP was added to the ionexchanged CNFs suspension for the construction of Ni−HHTP and Ni−HITP, respectively. Ultimately, a homogeneous suspension of CNF@c-MOF nanofibers was obtained. Furthermore, freestanding CNF@c-MOF nanopapers can be formed on a membrane filter (pore size: 0.1 μm) by vacuum filtration of the suspension, indicating the good processability of the nanofibrous materials (Figure S2). The contents of c-MOFs in CNF@c-MOF nanopapers were determined to be ∼15 wt % by thermogravimetric analysis (TGA) (Figure S3). More detailed synthesis procedures are provided in the Methods. C
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Figure 3. Electrochemical performances of CNF@c-MOF nanopaper electrodes in a three-electrode setup using 3 M KCl aqueous solution as electrolyte. (a) Photograph of a three-electrode setup used in this work. (b) Cyclic voltammetry curves of CNF@Ni−HITP at different scan rates. (c) Galvanostatic charge and discharge curves of CNF@Ni−HITP at different current densities. (d) Gravimetric capacitance of the CNF@c-MOF electrodes. (e) Comparison of gravimetric capacitances and (f) electrochemical impedance spectra of the CNF@c-MOF and CNF-c-MOF electrodes.
nickel(II) acetate (Figure 1d, Figure S5).40 In comparison, the BE for the Ni(II) species in CNF@Ni−HITP shifted negatively by ∼0.7 V upon growth of Ni−HITP nanolayers on CNFs. The Ni 2p3/2 spectrum of CNF@Ni−HITP can be deconvoluted into two peaks at 855.6 and 856.2 eV, indicating the presence of different coordination environments for Ni(II) species. The peak at 855.6 eV can be attributed to the Ni(II) species in the Ni−HITP layer, in which Ni(II) species were surrounded by four amine groups. The peak at 856.2 eV can be assigned to the Ni(II) species located at the interface joining the CNF and the Ni−HITP layers. These results strongly evidenced the interfacial growth of a Ni−HITP nanolayer on the surface of the CNF during which the Ni(II) species in CNF−Ni acted as bridges to bond CNF and Ni−HITP resulting in the integrated nanofibers of CNF@Ni−HITP. Meanwhile, scanning electron microscope (SEM) images of CNF@c-MOFs further confirm the nanofibrous morphology with CNFs entirely wrapped by continuous c-MOF nanolayers (Figure 1e and Figures S6 and S7). In addition, the microstructures of the assembled CNF@c-MOF nanopapers are revealed by their SEM images recorded at different magnifications (Figure 1f and Figure S6). The hybrid nanofibers interweave compactly with each other, forming a uniform network (Figure 1f and Figure S7). No individual c-MOF particles can be identified. Consequently, the typical nanostructure results in strong mechanical strength for the CNF@c-MOF nanopapers with tensile stresses and Young’s modulus of up to 350 MPa and 10 GPa, respectively (Figure S8). Meanwhile, the CNF@c-MOF nanopapers demonstrate good flexibility and foldability and can be folded into complex origami structures (Figure 1g, video S1). The high processability of CNF@c-MOF nanofibers, good flexibility, and strong mechanical strength of the fabricated nanopapers are of great importance for the development of their practical applications. The porosities of CNF@c-MOF nanopapers are evaluated by nitrogen sorption measurements (Figure 2a). The Brunauer−
Emmett−Teller surface areas of CNF@Ni−HHTP and CNF@ Ni−HITP were calculated to be 203 and 195 m2 g−1, respectively, which are significantly higher than that of pure CNFs (96 m2 g−1) due to the contribution of microporous cMOFs. The pore size distribution analyses reveal the hierarchical porous structure of CNF@c-MOF nanopapers, which contain both micropores (∼1.5 nm) and mesopores (15 nm−50 nm) arising from the c-MOFs and from interfiber cavities, respectively. The electrical conductivities of CNF@cMOF nanopapers, pure c-MOF pellets (pressed by c-MOF powders at a pressure of 1.5 GPa) (Figure S9), and CNF-c-MOF papers (prepared by direct mixing of c-MOF powders and CNFs in a weight ratio of 15:85, Figure S10) were measured by a standard four-probe method (Figure 2b, Figure S11). Noteworthy, the measured conductivities of the CNF@c-MOF nanopapers (9.4 S cm−1 for CNF@Ni−HHTP and 103 S cm−1 for CNF@Ni−HITP) are ∼1−2 orders of magnitude higher than the values of the corresponding pure c-MOF pellets (0.26 S cm−1 for Ni−HHTP and 39 S cm−1 for Ni−HITP) and ∼3−4 orders of magnitude higher than the values of the CNF-c-MOF papers (0.0023 S cm−1 for CNF-Ni−HHTP, 0.19 S cm−1 for CNF-Ni−HITP). The conductivity of the CNF@Ni−HITP nanopaper is even larger than most of the reported cellulosebased conductive papers (Table S1) and MOF-based conductive composites (Table S2). The significant improvement in conductivity for CNF@c-MOF nanopapers can be correlated to the continuous nucleated c-MOF nanolayers on CNFs that create fewer grain boundaries and smaller interparticle spaces compared to c-MOF powders. In addition, the flexible CNF@cMOF nanopapers demonstrate a very stable conductivity; the resistances are almost unchanged at various bending and folding angles (Figure 2c, Figure S12). We, hence, anticipated that the CNF@c-MOF nanopapers can be used as effective electrodes: the hierarchical porosity and high electrical conductivity are expected to facilitate electrolyte transport and charge transfer, respectively (Figure 2d). D
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MOF nanopapers are constructed by interconnected conductive nanofibers with continuous c-MOF layers compactly wrapped on CNFs. This nanostructure could most likely promote charge transfer within the electrode and at the interface of the electrode/electrolyte. In contrast, CNF-c-MOF papers consist of c-MOF particles dispersed in the CNFs matrix, where electron transfer is significantly limited between the individual cMOF particles suppressing their electrochemical performances. A detailed analysis of the EIS spectra strongly supports this hypothesis. As is clear from Figures 3f and S21, the EIS responses of all samples correspond well to that generated by a Randles equivalent circuit44 describing charge transfer and diffusion of the electroactive species with a finite-length Warburg diffusion element instead of an ordinary semi-infinite Warburg diffusion element (see the Supporting Information for further details).45 Such a circuit has been found to be able to model the response of a multitude of intercalation electrodes including different oxides45−47 and polypyrrole-covered nanocellulose.48 Table S3 summarizes the high-frequency resistance (high frequency intercept with real axis), the charge-transfer resistance (diameter of high frequency semicircle), as well as the frequencies for which diffusion (end of semicircle and start of 45° angled line) and finite-length diffusion (capacitive behavior; end of 45° line and start of line with steeper slope) start to dominate the EIS spectra. From this table, it is clear that both the high-frequency resistance of the electrodes and the chargetransfer resistance at the electrode/electrolyte interface are significantly lower for the CNF@c-MOF nanopapers than those for the corresponding CNF-c-MOF nanopapers. The fact that the onset frequencies for a diffusion-limited response as well as for a finite-length Warburg diffusion response are considerably higher for the CNF@c-MOF than those for the CNF-c-MOF counterparts shows that diffusion proceeds more rapidly in the former in which electrolyte ions are able to easily diffuse in the electrodes through their hierarchical pores. In contrast, the diffusion of electrolyte in the CNF-c-MOF papers is hampered because the interfiber pores of the CNF structure are mostly blocked by c-MOF particles as observed in the SEM images (Figure S10). Hence, we conclude that the nanostructure of the CNF@c-MOF nanopapers, as described above, enabling efficient electron transfer and fast electrolyte ion transport is of great benefit for their electrochemical performances. The observant reader may recall that the difference in conductivity between the CNF@c-MOF nanopapers and their CNF-c-MOF counterparts, as presented in Figure 2b, was ∼1000−10000 times (3−4 orders of magnitude). The corresponding differences between high-frequency resistances and charge-transfer resistances are ∼2−3 times and ∼6−16000 times, respectively. Since the conductivities are extracted under dc conditions while the high-frequency resistances and the charge-transfer resistances are extracted at high frequencies, it makes no sense to compare these values. Given the high flexibility, mechanical strength, and electrical conductivity as well as the excellent electrochemical performance, we assembled a symmetric supercapacitor based on CNF@Ni−HITP nanopapers (1.5 cm × 2.0 cm × 0.005 cm) for the evaluation of their potential applications in flexible electrochemical supercapacitor devices. A piece of filter paper, working as the separator, was sandwiched between two identical CNF@Ni−HITP nanopaper electrodes. Two pieces of graphite foil were used as current collectors. A gel prepared by dissolving poly(vinyl alcohol) (PVA) in aqueous KCl was used as the
The electrochemical performances of CNF@Ni−HITP and CNF@Ni−HHTP nanopapers were evaluated in an asymmetric three-electrode setup in an aqueous electrolyte (3 M KCl). The freestanding CNF@c-MOF nanopapers attached on a platinum ring served as working electrodes without any extra binders or conductive additives. The areal loading densities of c-MOFs on CNF@Ni−HITP and CNF@Ni−HHTP nanopapers were 0.578 and 0.714 mg cm−2, respectively. An Ag/AgCl electrode and a platinum wire were used as the reference electrode and counter electrode, respectively (Figure 3a). The CNF@c-MOF nanopaper electrodes display rectangular cyclic voltammetry (CV) curves at scan rates of 5−200 mV s−1 (Figure 3b) and highly symmetric triangular galvanostatic charge/discharge (GCD) curves (Figure 3c) at current densities of 0.33−33.33 A g−1 between the open-circuit potential of 0−0.7 V, which are characteristics of capacitive behaviors (Figure S13).23,24 Consistently, the electrode displayed similar CV curves when cycled catholically between 0 and −0.7 V. The CV studies revealed relatively broad working potential windows of up to 1.4 V for CNF@Ni−HITP (Figure S14). According to the GCD studies (Figure 3c), the gravimetric capacitances of the Ni− HITP and Ni−HHTP were calculated to be 125 F g−1 (current density: 0.33 A g−1) and 75 F g−1 (current density: 0.2 A g−1), respectively (Figure 3d), which are significantly higher than the corresponding values of most pristine MOFs.41−43 Notably, both electrodes show high rate capability. For example, the CNF@Ni−HITP nanopaper electrode retains ∼70% of its capacitance when the current density increases 100 times from 0.33 to 33 A g−1 (Figure 3d). In addition, we could tune the thickness of the c-MOF nanolayers on the CNFs during the synthesis and adjust the content of the c-MOF in the electrode. The CNF@Ni−HITP nanopapers consisting of thicker nanofibers have higher conductivity and retain good flexibility. More importantly, the increased thickness of the Ni−HITP nanolayers on the CNFs has no significant influence on the electrochemical performance including gravimetric capacitance, rate capability, and Coulombic efficiency of the nanopaper electrode, leading to a higher areal capacitance for the electrode (Figure S16). Furthermore, the electrochemical performances of CNF-c-MOF papers, prepared by mixing of the c-MOF particles and CNFs, were evaluated for comparison studies that may reveal the structure−property relationships of the electrodes. The CNF-c-MOF paper electrodes show significantly different electrochemical responses compared to those of the CNF@cMOF electrodes. The CNF-Ni−HITP electrode displays shuttle-shaped CV curves at low scan rates and very weak responses at high scan rates, while the CNF-Ni−HHTP paper did not show any significant electrochemical activity (Figure S17). Figure 3e compares the gravimetric capacitance of the cMOFs in the CNF@c-MOF nanopaper and the CNF-c-MOF paper electrodes. The values were calculated from the CV curves at the same scan rate of 50 mV s−1 (Figure S18 and S19). Noticeably, the capacitances of the c-MOFs in the CNF@cMOF nanopaper electrodes (Ni−HITP: 103 F g−1, Ni−HHTP: 39 F g−1) are >10 times higher than those of c-MOFs in the CNF-c-MOF paper electrodes (Ni−HITP: 8 F g−1, Ni−HHTP: 0.04 F g−1). To gain insight into the reason for the large difference in the capacitive performance of CNF@c-MOF nanopapers and CNFc-MOF papers, which are made up of exactly the same components, we compared their nanostructures in SEM images (Figure S10) and the results from electrochemical impedance spectroscopy (EIS) studies (Figure 3f). As mentioned, CNF@cE
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Figure 4. Electrochemical performances of the CNF@Ni−HITP double-layer supercapacitor. (a) Cyclic voltammetry curves at different potential windows at the same scan rate of 10 mV s−1. (b) Galvanostatic charge and discharge curves at a current density of 0.6 A g−1 at increasing potential window. (c)Calculated areal capacitances of the device at different current densities within 0−0.7 V (blue curve) and 0−1.0 V (orange curve). (d) Cyclic performance and capacitance retention data of the device within 0−0.7 V (blue curve) and 0−1.0 V (orange curve). (e) Cyclic voltammetry curves (scan rate: 100 mV s−1) under different folding angles. (f) Photograph of LED powered by the devices in series under different deformations.
increasing the current density from 0.2 to 13 mA cm−2, which gives higher rate capability than those of devices assembled by cMOF powders.23,24,26 The electrochemical cycle performance of the device was investigated by GCD cycling experiments at 1 A g−1. The device showed outstanding stability with no obvious reduction in capacitance after cycling 10000 times between 0 and 0.7 V, while only a 10% reduction in capacitance was observed when cycled between 0 and 1.0 V (Figure 4d). Consistently, the CNF@Ni−HITP nanopaper electrode possessed high structural stability as the crystalline structure and nanomorphology were unchanged during the cycling measurements (Figure S25). The high rate capability and stability of the device can be attributed to the integrated nanofibrous structure of CNF@Ni−HITP as mentioned above. In comparison, the device showed slightly higher capacitances at a wider potential window of 0−1.4 V (Figure S22), which resulted in significantly increased energy density and power density (Figure S24). However, the charge−discharge cyclic stability (Figure S26) and rate capability (Figure S22d) were sacrificed by extending the potential windows, which probably can be attributed to the minor faradaic process as revealed from the CV studies (Figure 4a). We anticipated that device optimization by, for instance, replacing the negative electrode with carbon-based materials, using organic or ionic liquid electrolytes, would further extend the working voltage and increase the cycling stability of the device. In addition to high capacitance and stability, mechanical flexibility is essential for energy-storage devices to power flexible electronics, such as foldable phones, curved displays, and wearable devices. Previous studies showed that pressing c-MOF powders into dense and thick pellets could dramatically increase the areal and volumetric capacitances of the devices.23,26 However, such devices are rigid and not optimal for use in flexible electronics. In comparison, the CNF@Ni−HITP supercapacitor developed in this study shows excellent flexibility. Bending (90, 120°) or even folding (180°) the device
electrolyte. Obviously, the use of aqueous electrolyte offers significant advantages of high rate capability, safety, and environmentally friendliness compared to the organic electrolyte. As expected, the device showed typical double-layer capacitive behaviors as illustrated by the CV and GCD curves (Figure 4a, b, Figure S22). The CV curves of the symmetric device recorded within the potential windows from 0−0.4 to 0− 1.4 V at a scan rate of 10 mV s−1 showed nearly rectangular shapes. It should be noted that a pair of broad and inconspicuous redox peaks at a cell voltage of 1.0 and 0.6 V were observed when operating the device at the potential windows of 0−1.2 and 0− 1.4 V, suggesting the presence of minor faradaic processes. Consistently, the GCD curves showed triangular traces within the potential windows from 0−0.4 to 0−1.4 V (Figure S22). However, the device had significantly decreased Coulombic efficiency when the potential window was extended above 1.0 V (Figure S23). To evaluate the effects of working potential window on the electrochemical performances, we prepared three identical symmetric devices for comparison studies in varying potential windows (0−0.7, 0−1.0, and 0−1.4 V). As illustrated in Figure 4c, the device showed relatively high areal capacitance of 96 mF cm−2 (current density: 0.2 mA cm−2) within 0−1.0 V. A high gravimetric capacitances of 141.5 F g−1 could be reached at a current density of 0.075 A g−1 for the device (Figure S22). As shown in Figure 4c, the overall device including working electrodes, separator, and current collectors is very thin (a thickness of 0.35 mm), giving a relatively high volumetric capacitance of up to 2800 mF cm−3. Accordingly, the maximum energy and power density of the device were 6.5 mW h cm−2 (185.7 mW h cm−3) and 0.013 mW cm−2 (0.37 mW cm−3), respectively (Figure S24). Noteworthy, the calculated values of capacitance, energy, and power density for the assembled device are comparable to the reported values for state-of-the-art MOFbased supercapacitors.17,20,49,50 In addition, it retained over 50% of its capacitance (59.6% for 0−0.7 V, 50.1% for 0−1.0 V) on F
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AT02) at an accelerating voltage of 200 kV. SEM images were collected on an FEG SEM instrument (Zeiss, Leo Gemini 1530) at an accelerating voltage of 3 kV. TGA was performed on a thermogravimetric analyzer (Mettler Toledo, /SDTA851e) under N2 or air flow (60 mL min−1) between 25 and 900 °C with a heating rate of 10 °C min−1. N2 sorption isotherms were measured in a Micromeritics ASAP 2020 surface area and pore size analyzer at 77 K. The samples were degassed at 100 °C under a kinetic vacuum (
