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Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries
Accepted Manuscript Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries Qiaoqiao Wang, Xiaoshu Zhu, Yuhan Liu, Yuyan Fang, Xiaosi Zhou, Jianchun Bao PII:
S0008-6223(17)31168-5
DOI:
10.1016/j.carbon.2017.11.054
Reference:
CARBON 12584
To appear in:
Carbon
Received Date: 21 September 2017 Revised Date:
13 November 2017
Accepted Date: 19 November 2017
Please cite this article as: Q. Wang, X. Zhu, Y. Liu, Y. Fang, X. Zhou, J. Bao, Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries, Carbon (2017), doi: 10.1016/ j.carbon.2017.11.054. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphical abstract
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Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries
a
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Qiaoqiao Wang a, Xiaoshu Zhu b, Yuhan Liu a, Yuyan Fang a, Xiaosi Zhou a, *, Jianchun Bao a Jiangsu Key Laboratory of New Power Batteries, Jiangsu Collaborative Innovation Center of
Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal
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University, Nanjing 210023, P. R. China
Center for Analysis and Testing, Nanjing Normal University, Nanjing 210023, P. R. China
*
Corresponding author.
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b
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E-mail address: [email protected] (X. Zhou).
ACCEPTED MANUSCRIPT Abstract Sodium-ion batteries (SIBs) have drawn ever-increasing attention for scalable electrical energy storage owing to the inexhaustible sources and wide distribution of sodium. However, to develop
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feasible anode materials still remains a great challenge for the practical application of SIBs. Here, we report hard carbons derived from a plentiful and deserted biomass of rice husk through a facile acid treatment and subsequent pyrolysis. The investigation illustrates that the electrochemical
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properties of the rice husk-derived hard carbons (RHHCs) are significantly influenced by the
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pyrolysis temperature because of the discrepancy in their microstructure. The RHHC pyrolyzed at 1300 °C (RHHC-1300) shows the highest reversible capacity of 372 mAh g−1 and good cycling stability and rate performance due to its large interlayer distance and suitable oxygen content. Moreover, full sodium-ion batteries are assembled to examine the application prospect using
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Na3V2(PO4)2F3/C and RHHC-1300 as cathode and anode materials, respectively, delivering a high-energy density of 185 Wh kg−1 and stable cycling performance. This work could intensify the
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fundamental understanding of the sodium storage mechanism in biomass-derived hard carbons.
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Keywords: Rice husk; Hard carbon; Sodium-ion battery; Anode; Storage mechanism
ACCEPTED MANUSCRIPT 1. Introduction As the most widely used energy storage system, lithium-ion batteries (LIBs) have dominated the power source market for electric vehicles and portable electronic devices due to their high energy
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density and long cycle lifespan [1]. However, the scarcity and uneven geographical distribution of lithium resource seriously hamper the large-scale applications of LIBs to energy storage on the grid [2]. Benefiting from the copious reserve and lower price of sodium with comparable physical and
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chemical properties to lithium, sodium-ion batteries (SIBs) have been considered as one of the most
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prospective candidates for large-scale stationary energy storage [3, 4]. In comparison with lithium ions, the radius of sodium ions is much larger (1.02 versus 0.76 Å) so that one cannot randomly choose electrode materials for SIBs from those for LIBs [5]. Hence, it is crucial to search for qualified host materials to supply enough space for sodium ions storage and transportation [6].
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Recently, a large number of layered oxides and polyanion compounds have been verified as main cathode materials for SIBs [7-9]. In terms of anode materials, commercial graphite cannot afford enough interlayer space to accommodate larger sodium ions and thus present low sodium storage
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[12].
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capability [10, 11]; in this aspect, it is extremely important to find suitable anode materials for SIBs
While many kinds of anode materials, such as carbon materials [13-28], alloys [29-34], metal oxides and sulfides [35-37], and organic compounds [38], have been reported to possess sodium storage property, most of them show poor cycle life or low specific capacity. Specifically, the structural disintegration caused by huge volume variation results in the loss of electrical contact and capacity decline in the case of alloy-type materials [39, 40]. The poor initial Coulombic efficiency (CE) of metal oxides or sulfides and the low sodium storage performance of organic compounds
ACCEPTED MANUSCRIPT may inhibit their practical applications in SIBs [41, 42]. Among all the anode materials, hard carbon with randomly oriented turbostratic nanocrystallites has demonstrated the most promising sodium storage performance [43-49]. Moreover, hard carbons derived from various biomass precursors and
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their fundamental sodium storage mechanisms have been widely studied recently [50-57]. Nonetheless, there are still two significant obstacles before realizing industrial production of hard
high price and low carbonization yield of precursors.
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carbons as anode materials for SIBs: unsatisfactory sodium storage capacity and high cost owing to
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Rice husk is a kind of agricultural deserted byproduct and is copiously existed in the world [58-62]. The abundant content of cellulose and hemicellulose makes rice husk a proper carbon source to fabricate hard carbons. Herein, we report a simple yet efficient approach to synthesize hard carbons by utilizing abandoned rice husk. The electrochemical properties of the resultant hard
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carbons as anode materials for SIBs are found to be considerably affected by pyrolysis temperature. The rice husk-derived hard carbon carbonized at 1300 oC (RHHC-1300) shows promising sodium storage performance with a high reversible capacity of 372 mAh g−1, a medium initial CE of 66%,
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and good cycling stability and rate performance. The practical application prospect of RHHC-1300
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is further corroborated by using full sodium-ion batteries with Na3V2(PO4)2F3/C as cathode material, exhibiting a high energy density of 185 Wh kg−1 and excellent cycling performance and rate capability. This contribution provides a route for exploiting deserted agricultural biomass to prepare hard carbon anode materials for SIBs. 2. Experimental 2.1. Materials Synthesis The compressed rice husks were bought from Taobao (3 kg in pack). Prior to the synthesis
ACCEPTED MANUSCRIPT process, impurities in the rice husks were removed by washing in 6 M HCl for 6 h and 10% HF for 12 h, rinsing with water, and drying at 120 oC overnight in a vacuum oven. For the pyrolysis process, 2 g of purified sample was loaded in a tube furnace and carbonized for 2 h under an argon
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flow with a heating rate of 5 oC min−1. The pyrolysis temperatures were 1100, 1300, and 1500 oC, and the obtained rice husk-derived hard carbons were denoted as RHHC-1100, RHHC-1300, and RHHC-1500, respectively.
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2.2. Materials Characterization
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Scanning electron microscopy (SEM) images were taken by a JEOL JSM-7600F scanning electron microscope operated at 10 kV. X-ray diffraction (XRD) patterns were collected using a Rigaku SmartLab diffractometer with Cu Kα irradiation (λ = 1.5406 Å). Raman spectra were determined on a Labram HR800 with a laser wavelength of 514 nm. Nitrogen sorption
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measurements were carried out on an ASAP 2050 surface area-pore size analyzer. High-resolution transmission electron microscopy (HRTEM) images were obtained by a JEOL JEM-2100F transmission electron microscope operated at 200 kV. X-ray photoelectron spectroscopy (XPS) was
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acquired on an ESCALab250Xi electron spectrometer from VG Scientific using 300 W Al Kα
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radiation. Scanning transmission electron microscopy (STEM) as well as elemental mapping measurements were performed on the JEOL JEM-2100F transmission electron microscope equipped with a Thermo Fisher Scientific energy-dispersive X-ray spectrometer. 2.3. Electrochemical Measurements Electrochemical tests were conducted using CR2032 coin cells. The working electrodes were prepared by coating the mixed slurry of RHHC, carbon black (Super-P), and polyvinyl fluoride (PVDF) binder in N-methyl-2-pyrrolidone with a weight ratio of 8:1:1 onto copper foil, and then
ACCEPTED MANUSCRIPT dried at 80 °C in a vacuum oven for 12 h. It is noteworthy that the RHHC particles were ground for 30 minutes using a mortar and pestle prior to direct use as electrode materials. The loading mass of active material was controlled between 1.0 and 1.5 mg cm−2. Glass fiber (GF/D) from Whatman and
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sodium metal were used as separators and counter electrodes, respectively. The electrolyte was composed of 1 M NaClO4 in a mixture of ethylene carbonate and diethyl carbonate (1:1 in volume). All coin cells were assembled in an argon-filled glovebox (H2O, O2 < 0.1 ppm, MBraun, Germany).
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Galvanostatic discharge and charge tests were carried out on Land CT2001A multichannel battery
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testing systems in the fixed voltage range of 0.01−2 V vs Na/Na+ at various current densities. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were conducted on a PARSTAT 4000 electrochemical workstation. CV was recorded at a scan rate of 0.1 mV s−1, and EIS was determined by applying a sine wave with an amplitude of 10.0 mV over the
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frequency range from 100 kHz to 100 mHz. Sodium-ion full batteries were fabricated using RHHC-1300 as the anode material and Na3V2(PO4)2F3/C as the cathode material in CR2032 coin-type cells. The Na3V2(PO4)2F/C nanocomposite was synthesized through a previously reported
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sol−gel method[63]. The weight ratio of the two electrodes (negative/positive) was 1:4. The full
g−1.
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cells were charged and discharged in the voltage range of 1.6−4.3 V at a current density of 25 mA
2.4. Ex situ XRD, HRTEM, and AES Analysis For the electrodes employed for ex situ XRD, HRTEM, and AES measurements, carbon black was not used, while RHHC-1300 was increased to 90 wt %. The cycled cells were disassembled in the argon-filled glovebox, and the RHHC-1300 electrodes were collected and washed with dimethyl carbonate (DMC) several times. The naturally dried RHHC-1300 electrode was put onto a glass
ACCEPTED MANUSCRIPT slide and covered with a layer of Kapton film prior to being placed into the Rigaku SmartLab diffractometer. In case of ex situ AES, the naturally dried RHHC-1300 electrode was protected under argon atmosphere before being put into a ULVAC-PHI PHI-700 scanning Auger microscope
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operated at 5 kV. As for ex situ HRTEM measurement, the rinsed electrode material was dispersed in DMC under the protection of argon atmosphere by sonication. The HRTEM samples were prepared by pipetting a few microliters of the resulting dispersion onto carbon grids in the
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argon-filled glovebox. The naturally dried carbon grids were placed in sample holder for ex situ
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HRTEM experiment. 3. Results and discussion
Fig. 1a,b displays scanning electron microscopy (SEM) images of the inner and outer surfaces of the starting rice husk precursor. The SEM image shows that the inner surface is smooth and the
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outer surface is rugged. After acid treatment and subsequent pyrolysis at elevated temperature under argon atmosphere, the rice husks first become yellow, shrink in size, and then turn black, implying
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its successful removal of impurities and carbonization (Fig. S1)[64]. When rice husks are carbonized at T oC, the resulting rice husk-derived hard carbon is designated as RHHC-T (Table 1).
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The annealed rice husks are found to possess irregular granular shape with a size distribution range of 3−25 µm and a dense structure (Fig. 1c,d and Fig. S2). An obvious decrease in the size of sample is observed during carbonization. The carbon yields of rice husks at 1100 oC, 1300 oC, and 1500 oC are around 26%, 24%, and 23%, respectively. The dense structure of rice husk-derived hard carbon (RHHC) helps to form a stable SEI film and promotes the transportation of electrons, boosting the electrochemical property of RHHC. To analyze the microstructure of RHHC, X-ray diffraction (XRD) and Raman spectroscopy were
ACCEPTED MANUSCRIPT performed and the results are demonstrated in Fig. 2a,b. All XRD patterns show broad peaks at 2θ around 22.5o and 43.6o, which correspond to the (002) and (101) planes of disordered carbon structure. There is no evident variation in the peak shape of RHHC with increasing the pyrolysis
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temperature, indicating the characteristic of hard carbon. The (002) peak position diverts to a higher angle with rising the carbonization temperature, suggesting the local structural evolution to a smaller d-spacing. As shown in Table 1, the interlayer spacing shrinks from 4.03 Å for RHHC-1100
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to 3.95 and 3.84 Å for RHHC-1300 and RHHC-1500, respectively. In Raman spectra (Fig. 2b), two
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detached characteristic peaks at ~1350 and ~1606 cm−1 can be assigned to D band and G band [47, 65], respectively, indicating the amorphous architecture of RHHCs. The full widths at half maximum of D and G bands in the Raman spectra diminish with increasing pyrolysis temperature, which further reveals the evolution of local long-range ordered structure, i.e., turbostratic
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nanocrystallite size. The turbostratic nanocrystallite size along the c-axis (Lc) of a rice husk-derived hard carbon can be computed by the Scherrer equation. On average, there are about 2.9, 3.1, and 3.4 layers of graphene sheets in the nanocrystallites of RHHC-1100, RHHC-1300, and RHHC-1500
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(Table S1), respectively. Detailed calculation can be found in Fig. S3. Furthermore, the
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nanocrystallite size along the ab-plane (La) was also calculated through Raman spectrum and the equation below [66]:
4 IG La = ( 2.4 × 10−10 ) λnm ID
where λnm is the laser wavelength, ID represents the integrated D band, and IG indicates the integrated G band (Fig. S4). The Raman results show that RHHC-1100 has the smallest turbostratic nanocrystallite size with La value of 11.0 nm, while the value rises to 12.0 nm for RHHC-1300, and 12.5 nm for RHHC-1500 (Table S2). Overall, the XRD and Raman analysis show that both
ACCEPTED MANUSCRIPT d-spacing and turbostratic nanocrystallite size of RHHCs are strongly affected by the carbonization temperature. Nitrogen adsorption−desorption isotherms in Fig. S5 displays the Brunauer−Emmett−Teller
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(BET) surface area and pore size distribution of RHHC. The BET surface area is closely associated with the pyrolysis temperature. The computed surface area diminishes with the increase of carbonization temperature, RHHC-1100 possesses the highest BET surface area of 2.68 m2 g−1
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while RHHC-1300 and RHHC-1500 encompass lower surface areas of 0.27 and 0.23 m2 g−1,
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respectively. These small surface areas of RHHC-1300 and RHHC-1500 could lead to finite SEI layers and hence enhance the initial CE [67], in accordance with the electrochemical results (Table 1). Besides, the pore size distribution in Fig. S5b also demonstrates different pore size relying on the annealing temperature. In terms of RHHC-1100, a few micropores and mesopores are existing
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with a total pore volume of 3.35 × 10−3 cm3 g−1, while these pore volumes remarkably decrease with increasing the heat treatment temperature. The pore size distribution of RHHC-1300 or RHHC-1500 is basically undetectable[68]. The pore size evolution of RHHCs with increasing temperature from
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1100 to 1500 oC is owing to the reduction of functional groups and the stronger stacking interaction
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between graphene layers.
High-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) techniques were adopted for a more careful investigation of the RHHC microstructure. The HRTEM images of RHHCs are demonstrated in Fig. 2c–e. As expected for hard carbons, an apparent disordered structure can be noticed, further corroborating the amorphous feature of our RHHCs. Upon increasing the pyrolysis temperature, we can clearly see the gradual growth of turbostratic nanocrystallite. At the same time, there is a decrease in the number of
ACCEPTED MANUSCRIPT nanopores encircled by some turbostratic nanocrystallites. The minimal number of nanopores can be seen in RHHC-1500, which might be correlated with the movement of turbostratic nanocrystallites with respect to each other at elevated temperatures. All SAED patterns show
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diffusive diffraction rings as a further indication of turbostratic nanocrystallite architecture. The diffraction rings turn sharper with rising the pyrolysis temperature, meaning the formation of a more ordered architecture. The HRTEM and SAED results are in good agreement with the XRD
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patterns and Raman spectra as well as the pore size distribution from BET measurement.
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Fig. 3a–c presents the cyclic voltammetry (CV) curves of the RHHC-1100, RHHC-1300, and RHHC-1500 electrodes at a scanning rate of 0.1 mV s−1 in the voltage range of 0.01−2 V. The irreversible peaks of RHHC-1300 at 0.45 and 0.92 V and RHHC-1500 at 0.47 V in the initial cathodic scan corresponds to the production of SEI films, while the prominent irreversible region of
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RHHC-1100 between 0.22 and 0.83 V can be assigned to the irreversible sodiation of surface functional groups and the generation of SEI layer [69]. The irreversible area of the RHHC-1300 or RHHC-1500 electrode is less than that of the RHHC-1100 electrode. This explains the higher initial
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CEs for the RHHC-1300 and RHHC-1500 electrodes. Furthermore, a pair of strong redox peaks
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emerge at around 0.02/0.12 V in the RHHC-1100, RHHC-1300, or RHHC-1500 electrode, which can be ascribed to the sodium ion insertion into/extraction from the graphene layers of turbostratic nanocrystallites [70].
Fig. 3d shows the representative initial discharge−charge curves of three RHHC electrodes at 25 mA g−1 (0.032 mA cm−2). The RHHC-1100 displays a low charge capacity of 323 mAh g−1 with a long sloping region in the electrochemical profiles and a relatively poor initial CE of 64%. This is because a lot of oxygen-containing functional groups readily cause the irreversible sodium ion
ACCEPTED MANUSCRIPT storage in the low temperature pyrolytic hard carbon, which is well supported by X-ray photoelectron spectroscopy (XPS) analysis (Fig. S6 and Table S3). Intriguingly, both RHHC-1300 and RHHC-1500 give high reversible sodium storage capacities and moderate initial CEs above 65%
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(Table 1). The RHHC-1300 electrode delivers the highest reversible capacity of 372 mAh g−1 (0.47 mAh cm−2) with a medium initial CE of 66%. RHHC-1500 has a larger initial CE of 68% as compared to RHHC-1300 owing to lower oxygen content and specific surface area, which diminish
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the irreversible sodium storage. Comparing the discharge and charge profiles of RHHC-1300 with
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RHHC-1500, one can clearly see that RHHC-1300 has a larger capacity contribution from the sloping region while RHHC-1500 possesses a larger capacity contribution from the plateau region. In addition, RHHC-1300 represents a slightly higher sodiation potential than that of RHHC-1500 based on their discharge curves. The discrepancy in sodium storage behavior in the sloping and
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plateau regions is originated from the difference in microstructure between two hard carbons. RHHC-1500 holds less defected sites and larger turbostratic nanocrystallite size, suggesting that the sloping region can be attributed to the sodium storage in defected sites while the plateau region can
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be ascribed to the sodium storage in between turbostratic graphene sheets.
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The rate capability was also surveyed to assess the kinetic performance of RHHCs. The results are demonstrated in Fig. 3e. The rate property increases first and then degenerates upon increasing the carbonization temperature, which is well consistent with the reported results [45]. Amusingly, RHHC-1300 exhibits a higher rate capability as compared to the kelp-derived hard carbon pyrolyzed at 1300 oC [71], which means that the larger interlayer spacing of RHHC is more favorable for facilitating the transportation of sodium ions and therefore enhancing the rate performance. Moreover, the rate capability of RHHC-1300 is better than that of RHHC-1100 or
ACCEPTED MANUSCRIPT RHHC-1500. As shown in Fig. 3e, at current densities of 500 and 1000 mA g−1, the average specific capacities of RHHC-1300 are 265 and 166 mAh g−1, respectively. Fig. 3f illustrates the cycling performances of RHHC-1100, RHHC-1300, and RHHC-1500 at 25 mA g−1 for 100 cycles. All the
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RHHC electrodes manifest stable cycling properties, especially RHHC-1100 has lower capacity decline during cycling, which should be related to the unique adsorption−desorption sodium ion storage behavior [72]. The cycling performances of RHHC-1300 and RHHC-1500 are influenced
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by the pyrolysis temperature owing to a variation in the size and interlayer distance of turbostratic
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nanocrystallites and the oxygen content, which leads to the change of reversible capacity in the plateau and sloping regions. RHHC-1300 maintains the highest specific capacity of 346 mAh g−1 after 100 cycles, corresponding to a 93% capacity retention based on the first-cycle reversible capacity. Noticeably, RHHC-1300 shows efficient sodium storage capability in comparison with
SIBs.
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previously reported carbon materials (Table S5), which make it an appealing anode material for
To understand the high rate capability of RHHC-1300, the redox pseudocapacitive contribution
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in the RHHC-1300 electrode is analyzed by separating the capacitive capacity and the
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diffusion-controlled capacity. The reaction kinetics of RHHC-1300 with sodium ions can be investigated by CV profiles at different scan rates (v) (Fig. S7a). The current response (i) at a fixed potential (V) can be expressed as a summation of capacitive contribution (k1v) and diffusion-dominated reaction (k2v1/2) according to the following equation: i(V) = k1v + k2v1/2[73]. By determining both k1 and k2 constants, the fraction of the current from surface capacitance and diffusion contributions can be differentiated. A representative profile (Fig. S7b) show the capacitive contribution (the red section) is prominent in the overall capacity. As illustrated in Fig. S7c, at scan
ACCEPTED MANUSCRIPT rates of 0.2, 0.4, 0.6, 0.8, and 1.0 mV s−1, the fractions of capacitive capacity are 39.6%, 49.3%, 59.9%, 65.8%, and 74.5%, respectively. The kinetics analysis indicates that the diffusion contribution are inhibited at high scan rates, which might be responsible for the high rate
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performance at high current densities. Ex situ XRD patterns of the RHHC-1300 electrode were recorded at various depths of discharge and charge at 25 mA g−1 to better reveal the structural evolution mechanism during
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sodiation/desodiation, as demonstrated in Fig. 4. As being discharged to 0.56 V (B), no apparent
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new peak is detected in the XRD pattern, suggesting no phase change occurring. As the continuous discharging to 0.08 V (C), the peak at 23.6o attenuates and a new diffraction peak appearing at 20.5o can be assigned to intercalation of sodium ions into turbostratic nanocrystallites of RHHC-1300 [67, 71]. The emerging XRD peak of the RHHC-1300 electrode progressively becomes stronger upon
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discharging to 0.03 V (D). When fully discharging to the cutoff voltage of 0.01 V (E), a coexistence of much weakened initial peak and largely increased new diffraction peak indicates an almost whole transformation from pristine RHHC-1300 to an intercalation compound. It is worth mentioning that
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the turbostratic structure of this hard carbon material can nearly restore from Na+-intercalated
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RHHC-1300 to its original state during the desodiation process as illustrated by the XRD spectra (E−I) in Fig. 4b, verifying that the sodium ion insertion/extraction processes take place with reversible phase transformation of this hard carbon. Note that the peak at ∼20.5° produced after the first sodiation partially maintains after the following desodiation, which suggests that part of the RHHC structure preserves dilating with sodium ions trapped in the structure. This result is supported by the moderate first-cycle CE. To elucidate the kinetic discrepancy among RHHCs particularly with relation to their different
ACCEPTED MANUSCRIPT rate performances, we have used galvanostatic intermittent titration technique (GITT) to study the apparent diffusion coefficient of sodium ions in the RHHC electrodes with a pulse current of 20 mA g−1 for 30 min between rest intervals for 60 min. According to the Fick's second law of diffusion,
[74]: 2
2
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4 m V ∆E DNa + = B M s πτ M B S ∆Et
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the diffusivity coefficient of sodium ions (DNa+) can be estimated based on the ensuing equation
where τ is the pulse duration, mB is the mass of active material, MB is the molar mass of carbon, VM
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is the molar volume, and S is the active surface area of the RHHC electrode. ∆Es and ∆Et can be gained from the GITT curves (Fig. 5a,b and Fig. S8). As demonstrated in Fig. 5c,d, sodium ion diffusivity coefficient attains a magnitude of 10−11 cm2 s−1. The diffusivity coefficient variation in the RHHC-1100, RHHC-1300, and RHHC-1500 electrodes is similar during the sodium ion
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intercalation/deintercalation process. During the sodiation process (Fig. 5c), the apparent diffusion coefficient of RHHCs first diminishes slowly followed by a fast decline. However, the diffusivity
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coefficient partially recovers before the cutoff voltage. Fig. 5c shows that apparent diffusion coefficient along with the sloping potentials is much higher than that with the plateau. It is
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reasonable to propose that the surface of the turbostratic nanocrystallites is easier to reach than the interlayer space. As these surface sites are progressively occupied, sodium ions begin to diffuse into the turbostratic nanocrystallites. However, the sodium ions have to conquer the repulsion from the previously anchored sodium ions on defect sites to intercalate in between the turbostratic graphene sheets, which accounts for the evident reduction in apparent diffusion coefficient in the plateau section. As the diffusion distance to sodium storage sites continues lengthening, the diffusivity values gradually become smaller and smaller. After attaining the lowest value of 6.7 × 10−16 cm2 s−1
ACCEPTED MANUSCRIPT the apparent diffusion coefficient reverse to increase, suggesting that the last storage mechanism of RHHCs is probably the adsorption of sodium atoms on the internal pore surfaces. During the desodiation process (Fig. 5d), the apparent diffusion coefficient variation is nearly opposite to that
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happening in the sodiation process, indicating that the sodiation and desodiation processes of RHHCs are highly reversible. In addition, we can notice that RHHC-1300 exhibits superior sodium ion apparent diffusivity coefficient in comparison with RHHC-1100 or RHHC-1500, which explains
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the slowest fading of reversible capacity at high current densities.
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The structural evolution and sodium storage mechanism in RHHCs were further examined using ex situ HRTEM and Auger electron spectroscopy (AES) techniques. HRTEM images before and after sodiation (Fig. 6a,b) demonstrate disordered carbon structures with two different interlayer spacing values of 3.95 and 4.21 Å, indicating that there is evident insertion of sodium ions in
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between the turbostratic graphene sheets [67, 71]. Similar phenomena are also observed in RHHC-1100 and RHHC-1500 (Fig. S9 and S10). The fringes and nanopores of RHHC-1300 become vague after the first discharge process, which is due to the generation of sodium storage.
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Moreover, SEM and TEM images of the RHHC-1300 electrode after 100 cycles displays a
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relatively uniform SEI film with a thickness of around 85 nm, which is thinner than that of the RHHC-1100 or RHHC-1500 electrode (Fig. S11−13). Fig. S14a compares the Nyquist plots for the RHHC electrodes. Obviously, the RHHC-1300 electrode displays a lower SEI film resistance than does the RHHC-1100 or RHHC-1500 electrode (45.5 vs 76.4 or 71.9 Ω, Table S4) based on the modified Randles equivalent circuit illustrated in the inset of Fig. S14a. This result suggests that the RHHC-1300 electrode has a more stable electrode interface, leading to the less side reactions and higher reversible capacity compared with other RHHC electrodes. Furthermore, the curves in the
ACCEPTED MANUSCRIPT low frequency region correspond to Warburg impedance with different slopes, suggesting the electrodes present different solid-state ion diffusion behaviors. Fig. S14b demonstrates the fitting results of the coefficient of Warburg impedance (σw), in which there is a lower Zre versus ω−1/2
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slope of RHHC-1300 than that of RHHC-1100 or RHHC-1500. With the above concerns, the diffusivity coefficient of sodium ions was calculated according to the following equation[75].
R 2T 2 2 A2 n 4 F 4CNa + 2σ w2
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DNa + =
Where R is the gas constant, T is the absolute temperature, A is the surface area of the electrode, n is
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the number of electrons transferred in the electrochemical redox reaction, F is the Faraday constant, and CNa+ is the molar concentration of sodium ions in the hard carbon. As a result, RHHC-1300 possesses a sodium ion diffusion coefficient of 1.24 × 10−13 cm2 s−1, which is slightly higher than that of RHHC-1100 (1.05 × 10−13 cm2 s−1) or RHHC-1500 (1.15 × 10−13 cm2 s−1) (Table S4). This
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result is completely consistent with the GITT analysis. Fig. 6c shows the ex situ AES analysis of C, O, and Na in the RHHC-1300 electrode at the fully discharged state with a sputtering rate of 13 nm
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min−1 (Fig. S15). In the initial process of argon ion sputtering, we observed rapid drop in sodium and oxygen contents and fast increase in carbon content, indicating that the SEI layer is mainly
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formed on the surface of the RHHC-1300 electrode. As the sputtering depth increases, the change in sodium content becomes relatively slow, suggesting that sodium ions homogeneously intercalate into turbostratic graphene sheets and nanopores, which is well supported by the elemental mapping measurement (Fig. S16). Owing to the lower binding energies for defect and nanovoid surface adsorption and the higher repulsion for sodium ion insertion into the turbostratic nanocrystallites, it is logical to conclude that the storage mechanism of sodium ions in RHHC consists of three parts (Fig. 6d): (1) sodium ion storage at defect sites corresponds to the sloping region above 0.12 V; (2)
ACCEPTED MANUSCRIPT storage of sodium ions between the graphene sheets inside the turbostratic nanocrystallites is associated with the plateau region around 0.1 V; (3) sodium ion filling of the nanovoids among the turbostratic nanocrystallites is related to the end region approaching 0.01 V.
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We further investigated the practical application prospect of RHHC through the evaluation of sodium-ion full batteries with Na3V2(PO4)2F3/C (Fig. S17−19) and RHHC-1300 as cathode and anode materials, respectively[63]. Fig. 7a displays the representative charge−discharge profiles of
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the Na3V2(PO4)2F3/C//RHHC-1300 full cell at 25 mA g−1 between 1.6 and 4.3 V. Based on the mass
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of RHHC-1300, the full battery exhibits a stable capacity of 258 mAh g−1 after the activation cycles with an initial CE of 50% and an average output voltage of 3.4 V. The energy density of this full cell based on the total mass of cathode and anode active materials is computed to be 185 Wh kg−1. Furthermore, the full cell also demonstrates good rate capability with a high specific capacity of 173
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mAh g−1 even at 100 mA g−1 and remarkable cycling stability with an 85% capacity retention after 100 cycles (Fig. 7b,c). These superb electrochemical performance of the full cell implies that
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RHHC-1300 has an appealing sodium storage capability as the anode material for SIBs. 4. Conclusions
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In summary, we have synthesized hard carbons from a biomass waste rice husk by a convenient route. The materials were fabricated through the same acid treatment and subsequent pyrolyzing at different temperatures from 1100 to 1500 oC and showed different electrochemical behaviors. These discrepancies
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the
sodium
storage
capability
were
studied
in
detail
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microstructure-electrochemical property relationship. RHHC-1300 delivers the highest reversible capacity of 372 mAh g−1 with representative electrochemical performance for hard carbons in SIBs. In addition to the relevance between microstructure and electrochemical property as identified by
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from the low sodium ion apparent diffusion coefficient of the RHHC materials. The practical viability of RHHC in full cells was further proved by assembling full batteries with Na3V2(PO4)2F3/C and RHHC-1300 as the cathode and anode materials, respectively. We have
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achieved a high-energy density of 185 Wh kg−1, superior rate capability, and stable cycling
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performance. Our work renders a simple model for understanding the electrochemical sodium storage behaviors in biomass-derived hard carbons. Acknowledgements
This work was supported by the National Natural Science Foundation of China (Grant No.
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51577094 and 21503112), the 100 Talents Program of Nanjing Normal University, the Priority Academic Program Development of Jiangsu Higher Education Institutions, and the Program of
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Jiangsu Collaborative Innovation Center of Biomedical Functional Materials.
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ACCEPTED MANUSCRIPT Fig. and Table Captions
Fig. 1. SEM images of (a) the inner surface and (b) the outer surface of a rice husk. (c) SEM image
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and (d) high-magnification SEM image of RHHC-1300. Table 1. Physical parameters and electrochemical properties for RHHCs.
Fig. 2. (a) XRD patterns and (b) Raman spectra of RHHCs. HRTEM images of (c) RHHC-1100, (d)
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RHHC-1300, and (e) RHHC-1500. The insets of (c–e) show the SAED patterns of corresponding
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samples.
Fig. 3. CV curves of (a) RHHC-1100, (b) RHHC-1300, and (c) RHHC-1500, (d) Galvanostatic discharge/charge profiles for the first cycle of RHHCs. (e) Rate capabilities of RHHCs. (f) Cycling performances of RHHCs.
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Fig. 4. (a) Discharge/charge profiles of RHHC-1300, the marks (A−I) show the various depths of discharge/charge. (b) XRD patterns of the RHHC-1300 electrode at different discharge/charge status.
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Fig. 5. GITT curves of RHHCs for (a) sodiation and (b) desodiation of the second cycle.
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Corresponding sodium ion apparent diffusion coefficients of RHHCs for (c) sodiation and (d) desodiation.
Fig. 6. HRTEM images of (a) fresh RHHC-1300 and (b) the RHHC-1300 sample after discharging to 0.01 V. (c) AES depth profiles of the RHHC-1300 electrode after discharging to 0.01 V. (d) Potentiogram and schematic representation of sodium storage mechanism. Fig. 7. (a) Charge and discharge profiles for the Na3V2(PO4)2F3/C//RHHC-1300 full cell. (b) Rate capability. (c) Cycling performance.
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Fig. 1.
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o
Interlayer
Lc (nm)
La (nm)
distance (Å)
RHHC-1100
1100
4.03
1.16
11.0
RHHC-1300
1300
3.95
1.23
12.0
RHHC-1500
1500
3.84
1.31
12.5
2
−1
area (m g ) 2.68
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Reversible
Initial CE
−1
capacity (mAh g )
(%)
323
64
0.27
372
66
0.23
327
68
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temp ( C)
BET surface
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Pyrolysis
Sample No.
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Table 1.
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Fig. 2.
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Fig. 3.
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Fig. 4.
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S0008-6223(17)31168-5
DOI:
10.1016/j.carbon.2017.11.054
Reference:
CARBON 12584
To appear in:
Carbon
Received Date: 21 September 2017 Revised Date:
13 November 2017
Accepted Date: 19 November 2017
Please cite this article as: Q. Wang, X. Zhu, Y. Liu, Y. Fang, X. Zhou, J. Bao, Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries, Carbon (2017), doi: 10.1016/ j.carbon.2017.11.054. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphical abstract
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Rice husk-derived hard carbons as high-performance anode materials for sodium-ion batteries
a
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Qiaoqiao Wang a, Xiaoshu Zhu b, Yuhan Liu a, Yuyan Fang a, Xiaosi Zhou a, *, Jianchun Bao a Jiangsu Key Laboratory of New Power Batteries, Jiangsu Collaborative Innovation Center of
Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal
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University, Nanjing 210023, P. R. China
Center for Analysis and Testing, Nanjing Normal University, Nanjing 210023, P. R. China
*
Corresponding author.
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E-mail address: [email protected] (X. Zhou).
ACCEPTED MANUSCRIPT Abstract Sodium-ion batteries (SIBs) have drawn ever-increasing attention for scalable electrical energy storage owing to the inexhaustible sources and wide distribution of sodium. However, to develop
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feasible anode materials still remains a great challenge for the practical application of SIBs. Here, we report hard carbons derived from a plentiful and deserted biomass of rice husk through a facile acid treatment and subsequent pyrolysis. The investigation illustrates that the electrochemical
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properties of the rice husk-derived hard carbons (RHHCs) are significantly influenced by the
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pyrolysis temperature because of the discrepancy in their microstructure. The RHHC pyrolyzed at 1300 °C (RHHC-1300) shows the highest reversible capacity of 372 mAh g−1 and good cycling stability and rate performance due to its large interlayer distance and suitable oxygen content. Moreover, full sodium-ion batteries are assembled to examine the application prospect using
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Na3V2(PO4)2F3/C and RHHC-1300 as cathode and anode materials, respectively, delivering a high-energy density of 185 Wh kg−1 and stable cycling performance. This work could intensify the
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fundamental understanding of the sodium storage mechanism in biomass-derived hard carbons.
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Keywords: Rice husk; Hard carbon; Sodium-ion battery; Anode; Storage mechanism
ACCEPTED MANUSCRIPT 1. Introduction As the most widely used energy storage system, lithium-ion batteries (LIBs) have dominated the power source market for electric vehicles and portable electronic devices due to their high energy
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density and long cycle lifespan [1]. However, the scarcity and uneven geographical distribution of lithium resource seriously hamper the large-scale applications of LIBs to energy storage on the grid [2]. Benefiting from the copious reserve and lower price of sodium with comparable physical and
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chemical properties to lithium, sodium-ion batteries (SIBs) have been considered as one of the most
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prospective candidates for large-scale stationary energy storage [3, 4]. In comparison with lithium ions, the radius of sodium ions is much larger (1.02 versus 0.76 Å) so that one cannot randomly choose electrode materials for SIBs from those for LIBs [5]. Hence, it is crucial to search for qualified host materials to supply enough space for sodium ions storage and transportation [6].
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Recently, a large number of layered oxides and polyanion compounds have been verified as main cathode materials for SIBs [7-9]. In terms of anode materials, commercial graphite cannot afford enough interlayer space to accommodate larger sodium ions and thus present low sodium storage
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[12].
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capability [10, 11]; in this aspect, it is extremely important to find suitable anode materials for SIBs
While many kinds of anode materials, such as carbon materials [13-28], alloys [29-34], metal oxides and sulfides [35-37], and organic compounds [38], have been reported to possess sodium storage property, most of them show poor cycle life or low specific capacity. Specifically, the structural disintegration caused by huge volume variation results in the loss of electrical contact and capacity decline in the case of alloy-type materials [39, 40]. The poor initial Coulombic efficiency (CE) of metal oxides or sulfides and the low sodium storage performance of organic compounds
ACCEPTED MANUSCRIPT may inhibit their practical applications in SIBs [41, 42]. Among all the anode materials, hard carbon with randomly oriented turbostratic nanocrystallites has demonstrated the most promising sodium storage performance [43-49]. Moreover, hard carbons derived from various biomass precursors and
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their fundamental sodium storage mechanisms have been widely studied recently [50-57]. Nonetheless, there are still two significant obstacles before realizing industrial production of hard
high price and low carbonization yield of precursors.
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carbons as anode materials for SIBs: unsatisfactory sodium storage capacity and high cost owing to
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Rice husk is a kind of agricultural deserted byproduct and is copiously existed in the world [58-62]. The abundant content of cellulose and hemicellulose makes rice husk a proper carbon source to fabricate hard carbons. Herein, we report a simple yet efficient approach to synthesize hard carbons by utilizing abandoned rice husk. The electrochemical properties of the resultant hard
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carbons as anode materials for SIBs are found to be considerably affected by pyrolysis temperature. The rice husk-derived hard carbon carbonized at 1300 oC (RHHC-1300) shows promising sodium storage performance with a high reversible capacity of 372 mAh g−1, a medium initial CE of 66%,
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and good cycling stability and rate performance. The practical application prospect of RHHC-1300
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is further corroborated by using full sodium-ion batteries with Na3V2(PO4)2F3/C as cathode material, exhibiting a high energy density of 185 Wh kg−1 and excellent cycling performance and rate capability. This contribution provides a route for exploiting deserted agricultural biomass to prepare hard carbon anode materials for SIBs. 2. Experimental 2.1. Materials Synthesis The compressed rice husks were bought from Taobao (3 kg in pack). Prior to the synthesis
ACCEPTED MANUSCRIPT process, impurities in the rice husks were removed by washing in 6 M HCl for 6 h and 10% HF for 12 h, rinsing with water, and drying at 120 oC overnight in a vacuum oven. For the pyrolysis process, 2 g of purified sample was loaded in a tube furnace and carbonized for 2 h under an argon
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flow with a heating rate of 5 oC min−1. The pyrolysis temperatures were 1100, 1300, and 1500 oC, and the obtained rice husk-derived hard carbons were denoted as RHHC-1100, RHHC-1300, and RHHC-1500, respectively.
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2.2. Materials Characterization
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Scanning electron microscopy (SEM) images were taken by a JEOL JSM-7600F scanning electron microscope operated at 10 kV. X-ray diffraction (XRD) patterns were collected using a Rigaku SmartLab diffractometer with Cu Kα irradiation (λ = 1.5406 Å). Raman spectra were determined on a Labram HR800 with a laser wavelength of 514 nm. Nitrogen sorption
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measurements were carried out on an ASAP 2050 surface area-pore size analyzer. High-resolution transmission electron microscopy (HRTEM) images were obtained by a JEOL JEM-2100F transmission electron microscope operated at 200 kV. X-ray photoelectron spectroscopy (XPS) was
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acquired on an ESCALab250Xi electron spectrometer from VG Scientific using 300 W Al Kα
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radiation. Scanning transmission electron microscopy (STEM) as well as elemental mapping measurements were performed on the JEOL JEM-2100F transmission electron microscope equipped with a Thermo Fisher Scientific energy-dispersive X-ray spectrometer. 2.3. Electrochemical Measurements Electrochemical tests were conducted using CR2032 coin cells. The working electrodes were prepared by coating the mixed slurry of RHHC, carbon black (Super-P), and polyvinyl fluoride (PVDF) binder in N-methyl-2-pyrrolidone with a weight ratio of 8:1:1 onto copper foil, and then
ACCEPTED MANUSCRIPT dried at 80 °C in a vacuum oven for 12 h. It is noteworthy that the RHHC particles were ground for 30 minutes using a mortar and pestle prior to direct use as electrode materials. The loading mass of active material was controlled between 1.0 and 1.5 mg cm−2. Glass fiber (GF/D) from Whatman and
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sodium metal were used as separators and counter electrodes, respectively. The electrolyte was composed of 1 M NaClO4 in a mixture of ethylene carbonate and diethyl carbonate (1:1 in volume). All coin cells were assembled in an argon-filled glovebox (H2O, O2 < 0.1 ppm, MBraun, Germany).
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Galvanostatic discharge and charge tests were carried out on Land CT2001A multichannel battery
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testing systems in the fixed voltage range of 0.01−2 V vs Na/Na+ at various current densities. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were conducted on a PARSTAT 4000 electrochemical workstation. CV was recorded at a scan rate of 0.1 mV s−1, and EIS was determined by applying a sine wave with an amplitude of 10.0 mV over the
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frequency range from 100 kHz to 100 mHz. Sodium-ion full batteries were fabricated using RHHC-1300 as the anode material and Na3V2(PO4)2F3/C as the cathode material in CR2032 coin-type cells. The Na3V2(PO4)2F/C nanocomposite was synthesized through a previously reported
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sol−gel method[63]. The weight ratio of the two electrodes (negative/positive) was 1:4. The full
g−1.
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cells were charged and discharged in the voltage range of 1.6−4.3 V at a current density of 25 mA
2.4. Ex situ XRD, HRTEM, and AES Analysis For the electrodes employed for ex situ XRD, HRTEM, and AES measurements, carbon black was not used, while RHHC-1300 was increased to 90 wt %. The cycled cells were disassembled in the argon-filled glovebox, and the RHHC-1300 electrodes were collected and washed with dimethyl carbonate (DMC) several times. The naturally dried RHHC-1300 electrode was put onto a glass
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operated at 5 kV. As for ex situ HRTEM measurement, the rinsed electrode material was dispersed in DMC under the protection of argon atmosphere by sonication. The HRTEM samples were prepared by pipetting a few microliters of the resulting dispersion onto carbon grids in the
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argon-filled glovebox. The naturally dried carbon grids were placed in sample holder for ex situ
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HRTEM experiment. 3. Results and discussion
Fig. 1a,b displays scanning electron microscopy (SEM) images of the inner and outer surfaces of the starting rice husk precursor. The SEM image shows that the inner surface is smooth and the
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outer surface is rugged. After acid treatment and subsequent pyrolysis at elevated temperature under argon atmosphere, the rice husks first become yellow, shrink in size, and then turn black, implying
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its successful removal of impurities and carbonization (Fig. S1)[64]. When rice husks are carbonized at T oC, the resulting rice husk-derived hard carbon is designated as RHHC-T (Table 1).
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The annealed rice husks are found to possess irregular granular shape with a size distribution range of 3−25 µm and a dense structure (Fig. 1c,d and Fig. S2). An obvious decrease in the size of sample is observed during carbonization. The carbon yields of rice husks at 1100 oC, 1300 oC, and 1500 oC are around 26%, 24%, and 23%, respectively. The dense structure of rice husk-derived hard carbon (RHHC) helps to form a stable SEI film and promotes the transportation of electrons, boosting the electrochemical property of RHHC. To analyze the microstructure of RHHC, X-ray diffraction (XRD) and Raman spectroscopy were
ACCEPTED MANUSCRIPT performed and the results are demonstrated in Fig. 2a,b. All XRD patterns show broad peaks at 2θ around 22.5o and 43.6o, which correspond to the (002) and (101) planes of disordered carbon structure. There is no evident variation in the peak shape of RHHC with increasing the pyrolysis
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temperature, indicating the characteristic of hard carbon. The (002) peak position diverts to a higher angle with rising the carbonization temperature, suggesting the local structural evolution to a smaller d-spacing. As shown in Table 1, the interlayer spacing shrinks from 4.03 Å for RHHC-1100
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to 3.95 and 3.84 Å for RHHC-1300 and RHHC-1500, respectively. In Raman spectra (Fig. 2b), two
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detached characteristic peaks at ~1350 and ~1606 cm−1 can be assigned to D band and G band [47, 65], respectively, indicating the amorphous architecture of RHHCs. The full widths at half maximum of D and G bands in the Raman spectra diminish with increasing pyrolysis temperature, which further reveals the evolution of local long-range ordered structure, i.e., turbostratic
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nanocrystallite size. The turbostratic nanocrystallite size along the c-axis (Lc) of a rice husk-derived hard carbon can be computed by the Scherrer equation. On average, there are about 2.9, 3.1, and 3.4 layers of graphene sheets in the nanocrystallites of RHHC-1100, RHHC-1300, and RHHC-1500
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(Table S1), respectively. Detailed calculation can be found in Fig. S3. Furthermore, the
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nanocrystallite size along the ab-plane (La) was also calculated through Raman spectrum and the equation below [66]:
4 IG La = ( 2.4 × 10−10 ) λnm ID
where λnm is the laser wavelength, ID represents the integrated D band, and IG indicates the integrated G band (Fig. S4). The Raman results show that RHHC-1100 has the smallest turbostratic nanocrystallite size with La value of 11.0 nm, while the value rises to 12.0 nm for RHHC-1300, and 12.5 nm for RHHC-1500 (Table S2). Overall, the XRD and Raman analysis show that both
ACCEPTED MANUSCRIPT d-spacing and turbostratic nanocrystallite size of RHHCs are strongly affected by the carbonization temperature. Nitrogen adsorption−desorption isotherms in Fig. S5 displays the Brunauer−Emmett−Teller
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(BET) surface area and pore size distribution of RHHC. The BET surface area is closely associated with the pyrolysis temperature. The computed surface area diminishes with the increase of carbonization temperature, RHHC-1100 possesses the highest BET surface area of 2.68 m2 g−1
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while RHHC-1300 and RHHC-1500 encompass lower surface areas of 0.27 and 0.23 m2 g−1,
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respectively. These small surface areas of RHHC-1300 and RHHC-1500 could lead to finite SEI layers and hence enhance the initial CE [67], in accordance with the electrochemical results (Table 1). Besides, the pore size distribution in Fig. S5b also demonstrates different pore size relying on the annealing temperature. In terms of RHHC-1100, a few micropores and mesopores are existing
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with a total pore volume of 3.35 × 10−3 cm3 g−1, while these pore volumes remarkably decrease with increasing the heat treatment temperature. The pore size distribution of RHHC-1300 or RHHC-1500 is basically undetectable[68]. The pore size evolution of RHHCs with increasing temperature from
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1100 to 1500 oC is owing to the reduction of functional groups and the stronger stacking interaction
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between graphene layers.
High-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) techniques were adopted for a more careful investigation of the RHHC microstructure. The HRTEM images of RHHCs are demonstrated in Fig. 2c–e. As expected for hard carbons, an apparent disordered structure can be noticed, further corroborating the amorphous feature of our RHHCs. Upon increasing the pyrolysis temperature, we can clearly see the gradual growth of turbostratic nanocrystallite. At the same time, there is a decrease in the number of
ACCEPTED MANUSCRIPT nanopores encircled by some turbostratic nanocrystallites. The minimal number of nanopores can be seen in RHHC-1500, which might be correlated with the movement of turbostratic nanocrystallites with respect to each other at elevated temperatures. All SAED patterns show
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diffusive diffraction rings as a further indication of turbostratic nanocrystallite architecture. The diffraction rings turn sharper with rising the pyrolysis temperature, meaning the formation of a more ordered architecture. The HRTEM and SAED results are in good agreement with the XRD
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patterns and Raman spectra as well as the pore size distribution from BET measurement.
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Fig. 3a–c presents the cyclic voltammetry (CV) curves of the RHHC-1100, RHHC-1300, and RHHC-1500 electrodes at a scanning rate of 0.1 mV s−1 in the voltage range of 0.01−2 V. The irreversible peaks of RHHC-1300 at 0.45 and 0.92 V and RHHC-1500 at 0.47 V in the initial cathodic scan corresponds to the production of SEI films, while the prominent irreversible region of
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RHHC-1100 between 0.22 and 0.83 V can be assigned to the irreversible sodiation of surface functional groups and the generation of SEI layer [69]. The irreversible area of the RHHC-1300 or RHHC-1500 electrode is less than that of the RHHC-1100 electrode. This explains the higher initial
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CEs for the RHHC-1300 and RHHC-1500 electrodes. Furthermore, a pair of strong redox peaks
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emerge at around 0.02/0.12 V in the RHHC-1100, RHHC-1300, or RHHC-1500 electrode, which can be ascribed to the sodium ion insertion into/extraction from the graphene layers of turbostratic nanocrystallites [70].
Fig. 3d shows the representative initial discharge−charge curves of three RHHC electrodes at 25 mA g−1 (0.032 mA cm−2). The RHHC-1100 displays a low charge capacity of 323 mAh g−1 with a long sloping region in the electrochemical profiles and a relatively poor initial CE of 64%. This is because a lot of oxygen-containing functional groups readily cause the irreversible sodium ion
ACCEPTED MANUSCRIPT storage in the low temperature pyrolytic hard carbon, which is well supported by X-ray photoelectron spectroscopy (XPS) analysis (Fig. S6 and Table S3). Intriguingly, both RHHC-1300 and RHHC-1500 give high reversible sodium storage capacities and moderate initial CEs above 65%
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(Table 1). The RHHC-1300 electrode delivers the highest reversible capacity of 372 mAh g−1 (0.47 mAh cm−2) with a medium initial CE of 66%. RHHC-1500 has a larger initial CE of 68% as compared to RHHC-1300 owing to lower oxygen content and specific surface area, which diminish
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the irreversible sodium storage. Comparing the discharge and charge profiles of RHHC-1300 with
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RHHC-1500, one can clearly see that RHHC-1300 has a larger capacity contribution from the sloping region while RHHC-1500 possesses a larger capacity contribution from the plateau region. In addition, RHHC-1300 represents a slightly higher sodiation potential than that of RHHC-1500 based on their discharge curves. The discrepancy in sodium storage behavior in the sloping and
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plateau regions is originated from the difference in microstructure between two hard carbons. RHHC-1500 holds less defected sites and larger turbostratic nanocrystallite size, suggesting that the sloping region can be attributed to the sodium storage in defected sites while the plateau region can
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be ascribed to the sodium storage in between turbostratic graphene sheets.
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The rate capability was also surveyed to assess the kinetic performance of RHHCs. The results are demonstrated in Fig. 3e. The rate property increases first and then degenerates upon increasing the carbonization temperature, which is well consistent with the reported results [45]. Amusingly, RHHC-1300 exhibits a higher rate capability as compared to the kelp-derived hard carbon pyrolyzed at 1300 oC [71], which means that the larger interlayer spacing of RHHC is more favorable for facilitating the transportation of sodium ions and therefore enhancing the rate performance. Moreover, the rate capability of RHHC-1300 is better than that of RHHC-1100 or
ACCEPTED MANUSCRIPT RHHC-1500. As shown in Fig. 3e, at current densities of 500 and 1000 mA g−1, the average specific capacities of RHHC-1300 are 265 and 166 mAh g−1, respectively. Fig. 3f illustrates the cycling performances of RHHC-1100, RHHC-1300, and RHHC-1500 at 25 mA g−1 for 100 cycles. All the
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RHHC electrodes manifest stable cycling properties, especially RHHC-1100 has lower capacity decline during cycling, which should be related to the unique adsorption−desorption sodium ion storage behavior [72]. The cycling performances of RHHC-1300 and RHHC-1500 are influenced
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by the pyrolysis temperature owing to a variation in the size and interlayer distance of turbostratic
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nanocrystallites and the oxygen content, which leads to the change of reversible capacity in the plateau and sloping regions. RHHC-1300 maintains the highest specific capacity of 346 mAh g−1 after 100 cycles, corresponding to a 93% capacity retention based on the first-cycle reversible capacity. Noticeably, RHHC-1300 shows efficient sodium storage capability in comparison with
SIBs.
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previously reported carbon materials (Table S5), which make it an appealing anode material for
To understand the high rate capability of RHHC-1300, the redox pseudocapacitive contribution
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in the RHHC-1300 electrode is analyzed by separating the capacitive capacity and the
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diffusion-controlled capacity. The reaction kinetics of RHHC-1300 with sodium ions can be investigated by CV profiles at different scan rates (v) (Fig. S7a). The current response (i) at a fixed potential (V) can be expressed as a summation of capacitive contribution (k1v) and diffusion-dominated reaction (k2v1/2) according to the following equation: i(V) = k1v + k2v1/2[73]. By determining both k1 and k2 constants, the fraction of the current from surface capacitance and diffusion contributions can be differentiated. A representative profile (Fig. S7b) show the capacitive contribution (the red section) is prominent in the overall capacity. As illustrated in Fig. S7c, at scan
ACCEPTED MANUSCRIPT rates of 0.2, 0.4, 0.6, 0.8, and 1.0 mV s−1, the fractions of capacitive capacity are 39.6%, 49.3%, 59.9%, 65.8%, and 74.5%, respectively. The kinetics analysis indicates that the diffusion contribution are inhibited at high scan rates, which might be responsible for the high rate
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performance at high current densities. Ex situ XRD patterns of the RHHC-1300 electrode were recorded at various depths of discharge and charge at 25 mA g−1 to better reveal the structural evolution mechanism during
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sodiation/desodiation, as demonstrated in Fig. 4. As being discharged to 0.56 V (B), no apparent
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new peak is detected in the XRD pattern, suggesting no phase change occurring. As the continuous discharging to 0.08 V (C), the peak at 23.6o attenuates and a new diffraction peak appearing at 20.5o can be assigned to intercalation of sodium ions into turbostratic nanocrystallites of RHHC-1300 [67, 71]. The emerging XRD peak of the RHHC-1300 electrode progressively becomes stronger upon
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discharging to 0.03 V (D). When fully discharging to the cutoff voltage of 0.01 V (E), a coexistence of much weakened initial peak and largely increased new diffraction peak indicates an almost whole transformation from pristine RHHC-1300 to an intercalation compound. It is worth mentioning that
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the turbostratic structure of this hard carbon material can nearly restore from Na+-intercalated
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RHHC-1300 to its original state during the desodiation process as illustrated by the XRD spectra (E−I) in Fig. 4b, verifying that the sodium ion insertion/extraction processes take place with reversible phase transformation of this hard carbon. Note that the peak at ∼20.5° produced after the first sodiation partially maintains after the following desodiation, which suggests that part of the RHHC structure preserves dilating with sodium ions trapped in the structure. This result is supported by the moderate first-cycle CE. To elucidate the kinetic discrepancy among RHHCs particularly with relation to their different
ACCEPTED MANUSCRIPT rate performances, we have used galvanostatic intermittent titration technique (GITT) to study the apparent diffusion coefficient of sodium ions in the RHHC electrodes with a pulse current of 20 mA g−1 for 30 min between rest intervals for 60 min. According to the Fick's second law of diffusion,
[74]: 2
2
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4 m V ∆E DNa + = B M s πτ M B S ∆Et
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the diffusivity coefficient of sodium ions (DNa+) can be estimated based on the ensuing equation
where τ is the pulse duration, mB is the mass of active material, MB is the molar mass of carbon, VM
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is the molar volume, and S is the active surface area of the RHHC electrode. ∆Es and ∆Et can be gained from the GITT curves (Fig. 5a,b and Fig. S8). As demonstrated in Fig. 5c,d, sodium ion diffusivity coefficient attains a magnitude of 10−11 cm2 s−1. The diffusivity coefficient variation in the RHHC-1100, RHHC-1300, and RHHC-1500 electrodes is similar during the sodium ion
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intercalation/deintercalation process. During the sodiation process (Fig. 5c), the apparent diffusion coefficient of RHHCs first diminishes slowly followed by a fast decline. However, the diffusivity
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coefficient partially recovers before the cutoff voltage. Fig. 5c shows that apparent diffusion coefficient along with the sloping potentials is much higher than that with the plateau. It is
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reasonable to propose that the surface of the turbostratic nanocrystallites is easier to reach than the interlayer space. As these surface sites are progressively occupied, sodium ions begin to diffuse into the turbostratic nanocrystallites. However, the sodium ions have to conquer the repulsion from the previously anchored sodium ions on defect sites to intercalate in between the turbostratic graphene sheets, which accounts for the evident reduction in apparent diffusion coefficient in the plateau section. As the diffusion distance to sodium storage sites continues lengthening, the diffusivity values gradually become smaller and smaller. After attaining the lowest value of 6.7 × 10−16 cm2 s−1
ACCEPTED MANUSCRIPT the apparent diffusion coefficient reverse to increase, suggesting that the last storage mechanism of RHHCs is probably the adsorption of sodium atoms on the internal pore surfaces. During the desodiation process (Fig. 5d), the apparent diffusion coefficient variation is nearly opposite to that
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happening in the sodiation process, indicating that the sodiation and desodiation processes of RHHCs are highly reversible. In addition, we can notice that RHHC-1300 exhibits superior sodium ion apparent diffusivity coefficient in comparison with RHHC-1100 or RHHC-1500, which explains
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the slowest fading of reversible capacity at high current densities.
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The structural evolution and sodium storage mechanism in RHHCs were further examined using ex situ HRTEM and Auger electron spectroscopy (AES) techniques. HRTEM images before and after sodiation (Fig. 6a,b) demonstrate disordered carbon structures with two different interlayer spacing values of 3.95 and 4.21 Å, indicating that there is evident insertion of sodium ions in
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between the turbostratic graphene sheets [67, 71]. Similar phenomena are also observed in RHHC-1100 and RHHC-1500 (Fig. S9 and S10). The fringes and nanopores of RHHC-1300 become vague after the first discharge process, which is due to the generation of sodium storage.
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Moreover, SEM and TEM images of the RHHC-1300 electrode after 100 cycles displays a
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relatively uniform SEI film with a thickness of around 85 nm, which is thinner than that of the RHHC-1100 or RHHC-1500 electrode (Fig. S11−13). Fig. S14a compares the Nyquist plots for the RHHC electrodes. Obviously, the RHHC-1300 electrode displays a lower SEI film resistance than does the RHHC-1100 or RHHC-1500 electrode (45.5 vs 76.4 or 71.9 Ω, Table S4) based on the modified Randles equivalent circuit illustrated in the inset of Fig. S14a. This result suggests that the RHHC-1300 electrode has a more stable electrode interface, leading to the less side reactions and higher reversible capacity compared with other RHHC electrodes. Furthermore, the curves in the
ACCEPTED MANUSCRIPT low frequency region correspond to Warburg impedance with different slopes, suggesting the electrodes present different solid-state ion diffusion behaviors. Fig. S14b demonstrates the fitting results of the coefficient of Warburg impedance (σw), in which there is a lower Zre versus ω−1/2
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slope of RHHC-1300 than that of RHHC-1100 or RHHC-1500. With the above concerns, the diffusivity coefficient of sodium ions was calculated according to the following equation[75].
R 2T 2 2 A2 n 4 F 4CNa + 2σ w2
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DNa + =
Where R is the gas constant, T is the absolute temperature, A is the surface area of the electrode, n is
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the number of electrons transferred in the electrochemical redox reaction, F is the Faraday constant, and CNa+ is the molar concentration of sodium ions in the hard carbon. As a result, RHHC-1300 possesses a sodium ion diffusion coefficient of 1.24 × 10−13 cm2 s−1, which is slightly higher than that of RHHC-1100 (1.05 × 10−13 cm2 s−1) or RHHC-1500 (1.15 × 10−13 cm2 s−1) (Table S4). This
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result is completely consistent with the GITT analysis. Fig. 6c shows the ex situ AES analysis of C, O, and Na in the RHHC-1300 electrode at the fully discharged state with a sputtering rate of 13 nm
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min−1 (Fig. S15). In the initial process of argon ion sputtering, we observed rapid drop in sodium and oxygen contents and fast increase in carbon content, indicating that the SEI layer is mainly
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formed on the surface of the RHHC-1300 electrode. As the sputtering depth increases, the change in sodium content becomes relatively slow, suggesting that sodium ions homogeneously intercalate into turbostratic graphene sheets and nanopores, which is well supported by the elemental mapping measurement (Fig. S16). Owing to the lower binding energies for defect and nanovoid surface adsorption and the higher repulsion for sodium ion insertion into the turbostratic nanocrystallites, it is logical to conclude that the storage mechanism of sodium ions in RHHC consists of three parts (Fig. 6d): (1) sodium ion storage at defect sites corresponds to the sloping region above 0.12 V; (2)
ACCEPTED MANUSCRIPT storage of sodium ions between the graphene sheets inside the turbostratic nanocrystallites is associated with the plateau region around 0.1 V; (3) sodium ion filling of the nanovoids among the turbostratic nanocrystallites is related to the end region approaching 0.01 V.
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We further investigated the practical application prospect of RHHC through the evaluation of sodium-ion full batteries with Na3V2(PO4)2F3/C (Fig. S17−19) and RHHC-1300 as cathode and anode materials, respectively[63]. Fig. 7a displays the representative charge−discharge profiles of
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the Na3V2(PO4)2F3/C//RHHC-1300 full cell at 25 mA g−1 between 1.6 and 4.3 V. Based on the mass
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of RHHC-1300, the full battery exhibits a stable capacity of 258 mAh g−1 after the activation cycles with an initial CE of 50% and an average output voltage of 3.4 V. The energy density of this full cell based on the total mass of cathode and anode active materials is computed to be 185 Wh kg−1. Furthermore, the full cell also demonstrates good rate capability with a high specific capacity of 173
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mAh g−1 even at 100 mA g−1 and remarkable cycling stability with an 85% capacity retention after 100 cycles (Fig. 7b,c). These superb electrochemical performance of the full cell implies that
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RHHC-1300 has an appealing sodium storage capability as the anode material for SIBs. 4. Conclusions
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In summary, we have synthesized hard carbons from a biomass waste rice husk by a convenient route. The materials were fabricated through the same acid treatment and subsequent pyrolyzing at different temperatures from 1100 to 1500 oC and showed different electrochemical behaviors. These discrepancies
in
the
sodium
storage
capability
were
studied
in
detail
by
microstructure-electrochemical property relationship. RHHC-1300 delivers the highest reversible capacity of 372 mAh g−1 with representative electrochemical performance for hard carbons in SIBs. In addition to the relevance between microstructure and electrochemical property as identified by
ACCEPTED MANUSCRIPT XRD, Raman, XPS, nitrogen adsorption, and HRTEM, we also used ex situ XRD to ascertain the expandable structure and operated GITT to determine the apparent diffusion coefficient of sodium ions in the RHHC electrodes. The results show that the poor rate performance is mainly resulted
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from the low sodium ion apparent diffusion coefficient of the RHHC materials. The practical viability of RHHC in full cells was further proved by assembling full batteries with Na3V2(PO4)2F3/C and RHHC-1300 as the cathode and anode materials, respectively. We have
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achieved a high-energy density of 185 Wh kg−1, superior rate capability, and stable cycling
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performance. Our work renders a simple model for understanding the electrochemical sodium storage behaviors in biomass-derived hard carbons. Acknowledgements
This work was supported by the National Natural Science Foundation of China (Grant No.
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51577094 and 21503112), the 100 Talents Program of Nanjing Normal University, the Priority Academic Program Development of Jiangsu Higher Education Institutions, and the Program of
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References
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Jiangsu Collaborative Innovation Center of Biomedical Functional Materials.
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Fig. 1. SEM images of (a) the inner surface and (b) the outer surface of a rice husk. (c) SEM image
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and (d) high-magnification SEM image of RHHC-1300. Table 1. Physical parameters and electrochemical properties for RHHCs.
Fig. 2. (a) XRD patterns and (b) Raman spectra of RHHCs. HRTEM images of (c) RHHC-1100, (d)
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RHHC-1300, and (e) RHHC-1500. The insets of (c–e) show the SAED patterns of corresponding
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Fig. 3. CV curves of (a) RHHC-1100, (b) RHHC-1300, and (c) RHHC-1500, (d) Galvanostatic discharge/charge profiles for the first cycle of RHHCs. (e) Rate capabilities of RHHCs. (f) Cycling performances of RHHCs.
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Fig. 4. (a) Discharge/charge profiles of RHHC-1300, the marks (A−I) show the various depths of discharge/charge. (b) XRD patterns of the RHHC-1300 electrode at different discharge/charge status.
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Fig. 5. GITT curves of RHHCs for (a) sodiation and (b) desodiation of the second cycle.
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Corresponding sodium ion apparent diffusion coefficients of RHHCs for (c) sodiation and (d) desodiation.
Fig. 6. HRTEM images of (a) fresh RHHC-1300 and (b) the RHHC-1300 sample after discharging to 0.01 V. (c) AES depth profiles of the RHHC-1300 electrode after discharging to 0.01 V. (d) Potentiogram and schematic representation of sodium storage mechanism. Fig. 7. (a) Charge and discharge profiles for the Na3V2(PO4)2F3/C//RHHC-1300 full cell. (b) Rate capability. (c) Cycling performance.
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Fig. 1.
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o
Interlayer
Lc (nm)
La (nm)
distance (Å)
RHHC-1100
1100
4.03
1.16
11.0
RHHC-1300
1300
3.95
1.23
12.0
RHHC-1500
1500
3.84
1.31
12.5
2
−1
area (m g ) 2.68
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Reversible
Initial CE
−1
capacity (mAh g )
(%)
323
64
0.27
372
66
0.23
327
68
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temp ( C)
BET surface
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Pyrolysis
Sample No.
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Table 1.
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Fig. 2.
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Fig. 3.
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Fig. 4.
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Fig. 5.
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Fig. 6.
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Fig. 7.
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