N-doped Carbon Composite Derived from Bamboo

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3D Porous Silicon/N-doped Carbon Composite Derived from Bamboo Charcoal as High Performance Anode Material for Lithium-ion Batteries Congcong Zhang, Xin Cai, Wenyan Chen, Siyuan Yang, Donghui Xu, Yueping Fang, and Xiaoyuan Yu ACS Sustainable Chem. Eng., Just Accepted Manuscript • DOI: 10.1021/ acssuschemeng.8b01189 • Publication Date (Web): 27 Jun 2018 Downloaded from http://pubs.acs.org on June 27, 2018

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ACS Sustainable Chemistry & Engineering

3D Porous Silicon/N-doped Carbon Composite Derived from Bamboo Charcoal as High Performance Anode Material for Lithium-ion Batteries Congcong Zhanga,b, Xin Caia,*, Wenyan Chena, Siyuan Yanga, Donghui Xua, Yueping Fang a, Xiaoyuan Yu a,* a

College of Materials and Energy, South China Agricultural University, Guangzhou, 510642, China. E-mail: [email protected]; caixin2015@ scau.edu.cn b

Guangzhou Institute of Energy Conversion, Chinese Academy of Sciences, Guangzhou 510640, China.

Abstract: The exploitation of renewable biomass resources toward green, high value-added functional nanomaterials is promising. Bamboo is a “natural Si reservoir” that contains considerable amount of silica. Available bamboo charcoal is the carbonated bamboo material with attractive features. Herein, 3D porous silicon is extracted from bamboo charcoal by simple calcination and magnesiothermic reduction. To improve the electrochemical stability, the porous Si is further coated with a layer of N-doped amorphous carbon by carbonizing the polyacrylonitrile precursor. The obtained silicon/ nitrogen-doped carbon composite possesses a 3D porous structure and exhibits significantly improved cycling performance along with high rate capabilities. Optimized Si/N-doped carbon composite delivers a reversible capacity of 603 mAh g-1 after 120 cycles at 200 mA g-1 and a high capacity of 360 mAh g-1 at 1.6 A g-1. It demonstrates that the N-doped amorphous carbon layer can effectively accommodate the volume change of the 3D porous silicon and reduce the

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charge transfer resistance so as to improve the cycling durability of the Si/carbon composite. The overall preparation process of the Si/N-doped carbon composite is economic, environmental-friendly and scalable. This work provides a sustainable solution from bamboo charcoal resource to hierarchically porous silicon/carbon composites for cost-efficient lithium-ion battery anode materials. Keywords: Bamboo charcoal; Lithium-ion batteries; Porous silicon; Silicon/carbon composite; Anode materials

Introduction Our mobile daily life requires portable and high performance energy supply sources, such as cleaner and more efficient rechargeable lithium-ion batteries (LIBs) with higher energy density, lower price and longer lifetime to run diverse mobile electronic devices and the emerging electric vehicles/hybrid vehicles. To further improve the actual energy density of LIBs, the development of high-capacity durable anode materials for new-generation LIBs has drawn extensive and continuous attention 1. Among numerous existing anode materials, silicon exhibits exceptionally high theoretical specific capacity (4200 mAh g-1 for Li22Si5) for lithium storage 2. Along with non-toxicity, natural abundance and appropriate window potential for lithium insertion (< 0.4 V vs Li/Li+), silicon has been proposed as one of the most promising anode materials for next-generation LIBs 3. However, the drastic volume change (> 300%) and poor electronic conductivity of silicon often result in the pulverization and the

electrical

contact

loss

of

the

active

electrode

upon

repeated

Li+

alloying/dealloying process, resulting in sharp capacity decay and undesirable cycling


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stability. To overcome these obstacles, tremendous efforts have been made to improve the reversible performance of the Si-based anodes. One main strategy is to build efficient Si nanostructures (nanowires, nanotubes and hollow nanospheres, etc.) and 3D porous Si to reduce the critical fracture size of the silicon particles by mitigating the physical strain as well as shorten the distance of the charge transport/ion diffusion 4-8

. Also, the compositing of nano-Si with other conducting/protective materials can

improve the electronic conductivity, further buffer the volume expansion and suppress the formation of unstable solid electrolyte interface (SEI) of the active electrode during cycles

9-12

. Conventional synthetic methods to prepare well-defined nano-Si

anodes including chemical vapor deposition of silane precursors and the laser ablation of bulk Si are usually high energy consuming and relatively delicate, which largely restrict their scalable manufacture and practical application

13-15

. So far, developing

facile, low-cost and scalable way to prepare efficient nano-Si composites is still highly demanded. Interestingly, many plant materials can transform into high-level intricate and available nanostructures (e.g. carbon nanotubes, nanosheets and nanodots) after suitable treatments 16-19. It is quite promising to exploit these inexhaustible biomass or biomaterials toward high value-added functional materials. In particular, many plants contain abundant silica since they can introduce soluble silica or silicic acids (Si(OH)4, Si(OH)O-) from the soil by their roots

20, 21

. The sustainable derivation and

cost-effective recycle of Si materials from natural plant resources are extremely appealing. For instance, several nano-Si based anode materials with hierarchical



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porous structures have been produced from rice husk and reed leaves, which exhibited attractive capacity performance

22-25

. Bamboo is a “natural Si reservoir”, which can

absorb water-soluble silicic acids in the earth and then convert them into hydrated silica (SiO2⋅nH2O) particles, hence the content of silica in bamboo is impressive

26

.

Bamboo charcoal (BC), the carbonated product of natural bamboo under high temperature (>800°C), is inexpensive, electrically conductive and rich in silicon. In virtue of the highly porous network structure and large specific area, BC has been widely used as effective adsorbents or supports in various fields such as pollutant removals, moisture absorption, anti-bacterial and soil nutrients

27, 28

. The rational

utilization of the renewable BC as a silicon resource for advanced functional materials is thus intriguing. Herein, we report the continuous extraction and carbon coating of 3D porous silicon from BC via a facile and controllable process for the first time. When employed as the anode active materials for LIBs, the initial porous silicon originated from BC displayed a high lithium storage capacity of over 2810 mAh g-1 at 200 mA g-1, but quickly decayed in subsequent cycles. When further coated by a nitrogen-doped

amorphous

carbon

layer

through

polyacrylonitrile

(PAN)

carbonization, the final porous silicon/amorphous carbon (Si@N/C) composite exhibits significantly enhanced capacity and cycling performance. It can still maintain a reversible capacity of 603 mAh g-1 after 120 cycles at 200 mA g-1. The excellent electrochemical performance of the BC-derived porous Si@N/C composite can be attributed to the following two aspects: i) The N-doped amorphous carbon layer can


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improve the electrical conductivity of the composite and be elastic/protective to maintain the structural stability of the Si@N/C composite during repeated electrochemical cycles; ii) The 3D porous and hierarchical structured silicon extracted from bamboo charcoal can effectively accommodate the volume change and benefit the penetration/diffusion of the electrolyte ions;

Experimental Preparation of the 3D porous silicon: Firstly, 10 g of commercial BC were added into 1000 ml of 1 M HCl solution and then the mixture was allowed to proceed for 4 h at 95 °C under vigorous stirring. After cooling down to room temperature, the cleaned BC was filtered and washed with deionized water and ethanol for three times, respectively, and dried at 60 °C for 12 h. Afterwards, the cleaned BC was mechanically milled in ethanol for 12 h at 300 r min-1 in a high-energy planetary ball mill (the ball-to-material mass ratio was 16:1). After milling and drying, uniform BC powder was obtained. Subsequently, the BC powder was calcinated in air at 700 °C for 2 h in a muffle furnace to obtain the high purity SiO2 precursor. Then, the as-received SiO2 precursor was mixed with magnesium powder by an agate mortar, according to the ratio of nSiO2 : nMg = 1:2.5. The mixed powder was devolved into a closed reaction container under Ar atmosphere. Later, the closed reaction container was put into a tube furnace and calcinated at 650 °C for 2 h with a heating rate of 3 °C min-1 under the condition of Ar atmosphere. After cooling down to room temperature, the reacted product was slowly added into 1000 ml of mixed solution containing 1 M HCl and 2 ml of hydrofluoric acid (40wt%), and then was mechanically stirred for 2 h



at room temperature. After being filtered and washed with deionized water, and dried at 60 °C for 12 h, the porous silicon was finally collected and obtained. Preparation of the porous Si@N/C composites: The Si@N/C composites were obtained via a simple and consecutive procedure following above process. Firstly, suitable amount of polyacrylonitrile (PAN) was dissolved in 15 ml of N, N-dimethyl formamide (DMF) with magnetic stirring at 80 °C for 2 h. After that, porous silicon was dispersed into the PAN solution by ultrasonic for 2 h, and then was magnetically stirred for another 6 h. The added total mass of PAN and Si was 1g. After drying in the agate mortar under irradiation by an infrared lamp, the precursor mixture was transferred into a quartz tube furnace and calcinated at 700 °C for 2 h with a heating rate of 3 °C min-1 under the condition of mixed Ar/H2 atmosphere (vAr/vH2, 95:5). It is noteworthy that the weight ratio of the porous silicon adding into the PAN precursors was adjusted by 10%, 15%, 20%, respectively, and the as-prepared Si@N/C sample was marked with Si@N/C-10, Si@N/C-15 and Si@N/C-20, correspondingly. Characterization of materials: The morphology and microstructure of the obtained samples were imaged by scanning electron microscopy (SEM, JEOL, JSM-6380LA, Hitachi S-4800) and transmission electron microscopy (TEM, JEOL JEM-2100HR). The crystal phases of the samples were characterized by X-ray diffraction (XRD, Rigaku D/Max-Ш diffractometer) with Cu Kα radiation (λ = 1.5418 Å). The laser Raman spectroscopy was recorded by a Renishaw inVia reflex Raman microscope using the Ar-ion laser with 523 nm line as the excitation source. Further N2 adsorption/desorption measurements (Geminimore-2390, Micromeritics) of the


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samples were evaluated with the Brunauer-Emmett-Teller (BET) method. Moreover, thermogravimetric analysis (TGA) was acquired on a Netzsch TG 209 Thermo Analyzer from room temperature to 800 °C at a heating rate of 10 °C min-1 under air condition. Electrochemical measurement: In the tested coil-type cells, the working electrodes was composed of the as-prepared active materials (the porous Si, Si@N/C-10, Si@N/C-15 and Si@N/C-20, respectively), carbon black (electronic conductive additive) and polyvinylidene fluoride (binder) at the weight ratio of 7:2:1. The average weight of the electrodes was ~1.5 mg, and all the specific capacity in this study is calculated based on the total weight of the as-prepared active materials in the electrode. For example, the specific capacity of the Si@N/C-based electrode is calculated based on the total mass of the Si@N/C composite. Meanwhile, using lithium metal as the counter electrode, 1 M LiPF6 dissolved in a mixed solution of ethylene carbonate/dimethyl carbonate (EC: DMC) (1:1, v/v) as the electrolyte and a piece of micro-porous polypropylene film (Celgard 2400) as the separator. The cells were assembled in a glove box filled with ultra-pure argon. Galvanostatic charge/discharge tests were performed on a Neware battery test system with the cut-off voltage ranging from 0.01 to 1.2 V (vs. Li/Li+) at room temperature. Electrochemical impedance spectroscopy (EIS) tests were performed on an IM6e electrochemical workstation across the frequency range between 1 MHz and 10 mHz with an AC signal of 5 mV. The cyclic voltammograms (CV) experiments were carried out on Potentiostat/Gallanostat Model (Perkin-Elmer 273A, EG & E) with



voltages between 0.01 and 1.2 V (vs. Li/Li+) and a scan rate of 0.1 mV s-1.

Results and Discussion

Scheme 1 Flow diagram for the synthetic procedures of the Si@N/C composite derived from bamboo charcoal (BC). Scheme 1 illustrates the synthetic process of the Si@N/C composite using bamboo charcoal as the silicon source. First, impurities were removed from the countless pores in BC by hydrochloric acid pickling, and the cleaned BC was grounded into uniform and fine black powder with ball milling. Then the BC powder was calcined under air to remove the major organic matters and other carbon components to produce the white SiO2 precursor. The SiO2 precursor then experienced magnesiothermic reduction and was further washed by a mixed HCl/HF solution to remove the MgO by-product and the residual SiO2, and high purity porous silicon in brown color was received 29. Finally, the high temperature decomposition and carbonization of the PAN polymer precursor generated the amorphous carbon-coated Si (Si@N/C) composite. Fig. 1 shows the SEM images of BC, the SiO2 precursor, the porous Si and the


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Si@N/C-15 composite. According to Fig. 1a, BC owns a unique 3D porous interconnected micro-tube structure consisting of highly ordered macropores with an average diameter of 20~30 µm, which were mainly originated from the porous nature of the bamboo plant. After calcination, the interconnected micro-tube structure of BC collapsed and the resulted SiO2 precursor became to many microflakes with a dominant sheet-like morphology, as shown in Fig. 1b. Noting that a small amount of SiO2 turned into the chip-like powders, which are possibly due to the silica component existing in the different parts of the initial BC material. Based on the SEM images presented in Fig. 1c and 1d, sheet-like SiO2 microflakes were transformed into purified Si possessing a 3D porous hierarchical structure after magnesiothermic reduction. The 3D porous Si is composed of vertically aligned nanobunches with worm-like network on the top part and thin nanowhisker bundles having a length of several hundred nanometers tightly connected at the bottom part. These nanowhiskers are basically made up from ultrafine Si nanoparticles. The 3D hierarchical morphology of Si was likely caused by the extensive magnesiothermic reaction across the sheet-like SiO2 precursor along with the subsequent gas release during the HCl/HF treatment 30. As demonstrated in Fig. 1e and Fig. 1f, the final Si@N/C composite is covered with relatively smooth carbon coating on the surface and gives a nanonest network following the main skeleton of the 3D porous Si. It is expected that the hierarchical porous structure of the 3D Si@N/C composite can offer sufficient voids to buffer the mechanical stress induced by the volume expansion of the active material. Shown in Fig.1g, the EDX spectrum and the corresponding element



mapping images reveal that3D Si can be evenly coated by the N-doped carbon layer. While the atomic content of Si, C, N and O in the Si@N/C composite is estimated to be 28.0%, 60.2%, 5.0% and

Figure 1 SEM images of (a) BC; (b) SiO2 precursor; (c), (d) 3D porous Si; (e), (f) Si@N/C-15; (g) EDX spectrum of the Si@N/C-15 composite and the elemental mapping image of Si, C, N and O, respectively.


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6.8%, respectively. It suggests that a moderate amount of nitrogen has been readily doped into the composite and the entire Si@N/C composite shows good homogeneity. Fig.2 displayed the TEM images of the as-prepared Si@N/C composite. From Fig. 2a and 2b, the interconnected porous structure and the amorphous carbon coating at the outside of the silicon component can be distinguished, which would be favorable for improving the structural integrity of the Si@N/C composite. HR-TEM in Fig. 2c further identified the amorphous carbon layer locating around Si, and the lattice spacing of 0.31 nm can be indexed to the characteristic (111) crystal plane of silicon with good crystallinity 25. The N-doped carbon coating layer is necessary to wrap Si in the hybrid but also can modify the poor electronic conductivity of the Si matrix 23.

Figure 2 (a), (b) TEM image of Si@N/C-15; (c) HR-TEM image of Si@N/C-15. XRD patterns of the products involved in the whole synthetic process of the Si@N/C composite are shown in Fig. 3a. It can be seen that the main peak of the SiO2 precursor appeared at 2θ = 26.8° is relatively sharp and strong, indicating that the produced SiO2 has high crystallinity and good crystal structure (JCPDS 46-1045). The diffraction peaks of the porous Si extracted from magnesium reduction are prominent and centered at 2θ = 28.6°, 47.3°, 56.1°, 69.1° and 76.4° that can be assigned to the (111), (220), (311), (400) and (331) plane of the silicon crystallites, respectively



(JCPDS 27-1402)

23, 31

. No significant impurity peaks are found, indicating the high

purity of the received porous Si material. Moreover, the XRD patterns of the Si@N/C composites exhibit that the intensity of the characteristic peak centered at 2θ = 28.6° is successively weakened

Figure 3 (a) XRD pattern of SiO2, Si and the Si@N/C composites (Si@N/C-10, Si@N/C-15, Si@N/C-20), respectively; (b) Raman spectrum of Si and Si@N/C-15; (c) TG curves of the as-prepared Si@N/C composites. Inset: TG curves of BC; (d) N2 adsorption/desorption isotherm curves of SiO2, Si and Si@N/C-15. Inset: pore-size distributions of the corresponding samples. with the decrease of the silicon content in the composite, also due to the coverage of the carbon layer on the surface. At the same time, a broad peak appears at 2θ = 25.8°


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becomes stronger that is in accordance with the increased content of carbon in the material 32. Moreover, the Raman spectra of the porous Si and the Si@N/C composite are displayed in Fig. 3b. The characteristic peak of the porous Si around 520 cm-1 and a small peak below 1000 cm-1 belong to the typical vibration modes of the crystalline Si 26, 33

. The two peaks centered at 1357 cm-1 and 1597 cm-1 corresponds to the D band

and the G band of the carbon species, respectively, which verified an intrinsic electronic conductivity of the delocalized sp2 π bonding 34. Compared with the porous Si, the reduced peak intensity at 520 cm-1 further confirms the amorphous carbon layer coating on the surface of the porous silicon. An intensity ratio (R value: ID/IG) is calculated to be 1.21 which implies a substantial amount of conductive carbon species involved in to contribute good electrical conductivity of the Si@N/C composite. In order to determine the exact content of the silicon component in relevant materials, thermogravimetric analysis of BC and the three Si@N/C composites (Si@N/C-10, Si@N/C-15 and Si@N/C-20) were carried out under air atmosphere from room temperature to 800 °C. In Fig. 3c, the dominant weight loss of 85.2% in the BC material occurred between 350 ~ 550 °C is attributed to the combustion of the organic matters and the carbon species in air. The content of the residual substance (silica) is about 14.8%, which corresponds to a silicon content of ca. 6.9% in BC. As for the Si@N/C composites, the mild weight loss below 100 °C is caused by the evaporation of the adsorbed water and small organic molecules. All the composites endured a major weight loss from 500 °C to 750 °C indicating the combustion of the



N-doped amorphous carbon and proving good thermal stability of the Si@N/C composites

35

. Since the SiO2 precursor from BC has been converted into Si after

magnesiothermic reduction, the residual weight of the final Si@N/C composite after calcination under air can be almost ascribed to silicon. Therefore, the silicon content of the Si@N/C-10, Si@N/C-15 and Si@N/C-20 composite is about 10.7%, 19.5% and 26.1%, respectively. The N2 adsorption/desorption isotherms of the samples are demonstrated in Fig. 3d. It is significant that the adsorption/desorption curves reveal the obvious porous character of the samples possessing dominant mesopores. The specific surface area of the SiO2 precursor, the porous Si and the Si@N/C-15 composite is derived to be 58.56, 79.51 and 111.23 m2 g-1, respectively. The introduction of N-doped carbon layer can apparently increase the specific area of the porous Si material. Meanwhile, the corresponding pore size distributions of the samples (inset of Fig. 3d) were analyzed by the Barret-Joyner-Halenda (BJH) method. It shows that the pores of the SiO2 precursor are mainly centered at around 3.7 nm, while the pore size distribution scope of the porous Si and the Si@N/C composite are distinctly expanded. Especially, the Si@N/C-15 composite has a majority of pores distributed in the mesoporous region with an average pore size of around 11.2 nm. These porous characters agree well with aforementioned 3D hierarchical porous morphology of the BC-derived Si and the Si@N/C composite, which will likely to facilitate the diffusion of electrolyte ions and alleviate the volume changes of the active silicon particles during lithium storage. XPS characterization was conducted to further investigate the surface states and the


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elemental composition of the Si@N/C composite. As shown in Fig. 4a, the survey spectrum verifies the existence of Si, C, O and N elements in the composite. Moreover, the O element belongs to the re-oxidized silicon species probably formed during the high temperature carbonization process, while the N element was introduced by the PAN carbon source. In Fig. 4b, the high-resolution XPS spectrum of Si 2p displays two

Figure 4 (a) The survey XPS spectrum of Si@N/C-15; (b) Si 2p spectrum; (c) C 1s spectrum; (d) N 1s spectrum. main peaks centered at 103.1 and 99.2 eV, corresponding to the Si-O bond (Si4+) and the Si-Si bond (Si0), respectively 36. The C 1s spectrum of Si@N/C-15 is shown in Fig. 4c. The peak centered at 284.7, 285.5, 286.5 and 288.2 eV can be attributed to the



C-C, C=C, C-N and C=N bond, respectively, suggesting strong interaction between the amorphous carbon and the N dopant

37

. Furthermore, the N 1s spectrum is

presented in Fig. 4d. The peak centered at 398.4 and 400.0 eV is ascribed to the pyridinic group (C-N=C) and the graphite-like substitutional structure (N coordinated with three C atoms), respectively

38

. These XPS results further proves that nitrogen

species is well incorporated into the amorphous carbon coating and the Si@N/C composite. In addition, the electrochemical performances of the active materials were evaluated by coin cells. Fig. 5 demonstrates the selected charge/discharge profiles (the 1st, 2nd, 10th and 50th cycle) of the porous Si, the Si@N/C-10, Si@N/C-15 and the Si@N/C-20 composite at a current density of 200 mA g-1, respectively. In Fig. 5a, the initial discharge capacity of the porous silicon reaches 2810 mAh g-1 and the initial coulombic efficiency is 80.0%. The first discharge curve of the porous silicon has a stable and wide span of discharge plateau at 0.1 V which corresponds to the alloying reactions of silicon with lithium

39

. As for the subsequent cycles, the reversible

capacity of the porous Si decayed seriously. In contrast, the charge/discharge curves of the Si@N/C-10, Si@N/C-15 and Si@N/C-20 composites for different cycles are displayed in Fig. 5b~d. The initial discharge capacity of the Si@N/C-10, Si@N/C-15 and Si@N/C-20 composite is 1340, 1697 and 1534 mAh g-1, corresponding to an initial coulombic efficiency of 59.6%, 67.4% and 68.1%, respectively. As the Si content increased, Si@N/C-15 shows the highest discharge capacity since a moderate amount of the amorphous carbon component is usually beneficial for achieving


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optimized lithium storage performance

40

. Compared with the porous silicon, the

lower initial coulombic efficiencies of the Si@N/C composites may be caused by the formation of the larger SEI film during the initial charge/discharge process 41, which is coincide with the higher specific surface area of the Si@N/C hybrid than that of Si (see Fig. 3d). Except the significant initial discharge plateau at 0.1 V, a gentle voltage slope between 0.1 to 0.2 V appears simultaneously owing to the positive effect brought by the N-doped amorphous carbon coating in the Si@N/C composite. As compared to the Si counterpart, the change of the



Figure 5 Discharge/charge profiles of the samples for the 1st, 2nd, 10th and 50th cycle, respectively at the current density of 200 mA g-1 in the voltage window of 0.01 ~ 1.2 V. (a) Porous Si; (b) Si@N/C-10; (c) Si@N/C-15; (d) Si@N/C-20. (e) Cycling performances of the porous Si and the Si@N/C composites at current density of 200 mA g-1; (f) Rate performance of the porous Si and the Si@N/C composites under different current densities (I: 100, II: 200, III: 400, IV: 800, V: 1600 mA g-1). charge/discharge profiles for the subsequent cycles of the three composites greatly


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decreases. Especially the discharge curves of all the Si@N/C composite overlaps well, suggesting good reversibility and cycling tolerance of the Si@N/C composites. It further confirms that the combination of the porous Si with N-doped carbon coating is reasonable and efficient for reversible lithium storage since nitrogen-doped carbon materials have been widely used to improve the electronic conductivity, the catalytic activity and even the long term stability of the active electrodes in various lithium batteries 42-44. In addition, the cycling performance of the porous Si and the Si@N/C composites at 200 mA g-1 are shown in Fig. 5e. It is obvious that the Si@N/C composites show much more enhanced capacity retention than that of the porous Si upon cycling. After 120 cycles, the reversible capacity of the porous Si, Si@N/C-10, Si@N/C-15 and the Si@N/C-20 composite is 155, 454, 603 and 430 mAh g-1, respectively. Among them, the porous silicon exhibits high capacities for the first few cycles but rapidly decays continuously, which is originated from the serious volume expansion of Si during repeated cycles 1. Comparatively, though the theoretical capacity of the Si@N/C composites decrease with the increased content of the carbon component, the cycling durability of the three Si@N/C composites is much better and impressive. In this regard, the N-doped amorphous carbon layer in the Si@N/C composites can protect the active Si particles from extensive exposure to the electrolyte and buffer the volume expansion of the active material during the charge/discharge process, resulting in significantly improved cycling stability. Particularly, the Si@N/C-15 composite achieves the best cycling stability with 0.2% drop of the reversible capacity per cycle.



Besides, the coulombic efficiency of the Si@N/C-15 composite rapidly rises to over 99% after the first charge/discharge cycle, and approaches to about 100% for the later cycles. It infers that relatively stable SEI film has been established during the initial cycle and the Si@N/C-15 composite owns excellent cycling ability. Fig. 5f shows the rate performances of the active materials at various current densities, i.e., from 100 mA g-1 to 1600 mA g-1 and then restores to 100 mA g-1. The porous Si shows poor rate capabilities and thus inferior lithium storage performance under high current densities. Even there is nearly no electrochemical activity remained when the current density increases to 1600 mA g-1, the specific capacity of Si still decays seriously when the current density returns to 100 mA g-1. In accordance with above-mentioned prominent cycling stability, the Si@N/C-15 composite exhibits the best rate performance, delivering a discharge specific capacity of 788, 730, 628, 491 and 360 mAh g-1 at the current density of 100, 200, 400, 800 and 1600 mA g-1, respectively. Besides, the Si@N/C-15 composite realizes good recovery capacity and retains the highest discharge capacity when the current density returns back to 100 mA g-1. Both of the reversible capacities and the cycling performance of our BC-derived Si@N/C composite are comparable and even superior to many reported nano-Si/carbon composites

14, 45-48

. The optimal electrochemical performance of the

Si@N/C composite can attribute to its 3D hierarchical and porous structure affording favorable advantages. On one hand, the 3D porous feature can not only accelerate the electrolyte penetration and the diffusion of lithium ions, but also efficiently provide enough void space to accommodate the volume change of the BC-derived Si@N/C


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composite. On

Figure 6 Cyclic voltammograms of the porous Si (a) and the Si@N/C-15 composite (b) at a scanning rate of 0.1 mV s-1 for a few cycles; (c) EIS curves of the porous Si



and the Si@N/C composites after 100 cycles. Inset: the equivalent circuit to fit the EIS curves. the other hand, the N-doped amorphous carbon coating can further promote the electronic conductivity and the mechanical integrity of the composite, leading to enhanced electrochemical activity and stability of the Si@N/C composite toward continuous lithiation/delithiation cycles 45, 49. To gain better understanding of the reactions of the electrode materials during Li+ intercalation/deintercalation process, cyclic voltammetry (CV) of the porous Si and the Si@N/C-15 composite was measured at 0.1 mV s-1 between 0.01 and 1.2 V, which is shown in Fig. 6a and b, respectively. For the first cathodic cycle (Fig. 6a), the broad peak around 0.6 to 1.0 V is attributed to the formation of SEI film and the decomposition of the organic electrolyte

25

. This peak almost disappears for the

following cycles. Until the 3rd scan cycle, the distinct reduction peak centered at 0.18 V is caused by the alloying reaction of the crystalline silicon transforming to the amorphous LixSi 50. While the two significant oxidation peaks at 0.36 V and 0.54 V belonging to the dealloying process of LixSi 51. Nevertheless, the area change of the curves for different scan cycles is significant, which suggests the poor repeatability and stability of the Si electrode. It concludes that the porous Si undergoes inevitable formation of unstable SEI film and serious structural damage during the initial discharge process. Specially, the Si@N/C composite demonstrates different behavior during the first cycle, as shown in Fig. 6b. In the initial cathodic cycle, the broad protrusion (from 0.6 to 1.1 V) is not sharp, which implies that the Si@N/C composite


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has relatively mild lithiation kinetic of the SEI film due to the protection of the amorphous carbon coating on the surface

46

. However, the strong reduction peak

centered at 0.19 V and the synergetic peak near 0 V can be clarified due to the insertion of lithium ions into the crystalline Si, showing better activity than Si because of more stable SEI film for the initial cycle

37, 51

. At the same time, the remarkable

oxidation peak centered at 0.55 V is mainly attributed to the extraction of lithium (delithiation) from LixSi to generate pure silicon. As for the subsequent cycles, the main redox peaks of the Si@N/C composite are similar to those of Si. It is noteworthy that the CV curves almost coincide with each other and overlap well, indicating that the Si@N/C composite has excellent electrochemical repeatability and reversibility toward lithium storage. In order to further comprehend the electrochemical behavior of the active electrodes, electrochemical impedance spectroscopy (EIS) characterization were carried out. The EIS curves of the Si electrode and the Si@N/C composite electrodes after 100 charge/discharge cycles at current density of 200 mA g-1 are shown in Fig. 6c. The typical EIS curve consist of a semicircle in the high frequency region that is indexed to the charge transfer resistance (Rct) at the electrode/electrolyte interface and a straight line in the low frequency region belonging to the diffusion of the electrolyte ions into the active electrode (W, Warburg impedance)

23, 26, 35

. The Rct of the porous

Si, Si@N/C-10, Si@N/C-15 and Si@N/C-20 composite is estimated to be 248, 67, 83 and 110 Ω, respectively. It is apparent that the Si@N/C composites own much smaller Rct than Si, implying the faster charge transfer across the electrode/electrolyte



interface because of the higher electronic conductivity. It is thus verified that the N-doped amorphous carbon can effectively improve the electrical conductivity and the charge transfer kinetic of the Si-based composites. Impressively, the Si@N/C composite can still maintain good diffusion properties even after 100 cycles, indicating the significant electrochemical activity and stability of the Si@N/C composite

52

. These results are in consistent with the high reversible capacity

performance, high rate capabilities and excellent cycling stability of the Si@N/C composite.

Conclusions In summary, 3D porous silicon has been extracted from BC via simple calcination and magnesiothermic reduction. Combined with continuous PAN carbonization, N-doped amorphous carbon layer was further coated onto the porous Si and produced the 3D porous silicon/carbon-based composites. The whole synthetic process of the Si@N/C composite is simple, economic and controllable, which does not involve any expensive reagents and toxic silicon precursors. The obtained Si@N/C composite possessed a 3D hierarchical structure giving desirable micro/nano porous morphology, high specific surface area and mechanical integrity. Thanks to the porous structural benefits and the protective amorphous carbon layer on the surface, the Si@N/C composite realized improved reversible specific capacity, excellent cycling stability and high rate capabilities. It is reasonable that the N-doped amorphous carbon layer can effectively accommodate the volume change of the 3D porous silicon, as well as facilitate the diffusion of lithium ions in the Si@N/C composite. This


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biomass-initiated strategy is accessible for the design and preparation of many other functional hierarchical nanocomposites. Also, our BC-derived silicon/carbon composite is promising as cost-efficient anode materials for next-generation LIBs.

Acknowledgments This research was financial supported by the Guangzhou Science and Technology Planning Project (No. 201704030022), the Natural Science Foundation of Guangdong Province (No.2017A030313283 and 2017A030313083), the Guangdong Science and Technology Planning Project (No. 2015A020209147), and the National Natural Science Foundation of China (No. 51602109 and 21673083).

Conflicts of interest There are no conflicts to declare.

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Table of Content: Bamboo charcoal-derived 3D porous Si@N/C composite is obtained via an economic and scalable process. 205x108mm (144 x 144 DPI)



Scheme 1 Flow diagram for the synthetic procedures of the Si@N/C composite derived from bamboo charcoal (BC). 186x99mm (144 x 144 DPI)


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Figure 1 SEM images of (a) BC; (b) SiO2 precursor; (c), (d) 3D porous Si; (e), (f) Si@N/C-15; (g) EDX spectrum of the Si@N/C-15 composite and the elemental mapping image of Si, C, N and O, respectively. 254x313mm (144 x 144 DPI)



Figure 2 (a), (b) TEM image of Si@N/C-15; (c) HR-TEM image of Si@N/C-15. 348x116mm (144 x 144 DPI)


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Figure 3 (a) XRD pattern of SiO2, Si and the Si@N/C composites (Si@N/C-10, Si@N/C-15, Si@N/C-20), respectively; (b) Raman spectrum of Si and Si@N/C-15; (c) TG curves of the as-prepared Si@N/C composites. Inset: TG curves of BC; (d) N2 adsorption/desorption isotherm curves of SiO2, Si and Si@N/C15. Inset: pore-size distributions of the corresponding samples. 285x225mm (144 x 144 DPI)



Figure 4 (a) The survey XPS spectrum of Si@N/C-15; (b) Si 2p spectrum; (c) C 1s spectrum; (d) N 1s spectrum. 217x174mm (144 x 144 DPI)


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Figure 5 Discharge/charge profiles of the samples for the 1st, 2nd, 10th and 50th cycle, respectively at the current density of 200 mA g-1 in the voltage window of 0.01 ~ 1.2 V. (a) Porous Si; (b) Si@N/C-10; (c) Si@N/C-15; (d) Si@N/C-20. (e) Cycling performances of the porous Si and the Si@N/C composites at current density of 200 mA g-1; (f) Rate performance of the porous Si and the Si@N/C composites under different current densities (I: 100, II: 200, III: 400, IV: 800, V: 1600 mA g-1). 298x328mm (144 x 144 DPI)



Figure 6 Cyclic voltammograms of the porous Si (a) and the Si@N/C-15 composite (b) at a scanning rate of 0.1 mV s-1 for a few cycles; (c) EIS curves of the porous Si and the Si@N/C composites after 100 cycles. Inset: the equivalent circuit to fit the EIS curves. 144x326mm (144 x 144 DPI)


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N-doped Carbon Composite Derived from Bamboo

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