Bacterial

Research Article pubs.acs.org/journal/ascecg

Hierarchical Porous Carbons from Poly(methyl methacrylate)/ Bacterial Cellulose Composite Monolith for High-Performance Supercapacitor Electrodes Qiuhong Bai,† Qiancheng Xiong,† Cong Li,† Yehua Shen,*,† and Hiroshi Uyama*,†,‡ †

Key Laboratory of Synthetic and Natural Functional Molecule Chemistry of Ministry of Education, College of Chemistry and Materials Science, Northwest University, No. 1 Xuefu Avenue, Xi’an City, Shanxi Province 710127, China ‡ Department of Applied Chemistry, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita 565-0871, Japan S Supporting Information *

ABSTRACT: This study deals with hierarchical porous carbons from bacterial cellulose (BC), having a layered structure for high-performance application, such as supercapacitor electrodes, fabricated from a composite monolith with unique microscopic/macroscopic morphology. A poly(methyl methacrylate) (PMMA)/BC composite monolith was first synthesized by thermally induced phase separation using ethanol and deionized water as solvents, where BC acts as the main carbon source as well as matrix and PMMA acts as the activator source producing the necessary activation material. Scanning electron microscopy analysis showed that a monolithic skeleton of PMMA was loaded uniformly on the nanofibers of BC to form a three-dimensional entangled structure of the PMMA skeleton and BC nanofibers, as observed in the microscopic view. Furthermore, the macroscopic twodimensional layered structure of BC remained in the as-obtained composite. The specific surface area, structural features, and thermal stability were investigated by Brunauer−Emmett−Teller, X-ray diffraction, and thermogravimetric analysis studies. The resulting PMMA/BC composite was carbonized and activated by KOH at 850 °C. The electrochemical properties were characterized by cyclic voltammetry, galvanostatic charge−discharge, and electrochemical impedance spectroscopy showing that the carbonization product of the composite displayed a high specific capacitance of 266 F g−1 at a current density of 0.50 A g−1 and the energy density reached a maximum of 23.6 W h kg−1 at a power density of 200 W kg−1. Moreover, 95% of the capacitance was retained after 10,000 charge−discharge cycles, which implies exceptionally high cyclic stability. This compatible and excellent electrochemical performance of the composite, in terms of the energy density and capacitance retention, can be contributed to the characteristic porous structure of the precursor composite monolith. The present research delineates a new approach to fabricate high-performance supercapacitor materials and low-cost energy storage devices from inexpensive bioresources. KEYWORDS: Bacterial cellulose, Composite, Monolith, Poly(methyl methacrylate), Supercapacitor

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INTRODUCTION Electrochemical capacitors (ECs), also called as supercapacitors or ultracapacitors, along with batteries and fuel cells, are considered the most important electrochemical energy storage/ conversion equipment.1 Recently, ECs have attracted significant attention from academia and industry due to their superior power density (103−104 W kg−1), long cycling life (more than 100,000 cycles),2−5 fast charge/discharge rates (in seconds), and a perfect bridge function between the dielectric capacitors and batteries/fuel cells.6,7 Such excellent advantages have triggered interest from scientists and engineers for applications in which consumer portable devices, energy back-up systems, and electrical/hybrid electric vehicles offer the necessary accelerated power and recuperate brake energy.8−10 Depending on their different energy storage mechanisms, ECs can be © 2017 American Chemical Society

briefly classified as pseudocapacitors that store charge mainly by relying on fast and reversible Faradaic redox reactions, electrical double-layer capacitors (EDLCs) that use carbon-active electrode materials based on charge separation at electrode− electrolyte interface, and hybrid-capacitors.11 The EDLCs constitute a major chunk of commercial ECs because of their well-developed features. However, there are some universal disadvantages in the present ECs, which include a limited specific capacitance, low energy density, poor cycle life, and tedious electrode preparation processes involving high cost and complexity. The most intensive measures to address these issues are Received: July 22, 2017 Published: August 14, 2017 9390

DOI: 10.1021/acssuschemeng.7b02488 ACS Sustainable Chem. Eng. 2017, 5, 9390−9401

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ACS Sustainable Chemistry & Engineering

site was reported by a solution casting method.37 The preparation of a PMMA-reduced graphene oxide composite with high conductivity by the self-assembly of positively charged PMMA latex particles and negatively charged graphene oxide sheets through electrostatic interaction was also demonstrated.38 Porous polymeric monoliths with continuous macroporous structures and large surface areas can meet the demands of growing applications. We have achieved facile fabrication of porous monoliths from polymer solutions using thermally induced phase separation (TIPS). Monoliths of PMMA, poly(acrylonitrile), poly(vinyl alcohol), poly(ethylene-co-vinyl alcohol), and poly(γ-glutamic acid) were successfully prepared by selecting an appropriate mixed solvent.39−43 Additionally, organic−inorganic hybrid monoliths were also prepared by TIPS, which were applied for battery materials, adsorbents, separation media, and catalyst carriers.44−47 Herein, we report a PMMA/BC composite monolith with a unique porous structure, in which the BC nanofibers were intertwined with the PMMA monolith skeleton (Figure 1).

creation of novel porous carbon electrode materials with high specific surface area and packing density.12 Thus, the development of simultaneously sustainable, low cost, clean, renewable energy sources, and their related technologies is a priority. So far, conventional activated carbon electrode materials from biomass have been widely used due to their chemical inertness and good thermal and mechanical stability.13,14 However, the activated carbon materials often face electrode kinetic problems, resulting in a low energy of only 4−5 W h kg−15 and poor charge−discharge rates.15 Comparatively, activated carbon-containing composite materials as the carbon material precursor are expected to possess a much higher energy density and a greatly improved electrochemical stability.16−21 Considering the electrochemical properties required, and the EC fabrication cost, the suitable carbon material is a key factor in constructing composites. Bacterial cellulose (BC), a microbial polysaccharide, is produced by a bottom-up process in a bacteria cultivation experiment via biosynthesis by a class of acetic acid producing bacteria such as Acetobacter xylinus.22,23 BC is different from plant cellulose and has many remarkable properties such as a micro- and nanoporous network structure, high purity, crystallinity (of 70−80%) (BC does not contain lignin and hemicelluloses24), and excellent biocompatibility, which is the key to its effective utilization in biomedical research and other relevant research fields.23 The mechanical properties of BC are superior when compared to other fibers. It was reported that the Young’s modulus and tensile strength of BC can reach a maximum of 20.8 GPa and 357.3 MPa, respectively.25 Interestingly, BC is known to have a unique anisotropic morphology on the basis of the static culture condition; the scanning electron microscopy (SEM) image of a freeze-dried BC shows a clear layered structure in the side view and a network structure with ribbon-shaped ultrafine nanofibers with diameters less than 100 nm in the top view26,27 (Figure S1). Furthermore, BC is expected to be a conductive material for various applications,26 especially for those based on carbon.28−31 Here, we used a BC hydrogel with 99% water content for synthesizing high-performance carbon for ECs. BC is an attractive material for electrode precursors because it is an abundant, eco-friendly, and reproducible carbon source and possesses a unique three-dimensional structure. Under the currently selected conditions, neither shrinkage nor deformation took place in the solvent exchange process, which is useful for the preparation of BC-containing composites with a precise structure, for electrode applications. It is well known that poly(methyl methacrylate) (PMMA) is a typical example of transparent polymers with good processabitity. However, its mechanical properties and thermal stability are often insufficient for battery applications.32 Neither does its low conductivity help. An alternative method is to incorporate PMMA into a conductive carbon matrix. Cellulose and its derivatives are often used as the matrix for composites due to their special fibrous network structure.33−35 They are also good precursors of carbon materials, which can be obtained by the carbonization of cellulose. Therefore, cellulose has a high potential as a precursor for the conductive matrix of polymer composites, including those of PMMA. There have been many studies on PMMA composites since PMMA not only has excellent physical and chemical properties, such as transparency and amenability, but also a manipulation of its porosity is possible.36 A transparent PMMA/cellulose compo-

Figure 1. Schematic illustration of fabrication route of PMMA/BC composite monolith and its conversion to functional activated carbon.

This composite monolith with a unique 3D−3D interpenetrating network (IPN) morphology was prepared by an environmentally friendly and time-saving facile approach based on TIPS. The as-obtained PMMA/BC composite showed not only very low density but also great improvement in thermal stability and mechanical property. Next, the composite monolith including the cellulose nanofibers matrix was subjected to carbonization, leading to a novel carbon material consisting of the BC-originated layer structure with a large specific area and mesopores formed during PMMA pyrolysis. BC can act as the precursor for the conducting material and function as matrix, and PMMA acted as the activator source providing CO2, H2O, CH4, and H2.48,49 The hierarchical structure as well as the layered morphology of the resulting carbon could improve the conductive speed of the electron, while the high surface area promoted the capacitance. In the present study, the fundamental electrochemical properties of the resulting carbon for applications in the electrode materials of ECs were examined. Importantly, the as-obtained activated carbon showed an excellent and compatible electrochemical performance, in terms of the energy density and capacitance retention. These target functions in the present proposal are important in energy fields, and these features are highly significant in the development of novel carbon electrodes. 9391


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remove the activator and other chemicals and washed repeatedly by water until the washed water displayed neutral pH. Subsequently, it was dried at 105 °C overnight and ground to powder. The resultant porous carbons were named as a-PMMA/BC-T, where T represents the activation temperature (T = 750, 850, or 900 °C). For reference, the samples, which were named c-PMMA, c-BC, and c-PMMA/BC, were prepared by only carbonization without KOH treatment at 900 °C. Due to the weak mechanical property of c-PMMA and c-BC, activated carbon based on PMMA and BC (a-PMMA, a-BC, respectively) could not be prepared. Electrode Preparation. The as-carbonized active material was spread on glass GCE. First, the GCE electrode was polished with αAl2O3 power (0.05 μm diameter), then ultrasonically rinsed with ethanol and deionized water and dried at ambient temperature. Finally, 2 mg of the carbon power was dispersed in a mixture of 1 mL of N,Ndimethylformamide (DMF) and 0.025 mL of Nafion (0.05%) to form the carbon dispersion. Twenty microliters of the suspension (2 mg mL−1 concentration) was dropped onto the GCE surface, and the electrode was subsequently dried by irradiation with an infrared lamp. This electrode is the work electrode.51 The loading mass of all the electrodes was maintained at the same level to minimize the test deviation. Electrochemical Measurements. All the electrochemical measurements were conducted using a CHI660E electrochemical workstation (CH instruments, USA) in a traditional three-electrode cell with 1 M H2SO4 aqueous electrolyte at room temperature. Saturated Ag/AgCl and platinum wire electrodes were used as reference and auxiliary electrodes, respectively. The cyclic voltammetry (CV) tests were performed in the potential range from 0 to 0.8 V at different scan rates. The electrochemical impedance spectroscopy (EIS) analysis was conducted in the 1 M H2SO4 aqueous electrolyte with the frequency range from 0.1 Hz to 1000 kHz. Galvanostatic charge−discharge (GCD) measurements were carried out at different current densities from 0.5 to 20 A g −1 with the same potential window of 0−0.8 V. The specific capacitance of the prepared electrode can be counted using the following formula:

EXPERIMENTAL SECTION

Materials. PMMA with a molecular weight of 9.5 × 10 was purchased from Aladdin Industrial Corporation (Shanghai, China). Ethanol and sodium hydroxide were obtained from Tianjin Zhiyuan Chemical Industries, Ltd. Potassium hydroxide and concentrated sulfuric acid were bought from TianJin Yongsheng Fine Chemical Co., Ltd. (Tianjin, China). Commercial BC with a cubic structure of 0.8 cm × 0.8 cm × 0.8 cm was supplied by Yeguo Foods Co., Ltd. (Japan). Rod glass carbon electrodes (GCEs) were bought from Chenhua Co., Ltd. (Shanghai, China). The reference electrodes and auxiliary electrodes were purchased from Tianjin Gaoss Union Technology Co., Ltd. (Tianjin, China) and Wuhan Gaoss Union Technology Co., Ltd. (Wuhan, China), respectively. All other reagents were of analytical grade and used as received without further purification. One molar sulfuric acid solution was prepared as the supporting electrolyte. Deionized (DI) water for solution preparation was supplied from a Millipore Autopure system (18.2 MΩ, Millipore, Ltd., USA). Characterization. The surface morphology and structure of the PMMA/BC monoliths were examined by a SEM instrument (SU3500 HITACHI, Japan) at an acceleration voltage of 15 kV. Transmission electron microscopy (TEM) analysis was performed by using a Tecnai G2 F20 S-TWIN instrument (USA) operated at 200 kV accelerating voltage. Nitrogen adsorption/desorption isotherms were obtained by a TR2 Star3020 surface area and pore size analyzer. Before the measurement, the samples were degassed for 8 h at 120 °C under vacuum. The specific surface area was obtained by the Brunauer− Emmett−Teller (BET) method, and the pore size distribution (PSD) was calculated on the basis of the Barrett−Joyner−Halenda (BJH) model. Thermogravimetric analysis (TGA) was conducted with a STA 449C thermogravimetric analyzer (NETZSCH, Germany) in the temperature range of 33 to 900 °C under a nitrogen atmosphere to evaluate the thermal stability of the samples. The X-ray diffraction patterns of the power sample were generated by D8 ADVAHCL (Bruker, Germany, equipped with Cu Kα radiation, λ = 1.5406 Å) from 5° to 80° at a scan rate of 0.03° min−1. The infrared spectra were recorded on a TENSOR27 (Bruker, Germany) FT-IR spectrometer from 4000 to 500 cm−1. Purification of Bacterial Cellulose. The commercial BC hydrogel was immersed in a large quantity of DI water with gentle stirring, and the water was changed every 3 h for 10 times. Then, the resulting BC hydrogel was treated with 0.5 M NaOH at 80 °C for 3 h. Finally, the hydrogel was washed with water, in a manner similar to the above-described process, until a neutral pH was reached. The washed BC was stored in water that includes 1% ethanol at 5 °C.50 Fabrication of PMMA/BC Composite Monoliths. A typical fabrication protocol for the composite monolith is as follows.40 The medium of the purified BC hydrogel was exchanged from water to an ethanol/water mixture (80/20 v/v). The hydrogel was immersed in a large amount of the ethanol/water mixture for 24 h, and this process was carried out at least three times. Then, a certain amount of BC was immersed in a mixed solvent of ethanol (8 mL) and water (2 mL) under gentle stirring. PMMA power (0.8 g) was dissolved in the solution at 60 °C with gentle stirring for 24 h. After the removal of the BC gel from the PMMA solution, the gel was cooled to ambient temperature, at which point phase separation of the PMMA solution in the gel took place. After 12 h, the medium of the obtained gel was exchanged from the mixed ethanol/water solvent to water by immersion in an excess of water with gentle shaking at least three times, and the resulting composite monolith was dried under vacuum at room temperature. Fabrication of Activated Carbon Material. The as-obtained PMMA/BC composite monolith was immersed in 10 M KOH solution for 24 h. The weight ratio of the monolith and KOH was 0.5:1.0. The water in the monolith was dehydrated at 105 °C for 10 h. The sample was transferred into a tubular furnace and heated from 50 °C to a predicted temperature with a heating rate of 10 °C min−1 in an N2 atmosphere (flow rate of 20 mL min−1). Finally, the sample was kept for 2 h, cooled to room temperature, and maintained there for 80 min. It was boiled with a large amount of dilute HCl for 5 min to 4

Cs = (I × Δt)/(m × ΔV)

(1)

where Cs is specific capacitance, I is current discharge density (A), and m is the mass of the electroactive materials in the electrodes (g). ΔV is the potential window (V) after insulation resistance (IR) drop, and Δt is discharge time (s). The energy density (W h kg −1) and the average power density (W −1 kg ) were calculated by the following equation:

Ed = CsΔU2/(2 × 3.6)

(2)

P = E × 3600/Δt

(3)

where Cs (F g−1) is specific capacitance, E is energy density (W kg −1), and Δt is discharge time (s).52

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RESULTS AND DISCUSSION PMMA/BC Composite Monolith. Solvent selection is a crucial factor in the fabrication of a PMMA monolith by TIPS. It is well known that both water and ethanol are not solvents for PMMA. Interestingly, PMMA can be dissolved in a mixture of ethanol and water (80/20 v/v) by heating.40 The preparation of the PMMA/BC composite monolith was carried out under conditions similar to those used to fabricate the PMMA monolith. The medium of the BC gel can be easily changed from water to various organic solvents. First, BC was immersed in a mixed solvent of ethanol and water (80:20 v/v) for 24 h to exchange the medium water of BC gel; then, PMMA was added into the mixture and dissolved at 60 °C to afford the PMMA solution in the presence of the BC gel (Figure S2). PMMA could permeate well through the inner network of the BC fibers because the solvent composition inside and outside the BC gel 9392


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Figure 2. SEM images with different magnification: PMMA monolith (a) and PMMA/BC composite monolith (b).

after the phase separation, the weight ratio of PMMA and BC in the resulting composite was estimated to be 4:1. The morphology of the as-obtained composite monolith was observed by SEM (Figure 2). In the resulting composite, the layered structure of BC (Figure S1) remained at the macroscopic scale. Interestingly, the magnified SEM image showed that the bead-connected monolithic PMMA skeleton and BC fibers were entangled to form the IPN structure, as illustrated in Figure 1. This morphology with a 2D layered structure of BC and 3D−3D IPN structure of the PMMA skeleton and BC fiber was highly unique, which is significant for development of functional porous materials with a precisely controlled structure. Owing to this characteristic morphology, the resulting composite monolith showed good mechanical property (Figure S4). The PMMA monolith was easily broken by compression, whereas a similar compression on the PMMA/ BC composite monolith resulted in only a vertical deformation without the monolith breaking. This difference strongly suggests the good reinforcement effect by the BC nanofibers with the layered structure, supporting the uniform IPN structure of the PMMA skeleton and BC fibers. An IPN is one of the most typical polymer structures, in which one cross-linked polymer interpenetrated another polymer network at the molecular level.53 IPN is an important design for high-performance polymeric materials. The composites study in this work demonstrate a unique IPN

was the same. During the cooling step, the phase separation took place to form the PMMA monolith in the BC gel. The physical and chemical interaction between PMMA and BC can ensure the uniform distribution of PMMA on the surface of the BC nanofibers. The diffusion of the PMMA solution in the BC gel was further confirmed by FT-IR (Figure S3). In Figure S3(a), the lines labeled (1), (2), (3), and (4) represent cross sections of the right side, central side, orthogonal central side, and vertical central layer of the composite monolith, respectively. The FT-IR spectra of the four points were almost the same, strongly supporting that PMMA is homogeneously inserted in the BC gel. The fabrication of the composite monolith with different amounts of BC and varying concentrations of PMMA was examined. When a large amount of BC was used, most of the solvent was adsorbed by the BC gel, which might not be beneficial to the subsequent phase separation process. On the other hand, a high concentration of PMMA resulted in the shape deformation of BC. Thus, controlling the amount of BC and the concentration of PMMA were crucial to fabricate composite materials with a well-defined structure. A combination of eight or nine BC cubes (cube size: ca. 8 mm × 8 mm × 8 mm) and 100 mL of the PMMA solution (80 mg mL−1 concentration) was a typical condition for the fabrication of the porous composite. From the change in the weight before and 9393


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was prepared under similar conditions. The layered and wrinkled structure, which may be derived from BC, was observed in the SEM image of c-BC and c-PMMA/BC at a low magnification (Figure 4). The PMMA monolithic skeleton was not observed in c-PMMA/BC. The SEM image at high magnification showed different morphology for c-BC and cPMMA/BC; c-PMMA/BC had a honeycomb-like structure with a highly porous morphology. These data strongly suggest that the PMMA monolithic skeleton affected the morphology of the carbon material. The SEM image of the activated carbon from a-PMMA/BC900 with a low magnification shows that a 3D honeycomb structure was formed and the layered structure derived from BC did not remain, suggesting that the morphological change took place during the activation process, upon using KOH. The rough surface with the hierarchical porous structure, which may lead to an effective improvement of conductive path in aPMMA/BC-900, was observed in the SEM image at a high magnification. The microstructure observation of c-PMMA, c-BC, cPMMA/BC, and a-PMMA/BC-T (T = 750, 850, and 900 °C) was carried out by TEM (Figure 5). TEM images of c-BC, c-PMMA, c-PMMA/BC, and a-PMMA/BC-750 showed that the as-prepared porous carbons possessed similar amorphous structure and irregular morphology with macro/mesopore overlapping distribution, and especially, these characteristics were found in a-PMMA/BC-750, implying that the porous carbon materials with micropores, mesopores, and macropores were successfully prepared. In addition, porous aggregated carbon particles were observed in Figure 5(c) for c-PMMA/ BC. The aggregation of carbon particles is favorable for electrolyte penetration through the electrode materials. On the other hand, Figure 5(e) and (f) showed that a-PMMA/BC-850 and a-PMMA/BC-900, respectively, had wrinkles and folds of edges structure, which suggests that the mesopores with high graphitic degree are distributed uniformly in the carbon surface. This structure can shorten the diffusion distance of electrolyte ions. The plicated lamella of a-PMMA/BC-850 was thinner than that of a-PMMA/BC-900, which is beneficial for the

structure consisting of the monolithic skeleton and nanofiber at the macroscopic scale by utilizing phase separation of the polymer solution in the presence of BC. The thermal properties of the composite monolith were evaluated by thermogravimetric analysis in a nitrogen atmosphere. Figure 3 shows TGA curves of the PMMA

Figure 3. TG curves of PMMA monolith and PMMA/BC composite monolith under nitrogen atmosphere.

monolith and PMMA/BC composite monolith. The thermal decomposition behaviors of both monoliths was similar. In the initial stage, a slight weight loss was found, which may be due to the evaporation of the water contained in the sample. A large weight loss was observed from the 320−410 °C. In the case of the PMMA monolith, few traces of the sample remained beyond 500 °C, whereas the residual ratio was about 35% for the PMMA/BC composite monolith. Considering the weight ratio of the composite monolith, cellulose as well as PMMA were partially converted to carbon by pyrolysis under the present conditions, whereas all the PMMA chains in the PMMA monolith were decomposed. This interesting result might be because the PMMA skeleton was created on the surface of BC nanofibers, which might facilitate physical interaction between the BC nanofibers and PMMA backbone. Conversion of Composite Monolith to Activated Carbon. The carbon material from the present composite monolith (c-PMMA/BC) was prepared by pyrolysis. For reference, the carbon from BC and PMMA (c-BC, c-PMMA)

Figure 4. SEM images with different magnification: c-BC (a), c-PMMA/BC (b), and a-PMMA/BC-900 (c). 9394


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Figure 5. TEM images of c-PMMA (a), c-BC (b), c-PMMA/BC (c), a-PMMA/BC-750 (d), a-PMMA/BC-850 (e), and a-PMMA/BC-900 (f).

composite material was significantly larger than that of cPMMA and c-BC obtained at the same sintering temperature, which may be owing to the formation of gases such as CO2, H2O, CH4, and H2, and other activators during the pyrolysis of PMMA. The small molecule gases from the PMMA pyrolysis can effectively increase the specific surface area of the product. Pore size distribution was calculated by using the nonlocal density functional theory (NLDFT) method (Figure S5(b)). Obviously, c-PMMA/BC displayed a highly porous structure with a broad PSD. The pore size of c-PMMA/BC was in the range from 2 to 100 nm, suggesting the existence of mesopores and macropores in c-PMMA/BC. This specific structure may ascribe to the unique structure of the precursor composite monolith, and it is suitable for high-performance EC applications. The specific surface area and pore structure were significantly influenced by the activation process with KOH and activation temperature. Figure 6 shows N2 adsorption−desorption isotherms of a-PMMA/BC-750, a-PMMA/BC-850, and aPMMA/BC-900. A-PMMA/BC-T displayed an I/IV-type isothermal. A hysteresis loop existing in all three samples is related to mesopores, implying a capillary condensation process in the pore structure. At low relative pressures (P/P0 < 0.2), the adsorption capacity sharply rose with the increase in relative pressure, indicative of micropores existence. Moreover, the formation of macropores were found in the relative pressure range from 0.9 to 1.0 due to a warped-tail phenomenon (Figure 6(a)). Hierarchical porous carbons with well-developed porosity might be generated by the self-activation process to offer convenient channels for KOH entrance, making activation agents richly in contact with carbon materials, which becomes a more effective activation step. The pore size distributions of samples were examined by BJH method (Figure 6(b)). Mesopores and macropores with broad pore size distribution ranging from 2 to 100 nm were found. These mesopores and macropores could serve as diffusion pathways for electrolyte ions. The specific surface area, total pore volume, and average pore diameter are summarized in Table 1. Both the BET surface area and the pore volume increased with an increase in activation temperatures due to sufficient activation degree, from

diffusion of electrolyte ions into mesopores, which corresponds well with the excellent electrochemistry preformed with aPMMA/BC-850. It is well known that an appropriate pore size distribution and large surface area of carbon materials are highly pertinent for excellent capacitive performance. Nitrogen adsorption− desorption measurements were carried out on the prepared samples to evaluate their pore properties. Figure S5(a) shows the nitrogen adsorption and desorption isotherms of the PMMA/BC composite monolith as well as three pyrolytic samples (c-PMMA, c-BC, and c-PMMA/BC), which were obtained by thermal treatment at 900 °C. The isotherms of all three carbon samples showed a typical type IV behavior (BDDT classification) with a H2 hysteresis loop at a high relative pressure (P/P0 of 0.45−0.95), indicating a clear capillary condensation derived from the mesopores (2−50 nm) inside the material. The specific surface area, pore volume, and mean pore diameter are summarized in Table 1. The pore volume and specific surface area of the PMMA/BC composite monolith were much smaller than those of the carbons prepared in this study, suggesting that the carbonization process provided the product with larger pore volume and surface area. The pore volume of the carbon from the Table 1. Specific Surface Areas, Pore Volumes, and Mean Pore Diameters of PMMA/BC Monolith, c-BC, c-PMMA/ BC Composite, and a-PMMA/BC-750, a-PMMA/BC-850, and a-PMMA/BC-900

Sample

Special surface area (m2 g−1)

Smicro (m2 g−1)

Pore volume (cm3 g−1)

Mean pore diameter (nm)

PMMA/BC monolith c-BC c-PMMA c-PMMA/BC a-PMMA/BC-750 a-PMMA/BC-850 a-PMMA/BC-900

5.62 241 253 264 1339 2076 1652

2.82 89 105 110 606 792 721

4.80 × 10−3 0.35 0.27 0.54 0.76 1.20 0.92

1.93 3.13 4.04, 7.71 4.88, 12.6 2.27, 3.55 2.22, 3.11 2.23, 3.21 9395


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Figure 6. N2 adsorption/desorption isotherms (a) and pore size distributions curve (b) of a-PMMA/BC-750, a-PMMA/BC-850, and a-PMMA/BC900.

Figure 7. CV curves of c-PMMA, c-BC, and c-PMMA/BC at a scan rate of 50 mV s−1(a). GCD curves of c-BC and c-PMMA/BC (b).

Figure 8. CV curves of a-PMMA/BC-750, a-PMMA/BC-850, and a-PMMA/BC-900 at a scan rate of 50 mV s−1(a). GCD curves of a-PMMA/BC750, a-PMMA/BC-850, and a-PMMA/BC-900 (b).

1339 m2 g−1 and 0.76 cm3 g−1 for a-PMMA/BC-750 to 2076 m2 g−1 and 1.20 cm3 g−1 for a-PMMA/BC-850, and the average pore diameter decreased from 2.27 to 2.22 nm, indicating the formation of the smaller pores produced by KOH etching. However, with a further increase in activation temperature, the BET surface area and pore volume dramatically decreased for aPMMA/BC, suggesting hole collapse at a high activation temperature due to abundant active sites, resulting in the formation of larger pores and lower specific surface area (1652 m2 g−1 for a-PMMA/BC-900). Therefore, a-PMMA/BC-850 should have the highest specific capacitance and rate capability of the electrode materials used in this study. It was reported that the polymerization of aniline26 and pyrrole54 in the presence of BC resulted in a similar porous composited structure. However, the morphology of the c-BC and c-PMMA/BC electrodes were different from the abovementioned results. In general, the surface morphology of carbon materials is relatively rough and porous, which is the intrinsic nature of activated carbons. All of the carbon materials prepared in this study exhibited a well-developed porous structure with different pore sizes.

The carbon structure was analyzed by X-ray diffraction (XRD) (Figure S6). A broad peak centered at 2θ = 24.6° was found for all the carbon materials, showing the typical amorphous structure of porous carbon. The addition of BC and the activation process affected the XRD results. A broad diffraction peak was seen at 2θ = 44.8° for a-PMMA/BC-750, aPMMA/BC-850, and a-PMMA/BC-900, supporting the amorphous structure of the present porous carbon.55 Such an amorphous porous structure can improve the transmission rate of electrons in the electrolyte solution (1 M H2SO4). A good contact between the active materials and the electrolyte is beneficial for improving the capacitance performance.56 Electrochemical Performance of the As-Obtained Carbon. Figure 7 shows CV and GCD curves of the c-BC, c-PMMA, and c-PMMA/BC electrodes. The CV curves were quasi-rectangular in shape along the current−potential, indicating that all three electrodes were stable in the 1 M H2SO4 solution and that they showed ideal capacitive behavior and excellent reversibility. The CV date also implies that redox reactions rarely occur on the electrode.57 Generally, it is predicted that the capacitance is proportional to the effective surface area. From the results of the N2 adsorption/desorption 9396


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Figure 9. CV curves of a-PMMA/BC-850 at different scan rates (a). GCD curves at a-PMMA/BC-850 at different current densities (b). Ragone plots of a-PMMA/BC-850 electrode (c). Cycle performance of a-PMMA/BC-850 asymmetric supercapacitor at current density of 5 A g−1 with GCD curves after 10,000 cycles (inset) (d).

from 750 to 850 °C, the corresponding capacitance increased from 186 to 266 F g−1. The results can be attributed to the contractible pore diameter and enlarged micropore area in aPMMA/BC-850 (Table 1). In general, high micropore areas and small pores are key factors for improvement of capacitance because they could offer abundant electrochemical active sites. CV curves of a-PMMA/BC-900 were much narrower, and discharge time was much shorter than that of a-PMMA/BC850. This may be due to collapse of the porous structure during the activation process at a higher temperature, resulting in a decrease in specific surface area and capacitance. As shown in Figure 8, the electrochemical performance of aPMMA/BC-850 was superior to that of other carbons prepared in this study. Thus, the electrochemical properties of a-PMMA/ BC-850 were examined in detail (Figure 9). The electrochemical performance of the electrodes and their capacitive behaviors are assessed using potential sweep CV and GDC tests in three-electrode systems.58,59 The CV curves of the aPMMA/BC-850 electrode (Figure 9(a)) indicate that the present device can be reversibly charged and discharged, and it displayed a stable electrochemical performance in the potential range of 0−0.8 V. In the scanning rate range from 10 to 100 mV s −1, the rectangular shapes of the CV curves were still retained, which demonstrates a good specific capacitive performance, equivalent series resistance (ESR), and fast diffusion of the electrolyte ions within the electrode at low scan rates.60 However, the rectangular CV became slightly distorted and deformed when the scanning rate reached 300 mV s−1. This may be because the transfer of electrons is far too quick at high scan rates.61 Figure 9(b) shows the typical charge−discharge curves of the different materials of the different current densities. The galvanostatic charge/discharge curves of a-PMMA/BC-850 were symmetrical isosceles triangles, and the voltage−time responses followed a linear relationship, depicting ideal electrochemical reversibility and superior charge−discharge properties, as well as a perfect double-layer feature of the capacitances, which is consistent with the CV results.62,63 This

characterization, it was found that c-PMMA/BC had the highest Barrett−Emmett−Teller (BET) surface area and c-BC had the lowest value (Table 1). The improvement in the specific capacitance of c-PMMA/BC is probably due to the presence of PMMA in the composite. A comparison of the CV curves of the c-PMMA, c-BC, and c-PMMA/BC electrodes shows that the c-PMMA/BC electrode not only displayed a higher background current at potential sweep but also showed a wider rectangle window. This clear difference may be attributed to the decomposition of PMMA during carbonization, producing small molecule gases such as CO2, H2O, H2, and CH4, as well as the native layered structure of BC, which could improve the specific surface area and pore structure of cPMMA/BC.48 Because of the higher current in the voltammograms of the c-PMMA/BC electrode when compared to c-BC and c-PMMA, an excellent capacitance for capacitors equipped with the c-PMMA/BC electrode could be obtained. As shown in the Figure 7(b), the discharging duration time of c-PMMA/ BC was obviously higher than those of c-BC and c-PMMA, and the specific capacitance of the c-PMMA/BC composite electrode was 140 F g−1 at a current density of 0.5 A g−1, while the specific capacitance of the bare c-BC electrode and cPMMA electrode were 50 and 67 F g−1, respectively. When the material was activated by KOH, the electrode material showed a higher current in the voltammograms and a more excellent capacitance. CV curves of a-PMMA/BC-T (T = 750, 850, and 900 °C) at a scan rate of 50 mV s−1 are shown in Figure 8(a). All the activated samples exhibited nearly rectangular shapes with a blurry broad peak at about 0.3 V, confirming the electrical double-layer capacitors (EDLC) performance with pseudocapacitance. Especially, for aPMMA/BC-850 electrode, the CV curves presented the largest potential window, suggesting the highest capacitance among the three samples. The charge−discharge curves of a-PMMA/ BC-T electrodes were symmetrically triangular and closely linear at a current density of 0.5 A g−1 (Figure 8(b)), demonstrating a good capacitive behavior. The discharge time gradually increased with an increase in activation temperature 9397


Research Article


Figure 10. Specific capacitance and capacitance retention of a-PMMA/BC-850 at different densities (a). Nyquist plots of EIS of c-PMMA, c-BC, cPMMA/BC, a-PMMA/BC-750, a-PMMA/BC-850, and a-PMMA/BC-900 (b). The electrical equivalent circuit model used for fitting impedance spectra (c).

and 95.1% of capacitance retention was observed even after 10,000 cycles, indicating excellent cycling stability and durability. This performance was closely related to the superior conductivity and high surface area of the carbon material. Importantly, the present activated carbon (AC) showed high energy density and capacitance retention in comparison with other reported carbons (Table S1). Figure 10(a) demonstrates the variation and retention of the specific capacitance versus current density. With an increase in the current density, the capacitance decreased. A retention of 80% was noticed when the current density ranged from 0.5 to 5 A g−1. The specific capacitances calculated by eq 1 were found to be 266, 240, 230, 212, 200, 195, and 191 F g−1 at current densities of 0.5, 1.0, 2.0, 3.0, 5.0, 10.0, and 20.0 A g−1, respectively. The excellent performances of a-PMMA/BC-850 are ascribed to the following reasons. First, the higher specific area is favorable for enhancing the charge storage density by accumulating more electrolyte ions. Second, the meso/ macropores in a-PMMA/BC-850 facilitates the quick pass of electrolyte ions in the pore channels at high current densities. This performance of a-PMMA/BC-850 suggests that the developed carbon material is excellent for application in supercapacitor. EIS was used to further study the fundamental behavior of the electrode materials and the structural characteristics of the supercapacitor cell during the charge/discharge process (Figure 10(b)).66 A single semicircle in the high-frequency region and a straight line in the low-frequency region were observed.67 The Nyquist plots were plotted using the equivalent circuit shown in Figure 10(c). The corresponding fitting results are presented in Table 2. At high frequencies, the intercept of the real part (Z′) represented solution resistance (Re).6 The six carbon electrodes had different Re values (Figure 10(b)). The smallest Re of a-PMMA/BC-850 was related to good pore permeability of the electrolyte ions and electrical conductivity. The semicircle diameter in the high-to-medium frequency is corresponding to charge-transfer resistance (Rct), which was caused by the Faradaic reactions and the double-layer capacitance (Cdl) (Figure 10(c)) on the electrode surface. The Rct of the c-PMMA/BC electrode was smaller than that of

carbon electrode exhibited a high specific capacitance of 266 F g−1 at a current density of 0.5 A g−1, which was much higher than that of other previously reported systems based on activated carbons.64 The increase in the capacitance may be partially explained by the porous structure, which allows electrons to move easily. The pyrolysis of PMMA led to a larger specific surface area at the electrode/electrolyte interface and enough pores to accommodate KOH, ensuring maximum contact between KOH and the sample, which makes the activation step more effective. In addition, the charge/discharge curve had no obvious insulation resistance (IR) drop at a higher current of 20 A g−1, indicating little overall resistance.62 The specific energy density of a-PMMA/BC-850 was calculated to be 23.6 W h kg−1 with a specific power density of 200 W kg−1 at a current of 0.50 A g−1, which was calculated from the GCD curves (Figure 9(c)). It is important to mention that this data was much higher than that of existing activated carbon-based supercapacitors.5 Furthermore, the energy density still remained at the ideal value of 17 W h kg −1 at a power density of 8.2 kW kg−1. Such high specific capacitance, specific energy, and power density could be facilitated by the structure of macro- and mesopores and high surface area of c-PMMA/ BC. Therefore, the ion transfer efficiency was improved, leading to the shortening of the diffusion distance.31 Cycling stability plays an important role in practical applications for supercapacitors. Some reports demonstrate that EDLCs consisting of porous carbons possessed excellent cycling stability due to the nonexistence of pseudocapacitance. However, their specific capacitance was relatively small. Although the specific capacitance of functionalized porous carbons might be as high as 150−300 F g−1, it often exhibited poor cycling stability65 due to the presence of oxygenated groups. Therefore, the development of new carbon materials with high specific capacitance and excellent cyclic stability is important. The cycling life of the a-PMMA/BC-850 electrode was analyzed by the GCD method at a strong current density of 5 A g−1 in the voltage range of 0−0.8 V for 10,000 cycles. The specific capacitance decreased by only 1.5% of its original capacitance after 5000 cycles as shown in the inset in Figure 9(d). The latter GCD curves were still symmetrical triangles, 9398



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Table 2. Calculated Values of Re, Rct, and Zw through CNLS Fitting of Experimental Impedance Spectra Based upon the Proposed Equivalent Circuit in Figure 10(c) Sample

Re (Ω)

Rct (Ω)

c-BC c-PMMA/BC c-PMMA a-PMMA/BC-750 a-PMMA/BC-850 a-PMMA/BC-900

72.0 42.0 45.1 33.03 27.6 28.9

35.0 30.0 33.1 29.0 26.1 28.0

× × × × × ×

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acssuschemeng.7b02488. Appearance image, SEM image of side view, and SEM image of top view of BC (Figure S1). Fabrication process of PMMA/BC composite monolith (Figure S2). Digital photograph of PMMA/BC composite monolith and FTIR spectrum at different views of the composite (Figure S3). Compression test of PMMA monolith and PMMA/ BC composite monolith (Figure S4). N2 adsorption/ desorption isotherms and pore size distributions curve of c-BC, c-PMMA, and c-PMMA/BC (Figure S5). XRD patterns of PMMA/BC composite monolith, c-BC, cPMMA, c-PMMA/BC and a-PMMA/BC-750, a-PMMA/ BC-850, and a-PMMA/BC-900 (Figure S6). Comparison of energy density, power density, capacitance retention of various carbon materials with a-PMMA/BC-850 (Table S1). (PDF)

Zw (Ω) 1.62 3.06 1.62 1.23 4.70 1.15

Research Article

10−4 10−3 10−3 10−3 10−4 10−3

c-BC and c-PMMA, which means that the contact area between electrode interface and electrolyte increased due to the developed macro/mesoporous structure of c-PMMA/BC. The electrode of a-PMMA/BC-850 exhibited the smallest semicircular diameter, and the Rct of a-PMMA/BC-850 was 26 Ω, which is probably owing to the better electrical conductivity and generation of meso/macropores by the self-activation process. This result is in good agreement with the CV and GCD results. The 45° slope of the curves at low frequencies is associated with the semi-infinite Warburg impedance (Zw).6 Meanwhile, a more vertical line at the low frequencies for aPMMA/BC-850 demonstrated the pure capacitive behavior of the porous materials and the charge-transfer behavior of the carbon materials. According to equivalent circuit results, Zw reached the lowest value for a-PMMA/BC-850. These EIS results show strong relationships with the fitting data of an equivalent circuit.

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AUTHOR INFORMATION

Corresponding Authors

*Tel.: 029-88302635. Fax: 029-88303527. E-mail: yhshen@ nwu.edu.cn (Y. Shen). *Tel: +81-6-6879-7364. Fax: +81-6-6879-7367. E-mail: [email protected] (H. Uyama). ORCID

Hiroshi Uyama: 0000-0002-8587-2507

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Author Contributions

The paper was written through contributions of all authors. All authors have given approval to the final version of the paper.

CONCLUSION In this work, hierarchical porous carbons with excellent performance as supercapacitor electrodes were developed. A novel PMMA-BC composite monolith with a characteristic IPN structure, consisting of a PMMA monolith skeleton and BC nanofibers, was prepared by TIPS. This composite with its unique morphology produced a carbon material with a large specific surface area and pore volume. By utilizing PMMA as the activator source, a porous structure was obtained, which enhanced the electrode performance in EDLC applications. Electrochemical analyses of the activated carbon from the composite monolith revealed that the materials have a high specific discharge capacitance of 266 F g−1 at a current density of 0.5 A g−1 in the potential of 0 V−0.8 V in a traditional threeelectrode system. The energy density can reach up to 23.6 W h kg−1 at a power density of 200 W kg−1; 95.1% of the initial capacitance was retained even after 10,000 cycles. These excellent electrochemical performances, especially the compatible high energy density and capacitance retention, have not been achieved for the activated carbons reported thus far. This might be because the IPN structure of the starting composite provides the hierarchical porous carbon with a specific mesopore/macropore structure. These distinguishing features of the carbon materials reported in this work make them promising as electrodes in supercapacitors and energy storage devices. However, additional improvement in the capacitance and flexibility is required for future applications. Further studies on BC-based composite materials to address these issues are in progress in our laboratory.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This study is financially supported by the Natural Science Foundation of China (No. 21675125), Amygdalus pedunculata Engineering Technology Research Center of State Forestry Administration, Key laboratory of Yulin Desert Plants Resources, a Grant-in-Aid for Scientific Research from the Japan Society for the Promotion of Science (No. 16K14081).

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REFERENCES

(1) Zhong, C.; Deng, Y.; Hu, W.; Qiao, J.; Zhang, L.; Zhang, J. A Review of Electrolyte Materials and Compositions for Electrochemical Supercapacitors. Chem. Soc. Rev. 2015, 44, 7484−7539. (2) Yan, K.; Kong, L.-B.; Shen, K.-W.; Dai, Y.-H.; Shi, M.; Hu, B.; Luo, Y.-C.; Kang, L. Facile Preparation of Nitrogen-Doped Hierarchical Porous Carbon With High Performance in Supercapacitors. Appl. Surf. Sci. 2016, 364, 850−861. (3) Yan, J.; Wang, Q.; Wei, T.; Fan, Z. Recent Advances in Design and Fabrication of Electrochemical Supercapacitors with High Energy Densities. Adv. Energy Mater. 2014, 4, 1300816−1300859. (4) Burke, A. R&D Considerations for the Performance and Application of Electrochemical Capacitors. Electrochim. Acta 2007, 53, 1083−1091. (5) Zhu, Y.; Murali, S.; Stoller, M. D.; Ganesh, K. J.; Cai, W.; Ferreira, P. J.; Pirkle, A.; Wallace, R. M.; Cychosz, K. A.; Thommes, M.; et al. Carbon-Based Supercapacitors Produced by Activation of Graphene. Science 2011, 332, 1537−1541.

9399


Research Article

ACS Sustainable Chemistry & Engineering (6) Wang, G.; Zhang, L.; Zhang, J. A Review of Electrode Materials for Electrochemical Supercapacitors. Chem. Soc. Rev. 2012, 41, 797− 828. (7) Kandalkar, S. G.; Dhawale, D. S.; Kim, C. K.; Lokhande, C. D. Chemical Synthesis of Cobalt Oxide Thin Film Electrode for Supercapacitor Application. Synth. Met. 2010, 160, 1299−1302. (8) Miller, J. R.; Simon, P. Materials science. Electrochemical Capacitors for Energy Management. Science 2008, 321, 651−652. (9) Pech, D.; Brunet, M.; Durou, H.; Huang, P.; Mochalin, V.; Gogotsi, Y.; Taberna, P. L.; Simon, P. Ultrahigh-Power Micrometresized Supercapacitors Based on Nnion-like Carbon. Nat. Nanotechnol. 2010, 5, 651−654. (10) Simon, P.; Gogotsi, Y. Materials for Electrochemical Capacitors. Nat. Mater. 2008, 7, 845−854. (11) Brousse, T.; Belanger, D.; Long, J. W. To Be or Not To Be Pseudocapacitive? J. Electrochem. Soc. 2015, 162, A5185−A5189. (12) Gogotsi, Y.; Simon, P. True Performance Metrics in Electrochemical Energy Storage. Science 2011, 334, 917−918. (13) Chen, H.; Wang, F.; Tong, S.; Guo, S.; Pan, X. Porous Carbon with Tailored Pore Size for Electric Double Layer Capacitors Application. Appl. Surf. Sci. 2012, 258, 6097−6102. (14) Sun, F.; Gao, J.; Liu, X.; Pi, X.; Yang, Y.; Wu, S. Porous Carbon with a Large Surface Area and an Ultrahigh Carbon Purity Via Templating Carbonization Coupling With KOH Activation as Excellent Supercapacitor Electrode Materials. Appl. Surf. Sci. 2016, 387, 857−863. (15) Lota, G.; Centeno, T. A.; Frackowiak, E.; Stoeckli, F. Improvement of the Structural and Chemical Properties of a Commercial Activated Carbon for Its Application in Electrochemical Capacitors. Electrochim. Acta 2008, 53, 2210−2216. (16) Zhang, J.; Yu, Y.; Liu, L.; Wu, Y. Graphene-hollow PPy Sphere 3D-Nanoarchitecture with Enhanced Electrochemical Performance. Nanoscale 2013, 5, 3052−3057. (17) Liu, J.; An, J.; Ma, Y.; Li, M.; Ma, R. Synthesis of a GraphenePolypyrrole Nanotube Composite and Its Application in Supercapacitor Electrode. J. Electrochem. Soc. 2012, 159, A828−A833. (18) Mao, L.; Chan, H. S. O.; Wu, J. Cetyltrimethylammonium Bromide Intercalated Graphene/Polypyrrole Nanowire Composites for High Performance Supercapacitor Electrode. RSC Adv. 2012, 2, 10610−10617. (19) Xu, C.; Sun, J.; Gao, L. Synthesis of Novel Hierarchical Graphene/Polypyrrole Nanosheet Composites and Their Superior Electrochemical Performance. J. Mater. Chem. 2011, 21, 11253−11258. (20) Li, J.; Xie, H. Synthesis of Graphene Oxide/Polypyrrole Nanowire Composites for Supercapacitors. Mater. Lett. 2012, 78, 106− 109. (21) Biswas, S.; Drzal, L. T. Multilayered Nanoarchitecture of Graphene Nanosheets and Polypyrrole Nanowires for High Performance Supercapacitor Electrodes. Chem. Mater. 2010, 22, 5667−5671. (22) Ha, J. H.; Shah, N.; Ul-Islam, M.; Khan, T.; Park, J. K. Bacterial Cellulose Production from a Single Sugar α-linked Glucuronic Acidbased Oligosaccharide. Process Biochem. 2011, 46, 1717−1723. (23) Shah, N.; Ulislam, M.; Khattak, W. A.; Park, J. K. Overview of Bacterial Cellulose Composites: A Multipurpose Advanced Material. Carbohydr. Polym. 2013, 98, 1585−1598. (24) Phisalaphong, M.; Suwanmajo, T.; Tammarate, P. Synthesis and Characterization of Bacterial Cellulose/Alginate Blend Membranes. J. Appl. Polym. Sci. 2008, 107, 3419−3424. (25) Hu, W.; Chen, S.; Xu, Q.; Wang, H. Solvent-free Acetylation of Bacterial Cellulose under Moderate Conditions. Carbohydr. Polym. 2011, 83, 1575−1581. (26) Hu, W.; Chen, S.; Yang, Z.; Liu, L.; Wang, H. Flexible Electrically Conductive Nanocomposite Membrane Based on Bacterial Cellulose and Polyaniline. J. Phys. Chem. B 2011, 115, 8453−8457. (27) Fu, L.; Zhang, Y.; Li, C.; Wu, Z.; Zhuo, Q.; Huang, X.; Qiu, G.; Zhou, P.; Yang, G. Skin Tissue Repair Materials from Bacterial Cellulose by a Multilayer Fermentation Method. J. Mater. Chem. 2012, 22, 12349−12357.

(28) Liu, Y.; Lu, T.; Sun, Z.; Chua, D. H. C.; Pan, L. Ultra-thin Carbon Nanofiber Networks Derived from Bacterial Cellulose for Capacitive Deionization. J. Mater. Chem. A 2015, 3, 8693−8700. (29) Wu, Z. Y.; Li, C.; Liang, H. W.; Chen, J. F.; Yu, S. H. Ultralight, Flexible, and Fire-resistant Carbon Nanofiber Aerogels from Bacterial Cellulose. Angew. Chem., Int. Ed. 2013, 52, 2925−2929. (30) Lee, K.-Y.; Qian, H.; Tay, F. H.; Blaker, J. J.; Kazarian, S. G.; Bismarck, A. Bismarck, A. Bacterial Cellulose as Source for Activated Nanosized Carbon for Electric Double Layer Capacitors. J. Mater. Sci. 2013, 48, 367−376. (31) Chen, L.-F.; Huang, Z.-H.; Liang, H.-W.; Yao, W.-T.; Yu, Z.-Y.; Yu, S.-H. Flexible all-Solid-State High-power Supercapacitor Fabricated with Nitrogen-doped Carbon Nanofiber Electrode Material Derived from Bacterial Cellulose. Energy Environ. Sci. 2013, 6, 3331− 3338. (32) Delbert, E. D.; James, O. S.; John, M. B. Optically Transparent Composite Material and Process for Preparing Same. U.S. Patent US5665450 A, 1997. (33) Ul-Islam, M.; Khan, T.; Park, J. K. Nanoreinforced Bacterial Cellulose-Montmorillonite Composites for Biomedical Applications. Carbohydr. Polym. 2012, 89, 1189−1197. (34) Ul-Islam, M.; Khan, T.; Park, J. K. Water Holding and Release Properties of Bacterial Cellulose Obtained by in Situ and Ex Situ Modification. Carbohydr. Polym. 2012, 88, 596−603. (35) Ifuku, S.; Nogi, M.; Abe, K.; Handa, K.; Nakatsubo, F.; Yano, H. Surface Modification of Bacterial Cellulose Nanofibers for Property Enhancement of Optically Transparent Composites: Dependence on Acetyl-Group DS. Biomacromolecules 2007, 8, 1973−1978. (36) Sathya, S.; Murthy, P. S.; Das, A.; Gomathi Sankar, G.; Venkatnarayanan, S.; Pandian, R.; Sathyaseelan, V. S.; Pandiyan, V.; Doble, M.; Venugopalan, V. P. Marine Antifouling Property of PMMA Nanocomposite Films: Results of Laboratory and Field Assessment. Int. Biodeterior. Biodegrad. 2016, 114, 57−66. (37) Liu, H.; Liu, D.; Yao, F.; Wu, Q. Fabrication and Properties of Transparent Polymethyl methacrylate/Cellulose Nanocrystals Composites. Bioresour. Technol. 2010, 101, 5685−5692. (38) Pham, V. H.; Dang, T. T.; Hur, S. H.; Kim, E. J.; Chung, J. S. Highly Conductive Poly (methyl methacrylate) (PMMA)-Reduced Graphene Oxide Composite Prepared by Self-Assembly of PMMA Latex and Graphene Oxide through Electrostatic Interaction. ACS Appl. Mater. Interfaces 2012, 4, 2630−2636. (39) Okada, K.; Nandi, M.; Maruyama, J.; Oka, T.; Tsujimoto, T.; Kondoh, K.; Uyama, H. Fabrication of Mesoporous Polymer Monolith: a Template-free Approach. Chem. Commun. 2011, 47, 7422−7424. (40) Yoneda, S.; Han, W.; Hasegawa, U.; Uyama, H. Facile fabrication of Poly(methyl methacrylate) Monolith via Thermally Induced Phase Separation by Utilizing Unique Cosolvency. Polymer 2014, 55, 3212−3216. (41) Sun, X.; Fujimoto, T.; Uyama, H. Fabrication of a Poly(vinyl alcohol) Monolith via Thermally Impacted Non-Solvent-Induced Phase Separation. Polym. J. 2013, 45, 1101−1106. (42) Wang, G.; Uyama, H. Reactive Poly(ethylene-co-vinyl alcohol) Monoliths withTunable Pore Mmorphology for Enzyme Immobilization. Colloid Polym. Sci. 2015, 293, 2429−2435. (43) Park, S. B.; Fujimoto, T.; Mizohata, E.; Inoue, T.; Sung, M. H.; Uyama, H. Fabrication of Poly(γ-glutamic acid) Monolith by Thermally Induced Phase Separation and Its Application. J. Microbiol. Biotechnol. 2013, 23, 942−952. (44) Okada, K.; Maruyama, J.; Uyama, H. Fabrication and Electrochemical Capacitive Behaviors of a Carbon Nanotube-Coated Polymer Monolith. Electrochemistry 2013, 81, 789−791. (45) Wang, G.; Xin, Y.; Uyama, H. Facile Fabrication of Mesoporous Poly(ethylene-co-vinyl alcohol)/Chitosan Blend Monoliths. Carbohydr. Polym. 2015, 132, 345−350. (46) Wang, G.; Uyama, H. Facile Synthesis of Flexible Macroporous Polypropylene Sponges for Separation of Oil and Water. Sci. Rep. 2016, 6, 21265−21271. 9400


Research Article

ACS Sustainable Chemistry & Engineering (47) Wang, G.; Yoshikawa, H.; Tamiya, E.; Uyama, H. Mesoporous Poly(ethylene-co-vinyl alcohol) Monolith Captured with Silver Nanoparticles as a SERS Substrate: Facile Fabrication and Ultra-high Sensitivity. RSC Adv. 2015, 5, 25777−25780. (48) Lehmann, F. A.; Brauer, G. M. Analysis of Pyrolyzates of Polystyrene and Poly(methyl Methacrylate) by Gas Chromatography. Anal. Chem. 1961, 33, 673−676. (49) Chang, Z.; Laverne, J. A. The Gases Produced in Gamma and Heavy-ion Radiolysis of Poly(methyl Methacrylate). Radiat. Phys. Chem. 2001, 62, 19−24. (50) Si, H.; Luo, H.; Xiong, G.; Yang, Z.; Raman, S. R.; Guo, R.; Wan, Y. One-step in Situ Biosynthesis of Graphene Oxide-Bacterial Cellulose Nanocomposite Hydrogels. Macromol. Rapid Commun. 2014, 35, 1706−1711. (51) Li, W.; Ding, Y.; Zhang, W.; Shu, Y.; Zhang, L.; Yang, F.; Shen, Y. Lignocellulosic Biomass for Ethanol Production and Preparation ofActivated Carbon Applied for Supercapacitor. J. Taiwan Inst. Chem. Eng. 2016, 64, 166−172. (52) Yan, J.; Fan, Z.; Wei, T.; Qian, W.; Zhang, M.; Wei, F. Fast and Reversible Surface Redox Reaction of Graphene−MnO2 Composites as Supercapacitor Electrodes. Carbon 2010, 48, 3825−3833. (53) Sperling, L. H. Interpenetrating Polymer Networks: An Overview. Adv. Chem. Ser. 1994, 239, 3−38. (54) Wang, H.; Bian, L.; Zhou, P.; Tang, J.; Tang, W. Core−sheath Structured Bacterial Cellulose/Polypyrrole Nanocomposites with Excellent Conductivity as Supercapacitors. J. Mater. Chem. A 2013, 1, 578−584. (55) Zhao, H.; Han, G.; Chang, Y.; Li, M.; Li, Y. The Capacitive Properties of Amorphous Manganese Dioxide Electrodeposited on Different Thermally-Treated Carbon Papers. Electrochim. Acta 2013, 91, 50−57. (56) Zhang, Y.; Yao, Q.-q.; Gao, H.-l.; Zhang, L.-S.; Wang, L.-Z.; Zhang, A.-Q.; Song, Y.-H.; Wang, L.-X. Synthesis and Electrochemical Performance of MnO2/BC Composite as Active Materials for Supercapacitors. J. Anal. Appl. Pyrolysis 2015, 111, 233−237. (57) Tian, Y.; Liu, Z.; Xue, R.; Huang, L. An Efficient Supercapacitor of Three-Dimensional MnO2 Film Prepared by Chemical Bath M. J. Alloys Compd. 2016, 671, 312−317. (58) Nystrom, G.; Razaq, A.; Stromme, M.; Nyholm, L.; Mihranyan, A. Ultrafast all-Polymer Paper-Based Batteries. Nano Lett. 2009, 9, 3635−3639. (59) Wu, Q.; Xu, Y.; Yao, Z.; Liu, A.; Shi, G. Supercapacitors Based on Flexible Graphene/Polyaniline Nanofiber Composite Films. ACS Nano 2010, 4, 1963−1970. (60) Hughes, M.; Chen, G. Z.; Shaffer, M. S. P.; Fray, D. J.; Windle, A. H. Electrochemical Capacitance of a Nanoporous Composite of Carbon Nanotubes and Polypyrrole. Chem. Mater. 2002, 14, 1610− 1613. (61) Wang, J.; Xu, Y.; Wang, J.; Du, X.; Xiao, F.; Li, J. High Charge/ Discharge Rate Polypyrrole Films Prepared by Pulse Current Polymerization. Synth. Met. 2010, 160, 1826−1831. (62) Long, C.; Qi, D.; Wei, T.; Yan, J.; Jiang, L.; Fan, Z. NitrogenDoped Carbon Networks for High Energy Density Supercapacitors Derived from Polyaniline Coated Bacterial Cellulose. Adv. Funct. Mater. 2014, 24, 3953−3961. (63) Ko, J. M.; Kim, K. M. Electrochemical Properties of MnO2/ Activated Carbon Nanotube Composite as an Electrode Material for Supercapacitor. Mater. Chem. Phys. 2009, 114, 837−841. (64) Qu, Q. T.; Shi, Y.; Li, L. L.; Guo, W. L.; Wu, Y. P.; Zhang, H. P.; Guan, S. Y.; Holze, R. V2O5·0.6H2O Nanoribbons as Cathode Material for Asymmetric Supercapacitor in K2SO4 Solution. Electrochem. Commun. 2009, 11, 1325−1328. (65) Zhang, L. L.; Zhao, X. S. Carbon-based Materials as Supercapacitor Electrodes. Chem. Soc. Rev. 2009, 38, 2520−2531. (66) Sugimoto, W.; Iwata, H.; Yokoshima, K.; Murakami, Y.; Takasu, Y. Proton and Electron Conductivity in Hydrous Ruthenium Oxides Evaluated by Electrochemical Impedance Spectroscopy: The Origin of Large Capacitance. J. Phys. Chem. B 2005, 109, 7330−7338.

(67) Yang, W.; Gao, Z.; Wang, J.; Ma, J.; Zhang, M.; Liu, L. Solvothermal One-step Synthesis of Ni-Al Layered Double Hydroxide/Carbon Nanotube/Reduced Graphene Oxide Sheet Ternary Nanocomposite with Ultrahigh Capacitance for Supercapacitors. ACS Appl. Mater. Interfaces 2013, 5, 5443−5454.

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