Multi-Scale Buffering Engineering in Silicon-Carbon Anode for

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Multi-Scale Buffering Engineering in SiliconCarbon Anode for Ultrastable Li-Ion Storage Guolin Hou, Benli Cheng, Yijun Yang, Yu Du, Yihui Zhang, Baoqiang Li, Jiaping He, Yunzhan Zhou, Ding Yi, Nana Zhao, Yoshio Bando, Dmitri V. Golberg, Jiannian Yao, Xi Wang, and Fangli Yuan ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.9b03355 • Publication Date (Web): 19 Aug 2019 Downloaded from pubs.acs.org on August 19, 2019

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Multi-Scale Buffering Engineering in Silicon-Carbon Anode for Ultrastable Li-Ion Storage Guolin Hou,† Benli Cheng,† Yijun Yang,‡ Yu Du,†,§ Yihui Zhang,‡ Baoqiang Li,† Jiaping He,† Yunzhan Zhou,‡,‖ Ding Yi,‡ Nana Zhao,‡ Yoshio Bando,⊥ Dmitri Golberg,# Jiannian Yao,‖,⊥ Xi Wang*, ‡ and Fangli Yuan*,†,※

†

State Key Laboratory of Multi-Phase Complex Systems, Institute of Process

Engineering, Chinese Academy of Sciences (CAS), Zhongguancun Beiertiao 1 Hao, Beijing 100190, P. R. China. ‡

Key Laboratory of Luminescence and Optical Information, Ministry of Education,

School of Science, Beijing Jiaotong University, Beijing, 100044, P. R. China. § University

of Chinese Academy of Sciences (UCAS), No.19A Yuquan Road, Beijing

100049, P. R. China. ‖ Chemistry ⊥

and Chemical Engineering Guangdong Laboratory, Shantou 515031, China.

Institute of Molecular Plus, Tianjin Key Laboratory of Molecular Optoelectronic

Sciences, Department of Chemistry, Tianjin University, and Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), Tianjin 300072, P. R. China. #

Science and Engineering Faculty, Queensland University of Technology (QUT), 2

George St., Brisbane, QLD 4000, Australia. ※ Center

of Materials Science and Optoelectronics Engineering, University of Chinese

Academy of Sciences, Beijing 100049, P. R. China. 

Correspondence

should

be

addressed

ACS Paragon Plus Environment

to:

[email protected]

or

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[email protected]


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ABSTRACT Silicon-carbon (Si-C) hybrids have been proven to be the most promising anodes for the next-generation lithium ion batteries (LIBs) due to their superior theoretical capacity (~4200 mAhg-1). However, it is still a critical challenge to apply this material for the commercial LIB anodes, because of the large volume expansion of Si, unstable solid-state

interphase

(SEI)

layers,

and

huge

internal

stresses

upon

lithiation/delithiation. Here, we propose an engineering concept of multi-scale buffering, taking the advantage of a nano sized Si-C nanowire architecture through fabricating specific micro-sized wool-ball frameworks to solve all the abovementioned problems. The regarded wool-ball-like frameworks, prepared at high yields, nearly matching industrial scales (they can be routinely produced at a rate of ~300 g/h), are composed of Si/C nanowires (NWs) building blocks. As anodes, the Si-C wool-ball frameworks show ultra-stable Li+ storage (2000 mAhg-1 for 1000 cycles), high initial coulombic efficiency (ICE) of ~90% and volumetric capacity of 1338 mA h cm-3. In situ TEM proves that the multi-scale buffering design enables a small volume variation, only ~19.5%, reduces the inner stresses and creates very thin SEI. The perfect multi-scale elastic buffering makes the regarded material more stable compared to common Si nanoparticles-assembled counterpart electrodes.

KEYWORDS: multi-scale buffering engineering, silicon anode, in situ TEM, ultra-stable Li-ion storage, Li-ion batteries


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Lithium-ion batteries (LIBs) with increased energy densities are necessary for batteries applied in portable electronic devices, electric vehicles and energy storage stations.1 As compared with commercial graphite anodes,2 silicon (Si)-carbon material has been highlighted as a promising electrode for the next generation of LIBs because of its relatively low working potential,3-5 abundance in nature, and the highest reported theoretical specific capacity of ~3572 mA h g-1 (ten times of that of graphite).6 For example, a wide range of advanced Si-C structured anodes have been designed, such as Si-C yolk-shell and graphene-encapsulated Si particles,7-11 or pomegranate structures,12-16 demonstrating a significant improvement in capacity and cycle life. However, it is still a long journey to match the application standards using the current Si-C anode design,17-20 for instance, excessive volumetric expansion (e.g. 200 ~ 400%), uncontrollable SEI layers, enormous internal stresses, low tap density, initially low coulombic efficiency (ICE) and volumetric capacity limit the potentials for mass usage. Decent electrochemical performances of a Si-C anode are necessary for its commercial application as a LIB anode material, which should be satisfied by a high initial coulombic efficiency (ICE) (>85%), an enhanced volumetric capacity (e.g. higher than that of graphite, 550 mA h cm-3) and a superior rate capability. Therefore, a systematic work should be done to tackle the above-mentioned problems in order to make an ideal Si-C anode. Herein, inspired by the superb elasticity of a wool-ball, we demonstrate an interesting multi-scale buffering engineering principle for assembling a battery anode via fabricating a S/C NWs-constructed wool-ball-like framework (Scheme 1). Such a


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design has multiple advantages: (1) nano-sized Si/C nanowires have the ability to accommodate large strains without pulverization and capacity fading towards radial direction,21-23 while nanoscale carbon shell completely encapsulates the Si NWs and works as a suitable elastic buffer that accommodates the radial expansion, and ensures the formation of a stable and thin SEI layer;24 (2) a micro-sized wool-ball secondary particle structure with internal voids provides the space to further alleviate the inherent volume changes and inner stresses, and endows the anode with perfect elastic characteristics to maintain the integrity and continuity of the electrode framework and cell configuration (superior ICE, tap density and volumetric capacity); (3) Importantly, kilogram-scale manufacturing of Si/C nanowires-assembled balls (300 g/h) is realized using a radio-frequency induction thermal plasma (RF-plasma) system in-tandem with a spray drying procedure (RFP-SD). It worth noting that, to date, this has not been possible using any other route.18 Being tested as an anode for LIBs, the wool-ball-like spherical frameworks exhibited ultra-stable Li+ storage properties: superior cycling stability for 1000 cycles, ~90% ICE, and high volumetric capacity of 1338 mA h cm-3. Their good structural stability, thin SEI, and excellent elastic properties during cycling have been additionally investigated through in-situ HRTEM probing and theoretical calculations. This suggested that the developed approach is a promising way to fabricate practical Si-C anodes.


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Scheme 1. Illustration of a multi-scale buffering engineering design.

RESULTS AND DISCUSSION Synthesis and Characterizations of Si-C Wool Balls (Si-C WBs). The detailed synthesis setup and procedure for making Si-C wool balls are illustrated in Figure 1a and Figure S1 (Supporting Information (SI)), in which a RF-plasma system combined with the spray drying procedure (RFP-SD) have been used. It is worth to note that quasiindustrial-scale production of Si-C anode materials can be realized under this method, the production yield reaches ~300 g/h using micro-Si as a raw material, and therefore, we have successfully prepared 1.5 kilograms of the material in a single batch (Figure S1), as also seen in Figure 1e. This step is very important for the promotion of practical applications of Si-based anodes, as most Si nanomaterials reported to date have not yet been scalable. From the X-ray powder diffraction (XRD) patterns shown in Figure 1b, as-made products are identified as a crystalline silicon (JCPDS No. 00-027-1402) and carbon (a broad peak around 25°); it is noted that (111) and (220) peaks imply the oriented facets owing to their strongest peak intensity, while the obvious “bumpy” peak at ~25 suggests the existence of amorphous Si. The as-synthesized composite exhibits


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spheres with a diameter of 3 to 10 m (Figures 1c-d). The micro-size of the material generally induces high tap density and further guarantees the enhanced volumetric energy density. Si/C NWs are intertwined into wool-ball-like spherical frameworks (SiC WBs, Figures 1d, f-g). It could be also noted that there are many nano-sized pores within a Si/C NW framework, as revealed by high-magnification TEM images (Figure 1f), which are mainly caused by the spray drying and the pyrolysis of glucose. Brunauer-Emmett-Teller (BET) and Barett-Joyner-Halenda (BJH) measurements show that Si-C WBs possess a specific surface area of 80.105 m2 g-1 with a pore volume of 0.3 cm3 g-1 (Figure S2, SI). The well-defined internal void space within Si/C NWs allows Si to expand freely without disturbing the integrity of the secondary particles, which is really important to maintain the continuity of the whole electrode framework and overall cell configuration.25 Every Si/C NW has the central crystal Si (c-Si) core, the subshell amorphous Si (aSi) coating, and the outermost carbon shell, as shown in Figure 1g-h. The numerous Si/C NWs are intertwined together and assembled into a spherical framework with abundant well-defined internal voids, thus making the final wool-ball structure. Asprepared by RF-thermal plasma, Si nanowires have an average diameter of ~50 nm (Figure S3, SI), much less than the critical sizes of crystalline Si nanowires (300-400 nm) upon lithiation (with respect to fracture); and the axial orientation of these nanowires is along the orientation, which matches well the XRD results (Figure 1b). Note that the small NW diameter would allow for better accommodation of the large volume changes, whereas the direct 1D electronic pathways could facilitate an


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efficient charge transport.21,22 Electron diffractions (Figure S3d inset, SI) also substantiate the overall crystalline-amorphous core-shell morphology.

Figure 1. Synthesis and characterization of the Si-C WBs consisting of Si/C NWs. (a) Schematic of the fabrication process of Si-C WBs. (b) Typical XRD patterns of asprepared Si-C WBs. (c-d) SEM and TEM images of as-prepared Si-C WBs. (e) Photograph of a final product (1.5 kg). (f-h) TEM, HRTEM and element distributions of Si/C NWs. (i) The Si-C WBs thermogravimetry analysis (TGA) and differential scanning calorimetry (DSC) curves under air atmosphere at a 10 C min-1 heating rate.

The outermost carbon layer encapsulating each Si NW is 3~5 nm thick (Figure 1g). Raman spectroscopy (Figure S4, SI) reveals two typical characteristic peak of carbon


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materials at 1350 cm-1 (D band, which is assigned to the disordered carbon), and the peak at 1580 cm-1 (G band, which corresponds to the graphitic carbon).26 It can be clearly discerned that the ratio of ID/IG (= 0.797) is relatively low, suggesting that the graphitization degree of carbon for Si-C WBs is high and the Si-C WBs exhibit good electrical and elastic properties.27,28 Elemental mapping analysis (Figure S5, SI) suggests that carbon is uniformly coated onto Si NW surfaces. DSC-TGA measurements (Figure 1i) suggest that the carbon content in Si-C WBs is about 6.3 wt.%. Although the carbon layer in Si/C NWs is only of a few nanometer thick, it not only firmly connects Si NWs into an integral microsphere, which provides a 3D conductive framework and subsequently enhances the total electrical conductivity, but also serves as a suitable elastic buffer to accommodate the volume expansion of Si NWs, ensuring the formation of a stable SEI layer and the corresponding high ICE.12,29 Electrochemical Performances of Si-C WBs. Then Si-C WBs were electrochemically evaluated in detail. The voltage profiles of Si-C WBs (Figure 2a) exhibit typical electrochemical features of a Si anode: a broader discharge voltage plateau at about 0.05-0.1 V, which corresponds to the formation of LixSi alloy during the lithiation process. A charge voltage plateau, at about 0.35-0.50 V, can be observed, which should be attributed to the de-lithiation process.30 It is worthy to note that high ICE is much important to the practical application of the anode materials.5 This factor plays the key role because it accounts for the most of the Li-ion loss and electrolyte consumption during the formation of a SEI layer, which greatly affects the cycle life and rate performance.31 Remarkably, the Si-C WBs exhibit an ultrahigh ICE of ~90% which


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quickly increases to 95% during the second cycle (Figure 2a). High ICE value typically means the stability of SEI layer on the electrode surfaces and reversible Li+ lithiation/delithiation within the Si-C WBs.32,33 In fact, as demonstrated for the reported Si or Si-C composites, the carbon precursors (Table S1, SI) and binders (Table S2, SI) both have a significant effect on the ICE.34-36 However, to date, only a few reports have shown an ICE value of >85%.37 In order to make a fair comparison, we utilized the same method to fabricate the similar structure, 5-10 μm-sized Si nanoparticles-C balls (Si NPs-C balls) built from many 50-100 nanosized spheres through tuning the Si/C ratios (Figure S6, SI). Note that the BET and BJH tests also demonstrate that Si NPs-C balls obviously exhibit a smaller surface area and pore volume than those of Si-C WBs, and the ICE of Si NPs-C balls is ~75% (Table S2, Figure S7, SI). It suggests that the present multi-scale buffering engineering design in Si-C WBs may be the key to form a stable SEI layer during cycling, thus leading to the improvement of ICE. This can also be confirmed by the Nyquist curve from the electrochemical impedance spectroscopy (EIS) (Figure 2b), where the semicircle means the charge-transfer resistance.12 It can be discerned that the surface kinetics of Si-C WBs is much faster than that of bare Si NWs. In addition, this phenomenon exhibits little change even after 100 cycles, which further suggest that the SEI layer is stable during Si-C WBs cycling. Cyclic voltammetry (CV) tests (Figure 2c) provide a further evidence for the stable SEI layer formation and good reversibility during cycling. For the cathodic branch (lithiation process), two peaks at 0.01 and 0.16 V are discerned, which can be attributed to the formation of LixSi alloy, while the two peaks at 0.35 and 0.52 V of the anodic branch


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(delithiation process) could be assigned to the dealloying stage, which transforms the LixSi alloy back into amorphous Si.38 The CV curve is consistent with those reported previously. Due to the existence of the carbon shell at every individual Si NW, the peaks of Si-C WBs exhibit a left shift. 39 Rate performance is another important factor to the practical application of anode materials.40 As depicted in Figure 2d, the as-prepared Si-C WBs show an enhanced rate performance. The ultrahigh capacity of 2433 mA h g-1 is delivered at 0.2 C; under a higher current density of 0.5 C and 1.0 C, this value changes to 1884 and 1320 mA h g-1, respectively. Note that even at 2.0 C and 3.0 C, Si-C WBs still have a reversible capacity of 854 and 625 mA h g-1. More importantly, when the current density is returned back to 0.2 C, the reversible capacity raises to 1848 mA h g-1 quickly. This further indicates that Si-C WBs could maintain the structural stability even at a high current density.41 The high volumetric capacity and stable cycling performance are generally believed to be critical parameters in determining battery performance in LIB applications, especially in portable electronics.42,43 However, it remains a big challenge for Si nanostructured anodes to display these parameters, because these materials usually have the large mass specific capacity (by weight), but much smaller volumetric capacity because of their large surface area. Herein, with a relatively high tap density of 0.55 g/cm3, the initial volumetric capacity of Si-C WBs is measured to be as high as 1338 mA h cm-3, which is more than a doubled figure compared to graphite anodes (550 mA h cm-3) and is comparable to a value for the pomegranate-structured Si/C spheres.8,44 In


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addition, this structure affords a relatively stable battery cycling performance even under a higher current density of 1.0 C (Figure 2e), opposed to severe capacity fading for bare Si nanowires (Figure S8, SI). After 1000 cycles, almost 2000 mA h g-1 capacity was retained, which is more than five times of the theoretical capacity of graphite. This further confirms that the present multi-scale buffering engineering of Si-C WBs can be proposed as an advanced technology for making LIB anodes.

Figure 2. Electrochemical tests on Si-C WBs. (a) Their discharge and charge curves at a galvanostatic current density of 420 mA g-1 cycled between the voltage of 1.5-0.01 V vs Li/Li+. (b) Electrochemical impedance spectra of Si NW and Si-C WBs (consisting of Si/C NW) electrodes before and after cycling. (c) Cyclic voltammetry (CV) curves with the cut-off voltage of 1.5-0.01 V vs Li/Li+ at a scan rate of 0.1 mV s-1. (d) Their


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rate performance at various current densities of 0.2 C, 0.5 C, 1.0 C, 2.0 C and 3.0 C. (e) Cycling performance and coulombic efficiency at 1.0 C.

In situ TEM Observations. In order to uncover the storage mechanism of Si-C WBs at the atomic scale, we fabricated a single-particle-based nano-battery prototype utilizing an in-situ TEM technique.45 This nano-scale battery consists of a single Si-C WB and a piece of lithium metal, similarly to our previous reports.10,46 Figure 3 and Supplementary Video 1 show the structural evolution at different lithiation stages of a single Si-C WB. As shown in Figure 3a, at first, the lithiation process takes place at the outermost edge of the Si/C NW, and then gradually moves inward the core, until the darker crystalline core finally disappears. The lithiation process of Si/C NW is quite consistent with the previous studies. Importantly, it can be clearly discerned that the diameter of Si/C NW shows only a 19.5% increase, from 36.4 nm to 43.5 nm, during 150 s lithiation (Figures 3a1-a4). The volume change of Si/C NW here is much lower than that of bare alloying anode made of Si and carbon, and is as small as for the intercalation-type graphite anode.44 After being fully lithiated, the carbon layer is kept intact, without any breaks, and reveals a stable SEI layer forming on the surface (Figure 3b). The other interdependent characteristics of the wool-ball-like design that enable superior battery performance is the sufficient and well-defined internal void space. Such space within Si/C NWs allows the Si to expand freely without disturbing the integrity of a secondary particle, which is really important for maintaining the


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continuity of the electrode framework and overall cell configuration in practical application. Thus, the Si-C WB-based nanobattery was also prepared to analyze volume evolution utilizing in-situ TEM technique. As depicted in Figure 3c and Supplementary Video 2, with the lithiation going on, the internal voids within Si/C NWs gradually disappeared and filled up by the expanded Si/C NWs during 300 s lithiation. After being fully lithiated, the textural characteristic of a secondary particle is still well maintained and shows only a slight volume expansion of 17.4% with the diameter increasing from 3.45 m to 3.64 m (Figure S9, SI). Ex-situ TEM analysis of particle-level Si-C WBs (before and after lithiation) was also conducted, showing that the Si/C NW of a Si-C WB is not only thickened along the radial direction, but also elongated along the longitudinal direction (Figure 3d). Both in situ TEM and ex situ TEM studies of the structural changes under cycling indicate that the internal voids among Si-C WBs became compressed. That is, the special wool-ball-like design could provide enough voids to accommodate the volumetric expansion of Si and reduce the inner volume stresses through the internal voids change, as illustrated in Figure 3e. Moreover, the well-distributed 3D pores shorten the diffusion path of lithium ions and decrease the inner resistance of Si-C WBs, which could effectively promote the lithium ion diffusion during the charge/discharge process. And this fact is well consistent with the superior rate performance of Si-C WBs.


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Figure 3. In situ TEM observation of the lithiation process of an individual Si-C WB. Schematic illustration of in situ TEM electrochemical experimental setup. (a) Timelapse images of a Si/C NW before and after lithiation. (b) HRTEM images of carbon layer and SEI layers after fully lithiated. (c) In-situ TEM images of a Si-C WBs before lithiation and after lithiation for 30 s, 120 s and 300 s, respectively. (d) Ex-situ TEM images of particle-level Si-C WBs before and after cycles.

The superior electrochemical performance for Si-C WBs is believed to be attributed to the wool-ball-like architectures which could nicely buffer the volume expansion and relief the subsequent strains, as illustrated in Figure 4. Firstly, it should be mentioned that the hierarchical Si-C WBs are composed of dozens of coiled Si/C NWs, which have a tripled-like morphology: the central crystal Si (c-Si) core, the sub-shell amorphous Si (a-Si) coated onto the c-Si, and the outermost carbon shell, as was discussed above. Not only the fluffy wool-ball structure could facilitate the lithiation, but also the triple buffering structure of Si/C NW plays a vital role in the enhanced electrochemical


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performance. The in situ HRTEM experiments on a single Si/C NW further demonstrate the lithiation processes. During the initial lithiation, a thin SEI layer, about 2.1 nm thick, is formed, and the amorphous phase exhibits a significant volumetric expansion (Figure 4a2). The diffraction rings of a-Si became weaker and finally changed to a single typical a-LixSi pattern (Figure 4b3, b4), indicating that the a-Si phase converted to amorphous LixSi.47-50

Then the central c-Si core underwent a crystalline-to-amorphous phase

transition during the further lithiation process, and finally changed to a-LixSi phase after lithiation (Figure 4a3). The spots which represent the (220) and (111) planes disappeared quickly, suggesting that Li+ ions preferentially penetrated along and directions at the interface of a-Si and c-Si (Figures 4b1, 4b2). The (220) and (111) planes are also the dominant planes of Si-C WBs based on the XRD results (Figure 1b). The detailed lithiation process from a-Si sub-shell to c-Si core is illustrated in Figure 4c. In addition, it is reported that the diffusivities of lithium ions in a-Si (DLia-Si = 1.25 × 10-9 ~ 3.69 × 10-8 cm2 s-1) are larger than that of c-Si (DLic-Si = 1.67×10-10 ~ 4.88×10-9 cm2 s-1) at room temperature,51 suggesting that lithium ions diffuse faster in a-Si, which further confirms that the Si/C NWs with a-Si sub-shell and c-Si core are favorable for the lithium ions penetration from the surface to the core, thus leading to a perfect rate performance of the Si-C WBs. Furthermore, based on the simulation of volumetric expansion of c-Si and a-Si during lithiation (Figure S10, SI), both Si phases exhibited nearly linearly expansions as the function of Li content, which is consistent with previous reports.52 The similar volume expansions prevent the separation of a-Si and c-Si in the interface during lithiation.


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Secondly, the lithiation behaviors on the a-Si sub-shell and c-Si core were also investigated using DFT calculations. The electron density of states for pure amorphous and crystalline phases and selected LixSi alloy were calculated (Figure 4d-e). It is demonstrated that purely amorphous Si exhibits a larger band gap than that of the crystalline phase (Figure S11, SI), which is due to the strong sp3 hybridization and the tetrahedral structure. When the alloying Li content increases, the Si network loses its continuity and decomposes into smaller fragments, which is caused by decreasing SiSi bonding strength, making the Si material soft and flexible.52 The 3s and 3p states became narrower and the Li-Si alloy bandgap gradually disappeared with increasing lithiation time, leading to a more metallic character; this, in turn, further promotes the lithium ion diffusion. The bulk modulus of Li-Si alloys (as a function of the alloying Li content) was also calculated. As shown in Figure 4f, the bulk moduli of Si-Li alloys decrease with increasing Li content for both amorphous and crystalline Si phases, leading to significant elastic softening, which could relieve the strain during lithiation.53 It is also noted that for a given alloying Li content, the amorphous phase tends to be softer than its crystalline counterpart. We also investigated the lithiation process of the c-Si core. During the initial stages of lithiation (Figures 4a-b), lithium ions penetrate the outermost carbon layer and diffuse through the a-Si sub-shell. We assume that Li ions penetrate the crystalline Si through the crystal direction and induce anisotropic expansion of the Si/C tripled-structure NWs. As the dominant planes of Si-C WBs are (110) and (111), as based on TEM and XRD data, their corresponding vertical directions and


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were simulated as the axially oriented directions to investigate the volume expansion. When lithiation starts, lithium ions preferably diffuse along the direction and cSi exhibits the significant volume expansion on (110) planes, which is consistent with the results reported previously.54 For the axially oriented c-Si NW (Figure 4g), there are six directions arranged hexagonally and perpendicular to the axis. When lithium ions penetrate c-Si, the expansion along six directions leads to a small cross-sectional expansion of Si/C NWs. By contrast, for the axially oriented c-Si NW (Figure 4h), only one direction would heavily expand, leading to the apparently anisotropic expansion of Si/C NWs. This is why the triple-structured (hard-soft-soft) Si/C NWs exhibit a small volumetric expansion along the radial direction. Besides, the outermost carbon layer also restrains the volume expansion and maintains the integrity of Si-C WBs. Finally, thin SEI layers formed on the outermost carbon surface were also analyzed via ex situ XPS, ex situ Raman and ex situ FT-IR. The same techniques were also applied to Si NPs-C balls for comparison. Figures 4i and 4j show the C 1s XPS spectra of Si-C WBs before and after cycling, respectively. There are four peaks at 284.4, 285.6, 286.5, and 289 eV, which can be associated with C-C, C-OH, C-O-C, and O-C=O bonds, respectively. It is worth noting that the strong characteristic peaks at around 290 eV in the C1s spectrum (Figure 4j) and at 54.6 eV in the Li 1s spectrum (Figure S12b, SI) of Si-C WBs after cycling correspond to Li2CO3, suggesting the formation of SEI layers upon the outermost carbon layer. The C 1s and Li 1s spectra of Si NPs-C and Si NWs without carbon layer exhibited similar evolution during the lithiation process (Figures


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S12a, S13, S14). However, the area proportion of the peak at 290 eV in C 1s spectrum of Si-C WBs after cycling is smaller than that of Si NPs-C balls, implying that the SEI layer formed on the Si-C WBs is thinner than that of Si NPs-C balls and Si NWs without carbon layer, which enables the Si-C WBs to achieve excellent ICE. The ex situ Raman spectra of Si-C WBs before and after cycling are shown in Figure 4k. Selected Raman shift window of the Raman bands centered at 509.7, 923.3, 1337 and 1589.1 cm-1 are corresponding to crystalline Si-Si, amorphous Si-Si, D band, and G band. We can apparently observe that every Raman band becomes weaker after cycling. For the Raman spectra of Si NPs-C balls (Figure S15, SI), no obvious amorphous Si-Si bonds can be discerned, and the D and G bands significantly enhanced after cycling, which indicates that the disorder degree of carbon was increased. Ex situ FTIR spectra of the two Si-C materials (Figure S16, SI) also show that the characteristic bands of SEI layer (CO32- of Li2CO3 at 870 cm-1 and Li-O at 520 cm-1) in Si-C WBs spectrum after cycling are less pronounced than those in Si NPs-C balls, further confirming that the outermost carbon layer of a-Si sub-shell and c-Si core tend to form a thinner SEI layer, which facilitates the superior ICE.31 The structure well contributes to the ultra-high ICE of ~90% for Si-C WBs. Very recently, Cui et al.12 have proved that high CE (ICE 93.2%; increases to 99.5% within 5 cycles) can be realized using Si microparticles caging into multi-layered graphene, even for Si microparticles (~1-3 μm). This further suggests that coating of carbon layer on Si particles is a promising strategy to alleviate the huge volume expansion.


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Figure 4. Nano-level characterizations and simulations of Si-C WBs during lithiation. (a-b) In situ TEM characterization of structure and FFT changes of a single Si/C NW at different lithiation stages. (c) Illustrations of Li penetration from outer a-Si to core c-Si. (d-e) Electron density of states (DOS) of a-Li1.67Si and c-Li15Si4 alloys, respectively. The vertical dotted line is the Fermi level position. (f) Bulk moduli for cLi-Si and a-Li-Si alloys at various lithiation stages. (g-h) Entrance area for Li penetration on Si (110) and (111) planes and the schematic illustrations of anisotropic


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expansion of the different Si NWs when Li diffuses along the lateral directions, respectively. (i-k) High-resolution C 1s XPS and Raman spectra of Si-C WBs before and after cycling, respectively. Micro-sized Si-C Wool-Ball Architecture’s Merits toward Improvements of Electrochemical Performance. More importantly, the well-defined internal void space among Si-C WBs results in perfect elastic and restorable characteristics. This is the key to effectively address the urgent issues in LIB practical applications, e.g. particle fracture, electrode thickening and electrical mis-contacting. To substantiate these impressive characteristics, the mechanical behavior on a single-particle level under an external load was examined using a piezo-controlled, electrically biased TEMAFM holder. As shown schematically in Figure 5a, a circuit was built by sandwiching the Si-C WB secondary particle between a conducting Au substrate and a sharp W tip. From Figure 5b-c and Supplementary Video 3, it is obvious that the Si-C WB is resilient to an external load due to its well-defined internal void space and flexibility. Moreover, it completely returns to its original shape after the load is removed, revealing perfect elastic and restorable characteristics. Even after lithiation/delithiation, the Si-C WB secondary particle still shows good elastic resilience to external stress and returns to its original morphology after the load is removed (Figure S17 and Supplementary Video 4, SI). In contrast, when an external load is applied on a relatively compact Si NPs-C ball (Figure 5d-e and Supplementary Video 5), although it shows a certain mechanical strength, the material quickly cracks and crushes after application of an external load just for a few seconds. The distinct elastic quality of the Si-C WB secondary particle


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makes it well suited to address the internal stress induced by Si volume expansion during (de)lithiation and maintain the electrode framework stability. This was confirmed by the cross-sectional scanning electron microscope (SEM) images of the electrode before and after cycles. As shown in Figures 5f-g, the electrode still possesses an intimate contact with the copper current collector without any structural collapse and particle peeling off. The textural characteristics of Si-C WBs based electrodes after cycling are well preserved and show only a slight thickness increase, from 27 nm to 29.5 nm after 50 cycles. These inspiring characteristics can be due to the wool-ball-like design with well-defined internal voids within Si-C WBs, which allow the Si to expand freely without destroying the integrity and continuity of the electrode framework and cell configuration.


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Figure 5. Micro- and particle-level characterizations and simulations of Si-C WBs. (a) Diagram of the electrical circuit for external load testing. (b-e) Schematic and timelapse TEM images of external load testing on particle-level Si-C WBs (b-c, Supplementary Video 3) and relatively compact Si NPs-C balls (d-e, Supplementary Video 5). (f-g) Cross-sectional SEM images of Si-C WBs electrodes before (f) and after 50 cycles (g). Mass Production of Si NWs. Mass production is vital for the practical applications of Si-C WBs.


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Herein, the Si NW precursor was prepared by a radio-frequency thermal plasma (RF-plasma) system, which exhibited a production capacity of 300g/h (SI, Estimation of the maximum feeding rate) using micro-Si as a raw material. The process could yield a 1.5-kilogram product within a single batch (shown in Figure 1e). The thermal plasma synthesis process mainly involves phase conversion, interactions between thermo-fluid fields, electromagnetic fields, and particle concentration fields (Figures 6a-c, Figure S18, SI). Thus, the appropriate raw material size and parameters for the plasma process is of much importance. The raw irregular silicon particles should be firstly melted and evaporated. The subsequent physical and/or chemical reactions can proceed (Figure S19, SI). Therefore, it is important that all micro-sized particles fed in plasma could be evaporated with the assistance from the high enthalpy peculiar to thermal plasma. During the particle evaporation process, both fully-vaporized particles ratios and gasification ratios of raw materials are greatly decreased with the increase of the particles’ size (Figure 6d, Figure S20, SI; estimation of the maximum size of raw materials, SI). Considering that productions with nano-silicon (above 99.0%) are acceptable, the appropriate raw material size is about 30 μm. On the other hand, temperature gradients have a great influence on the morphologies of Si nanoparticles in the thermal plasma synthesis process (Figure 6e). Raw materials delivered into plasma typically go through three stages during the preparation process: vaporization, nucleation/crystal growth and condensation/ coagulation. At the first stage, irregularly micro-sized silicon particles were melted and rapidly evaporated with the help of high enthalpy supplied by a thermal plasma equipment. Then, the vaporized Si materials were transferred to one end of the plasma equipment, where the temperature decreased drastically and Si species uniformly nucleated with a high degree of supersaturation. At the second stage of crystal growth stage, Si nanoparticles with different shapes could be fabricated


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due to various temperature gradients (cooling rates). Under natural cooling condition (i.e. the cooling rate is relatively low), the freshly formed silicon nuclei began to crystalize and then formed irregular nanoparticles. The nuclei growth can be suppressed by introducing an argon quenching gas at high temperature gradients. Then the final product of silicon materials exhibits a spherical shape, because silicon materials endured the minimum surface energy. Whereas, by introducing a graphite tube (as in this work) the temperature gradients are considerably reduced and the crystal growth time is relatively prolonged, which facilitate the nuclei growing according to its growth habit to the thermodynamic equilibrium state, i.e. 1D nanowire. Note that the temperature gradients within the graphite tube are still very high, thus, the nuclei formed at a high-temperature region tend to be in the crystal state, when they transport down to the low-temperature region, while the latter nuclei, deposited on the crystal core, are mainly amorphous. Thus, the specific advantages of thermal plasma make it possible to form crystal-core and amorphous-shell in one-step and at large-scale.


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Figure 6. Theoretical simulations and the mass production process of Si NWs. (a-c) Heat flux distribution, temperature field and axial velocity field in a RF-plasma reaction system, respectively. (d) Gasification ratio of Si particles: 5.0 μm~50.0 μm. (e) Schematic illustration of the possible process of Si products with the different morphologies prepared by RF plasma.

To sum up, the reported superb cycling and rate performance of S-C WBs can be attributed to the specially designed special wool-ball-like design architectures. The core-shell/crystalline-amorphous Si/C NW, the 3D Si-C WB designed frameworks, can significantly reduce the volume expansion during the lithiation/delithiation process. This essentially relieves the inner stress caused by the volumetric expansion of silicon.


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The carbon layer could also accommodate the radial expansion and ensure high ICE through the formation of a stable and thin SEI layer. Moreover, the wool-ball-like design with abundant nanopores provides a sufficient void space to further accommodate the volume changes, and endow it with perfect elastic and restorable characteristics to address the inner volume stresses and maintain the electrode framework stability. In addition, the well-distributed 3D pores can shorten the diffusion path of Li ion and increase the inner resistance, effectively promoting the diffusion of Li ion during the lithiation/delithiation process. Therefore, the present original structure of the Si-C WBs helps to perfectly endure the volume expansion of Si NWs and retain the overall morphology integrity, thus contributing to a good electrochemical performance.

CONCLUSION In summary, being inspired by the structure of a wool-ball, the original wool-ball-like design of triple-structured Si-C WBs is implemented and the material is successfully prepared by a RF-plasma strategy followed by a spray drying treatment. All Si/C nanowires, composed of a c-Si core, a-Si sub-shell and the outermost carbon shell, are hierarchically embedded into a 3D porous wool-ball-like structure (Si-C WBs). This strategy enables scalable production of Si-C WBs at a rate of ~300 g/h. When tested as anode material of LIBs, Si-C WBs exhibit a superior ICE of ~90%, an enhanced capacity of 2433 mA h g-1, an improved volumetric capacity of 1338 mA h cm-3, and a perfect rate performance. In situ HRTEM probing indicates that the hierarchical


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frameworks exhibit a perfect elastic and restorable characteristics, and experience only small volume variations, as per requirements for the secondary particles in the practical applications. The regarded multi-buffering and related storage mechanism are thoroughly investigated. We assume that the structure-design and large-scale synthesis could facilitate the practical application of Si-C composites in LIB anode.

METHODS Synthesis of Si Nanowires (Si NWs). The Si NWs fabrication process is depicted in Figure S1. Si nanowires were prepared by a radio frequency (RF) thermal plasma system in a one-step continuous way. The RF plasma apparatus contains a RF generator (10 kW, 4 MHz), a raw material feeding system, a graphite-stainless steel reactor, an exhaust system, and gas supplying and control systems. Detailed apparatus are shown in the supporting information. The starting powders (Si particles, ~30 μm in average size) were continuously supplied into the plasma flame under a carrier gas flow (Ar, 99.9%). The evaporation occurred in the plasma jet and the as-formed seeds appeared in the graphite tube zone, then the yellowish-brown products were collected from the chamber. The fabrication of nanoparticles by thermal plasma mainly involves phase conversion and interactions among the thermo-fluid fields, electromagnetic fields and particle concentration fields in a short period of time, i.e. a few tens of milliseconds. Thus, the parameters for plasma processing are of much importance. The typical operating parameters are given in Table S4. Synthesis of Si-C WBs. The prepared Si NWs were used to fabricate Si-C WBs


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using a spray drying process. Certain amounts of Si NWs and glucose were dispersed in a deionized water (Si NWs: glucose: H2O = 1: 0.2: 10 wt%) using an ultrasonic bath for 1 h, subsequently a homogeneously dispersed brown suspension was prepared. Then, it was processed in a spray drying machine (YC-015, Shanghai Pilotech Instrument and Equipment Co., Ltd., China) with an inlet temperature of 230 ºC. When the precursor became atomized, a dry brown powder was instantly formed in the spray dryer collector. The precursor spheres (denoted as Si-C WB) were gathered and subsequently calcined at 1200 ºC for 3 h in argon atmosphere. Material Characterization. The materials were analyzed by X-ray diffraction (XRD, Philips X’ Pert PRO MPD) using the Cu Kα radiation at 40 kV and 30 mA. The product morphologies were studied using a field-emission scanning electron microscope (FESEM, JSM-6700F, Tokyo, Japan) equipped with an energy-dispersive X-ray spectrometer (EDS). The detailed structural features were uncovered by means of transmission electron microscopy (TEM, JEOL JEM-2100, Tokyo, Japan). The sample porosity was measured using physical adsorption of nitrogen at -196 C on a Brunauer-Emmett-Teller (BET) surface area analyzer (Micrometric, ASAP 2010). A Raman spectrometer (Renishaw, England) with an excitation wavelength of 514.5 nm and a beam spot size of 1-2 mm was employed to characterize the specimens. Thermogravimetric analyses (TGA) of Si-C WBs were conducted with a TG-209 F3 instrument (NETZSCH, Germany) at a heating rate of 10°C min-1 in air. For ex situ SEM/TEM analysis of the variations in electrode thickness and surface morphology after cycling, coin cells were disassembled and the electrodes were rinsed with dimethyl


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carbonate in an argon-filled glove box. Ex situ X-ray photoelectron spectroscopy (XPS, ESCALAB250XI, Thermo Scientific) and ex situ Fourier transform infrared spectroscopy (FTIR, TENSOR-27, Bruker) were used to trace the composition evolution of electrodes before and after lithiation. DFT Calculations. The first principles computations were peformed with the Vienna Ab-initio Simulation Package (VASP),55 where projected-augmented-wave (PAW) method was adopted.56 The functional of Perdew, Burke, and Ernzerhof (PBE) under the generalized gradient approximation (GGA) was used in the calculations.57 The energy cutoff was fixed as 350 eV. The convergence of energy and force were set as 10−4 eV and 0.02 eV/Å, respectively. In Situ TEM Analysis. In situ transmission electron microscopy (TEM) probing was carried out by means of a JEOL-3100 FEF TEM having an Omega filter and equipped with a “Nanofactory Instruments AB” STM-TEM holder. In order to prepare the prototype cell, an individual Si-C WB was attached to the gold electrode, which was placed on the piezotube. A tiny piece of Li foil was attached to the sharply etched W wire as a counter-electrode. A grown Li2O layer on the Li surface served as a solid electrolyte. The lithiation was processed at a negative bias of -3 V to 0 V with respect to Li. For mechanical load testing, the piezo-controller pushed the W tip toward the SiC WB. The reversible deformation confirms the material elasticity. Electrochemical Test. CR2025 coin-type half-cells were assembled in an argonfilled glove box to probe the material electrochemical properties. The working electrode was fabricated by mixing the active materials, acetylene black, and a


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carboxymethylcellulose (CMC) binder at a weight ratio of 8:1:1. A deionized water was used a solvent. Homogeneous slurry was prepared after mixture stirring for 45 min, and its casting onto a Cu foil current collector. A Si-C WB film on Cu foil was dried at 80C for 24 h and then slice into disks with diameters of 14.0 mm. The disks were then dried at 120C under vacuum for 24 h. Li foils were taken as the counter electrodes and polypropylene microporous films (Celgard 2400) as separators. The liquid electrolyte was 1 M LiPF6 in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1:1, v/v). The galvanostatic charge/discharge tests were performed using a CT2001A LAND testing instrument (Wuhan LAND electronics Co., Ltd, China) in a voltage range of 0.01 and 3.0 V at a current density of 0.1 C, 0.2 C, 0.5 C, 1.0 C, 2.0 C and 3.0 mA g-1. Cyclic voltammograms (CV) were recorded in a voltage range of 0.01-1.5 V at a scanning rate of 0.1 mV s-1 at room temperature.

ASSOCIATED CONTENT Supporting Information: The supporting information is available free of charge on the XXXX. Schematic illustration of the RF-plasma and spray drying system; N2 adsorptiondesorption isotherms and pore size distribution of the Si-C WBs; SEM and TEM images of the Si NWs; Raman spectra and elemental mapping of Si-C WBs and SiG WBs. SEM images and BET of Si NPs-C balls with different Si/C ratios; Tables of ICE of Si-C hybrid anodes using the different carbon precursors and binders; Cycling performance and efficiency of the bare Si NWs and Si NPs-C balls anode; In situ TEM imagines of the Si-C WB before and after full lithiation; The volume


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changes and electron density of states of a-Si and c-Si alloys upon lithiation as a function of Li content; High-resolution Li 1s and C 1s XPS, Raman and FT-IR of Si NPs-C balls and Si-C WBs before and after cycling; TEM images of external load testing on particle-level Si-C WBs after lithiation; Schematic illustration and theoretical simulations of the endothermic process of particles in plasma.

ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (NSFC Nos. 21805282; 11535003; 21878312; 51802013) and by the Beijing Natural Science Foundation (BNSF No.2184126). D.G. is grateful for granting the Australian Research Council (ARC) Laureate Fellowship FL160100089 and QUT Project 322170-0355/51.


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Multi-Scale Buffering Engineering in Silicon-Carbon Anode for

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