Ultrafast-Charging Silicon-Based Coral-Like Network Anodes for

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Ultrafast-Charging Silicon-Based Coral-Like Network Anodes for Lithium-Ion Batteries with High Energy and Power Densities Bin Wang,†,§,○ Jaegeon Ryu,‡,○ Sungho Choi,‡ Xinghao Zhang,§ Didier Pribat,∥,⊥ Xianglong Li,*,§ Linjie Zhi,*,§ Soojin Park,*,‡ and Rodney S. Ruoff*,†,# ACS Nano Downloaded from pubs.acs.org by EASTERN KENTUCKY UNIV on 02/03/19. For personal use only.

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Center for Multidimensional Carbon Materials, Institute for Basic Science (IBS), Ulsan 44919, Republic of Korea Department of Chemistry, Division of Advanced Materials Science, Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea § CAS Key Laboratory of Nanosystem and Hierarchical Fabrication, CAS Center for Excellence in Nanoscience, National Center for Nanoscience and Technology, Beijing 100190, PR China ∥ Nanomaterials for Energy laboratory, Department of Energy Science, Sungkyunkwan University, Suwon 440-746, Republic of Korea ⊥ Laboratoire de Physique des Interfaces et des Couches Minces (LPICM), Ecole Polytechnique, Palaiseau 91128, France # Department of Chemistry, School of Materials Science and Engineering, and School of Energy and Chemical Engineering, Ulsan National Institute of Science and Technology (UNIST), Ulsan 44919, Republic of Korea ‡

S Supporting Information *

ABSTRACT: Fast charging rate and large energy storage are becoming key elements for the development of nextgeneration batteries, targeting high-performance electric vehicles. Developing electrodes with high volumetric and gravimetric capacity that could be operated at a high rate is the most challenging part of this process. Using silicon as the anode material, which exhibits the highest theoretical capacity as a lithium-ion battery anode, we report a binder-free electrode that interconnects carbon-sheathed porous silicon nanowires into a coral-like network and shows fast charging performance coupled to high energy and power densities when integrated into a full cell with a high areal capacity loading. The combination of interconnected nanowires, porous structure, and a highly conformal carbon coating in a single system strongly promotes the reaction kinetics of the electrode. This leads to fast-charging capability while maintaining the integrity of the electrode without structural collapse and, thus, stable cycling performance without using binder and conductive additives. Specifically, this anode shows high specific capacities (over 1200 mAh g−1) at an ultrahigh charging rate of 7 C over 500 charge−discharge cycles. When coupled with a commercial LiCoO2 or LiFePO4 cathode in a full cell, it delivers a volumetric energy density of 1621 Wh L−1 with a LiCoO2 cathode and a power density of 7762 W L−1 with a LiFePO4 cathode. KEYWORDS: fast charging, volumetric energy density, silicon nanowires, interconnection, lithium-ion batteries

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educing the charging time and increasing the energy storage in lithium-ion batteries (LIBs) are urgent challenges to meet in view of the rapidly increasing demands of high-performance electric vehicles. The charging © XXXX American Chemical Society

Received: November 28, 2018 Accepted: January 28, 2019

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Figure 1. Characterization of a N-PSi@C electrode. (a) Illustration of the electrode structure. (b, c) SEM images show the morphology of the electrode, where the dashed lines in panel c indicate the connection between the wires; (d−f) TEM images show the detailed structure of the connections, porous silicon, and carbon sheaths. (g) Elemental maps of Si and C and (h) Raman spectroscopy show the components of the electrode and the amorphous structure of the carbon sheaths.

in local regions (that is, at and near the nanostructures), for the entire electrode, the conventionally used individual silicon− carbon units can be without electrical connections or in poor connection such as through polymeric binders and carbon additives, and these types of electrodes have poorer electron transport efficiency, poorer structural stability of the entire electrode, and lower energy density (due to the addition of binders, carbons, or conducting matrix).22,23,27 To address these issues and thus promote the fast-charging and high-energy and -power performance of batteries, we made an electrode that has a network of interconnected carboncoated porous silicon nanowires as a binder-free anode in LIBs (namely, N-PSi@C; “N” stands for “network”, and “P” stands for “porous”). The configuration of the electrode is illustrated in Figure 1a. The mechanism for the fast charging performance of the electrode is as follows: (i) the interconnected silicon nanowires (SiNWs) build a three-dimensional (3D) network that accelerates the electron transport in the entire electrode.27 (ii) The space between the nanowires and also the nanopores inside the nanowires provide channels for ion transport.22,28,29 (iii) Conformal coating of thin carbon layers reduces the resistance in the system and also provides a barrier to alleviate the thickening of the SEI and accommodate the volume change of the silicon nanostructures.30−33 (iv) The interconnection between the nanowires also prevents them from being individually detached from the current collector that is normally observed in the case of individual SiNWs due to the

rate of the commercial graphite anode in LIBs is limited by sluggish intercalation kinetics upon Li-ion insertion, which usually results in Li plating at a high charging rate.1−3 Moreover, the gravimetric capacity of pure graphite is 370 mAh g−1, which is rather low. Alternatively, silicon has been considered as a promising anode material due to its high specific capacity of 3572 mAh g−14,5 and also a working potential (0.22 V versus Li/Li+) that is comparable to that of graphite anodes.6−8 Although the lithium storage capacity in silicon is high, the fast charging performance of silicon-based batteries suffers from a structural instability during cycling.7 Silicon experiences a large volume change (∼300%) during lithiation and delithiation processes, which leads to the fracture of silicon particles, as well as continuous electrolyte consumption because of repeated solid electrolyte interphase (SEI) formation on freshly exposed material surface, inducing decreased electron- and ion-transport kinetics and, thus, fast capacity fading and poor charging rate capability.9−12 To promote fast charging performance, various strategies have been reported, e.g., the use of porous silicon nanostructures that expose more active surface area to the electrolyte and thus accelerate ion diffusion.13−20 Also, the combination with carbon materials to make core−shell structures with the goals of promoting electron transport and stabilizing the silicon structures during electrochemical reactions has been investigated.7,21−26 While porous silicon and the presence of carbon can improve the reaction kinetics B

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Figure 2. Lithium-storage properties of N-PSi@C and the control anodes. (a) Galvanostatic first charge/discharge curves at 0.05 C-rate with a potential window of 0.005−1.5 V. (b) Rate capability tests followed by cycling stability test at 1 C-rate. (c) Voltage profiles at different Crates of N-PSi@C and control electrodes (insets are electrode pictures after 7 C-rate charge/discharge test). (d) Charge volumetric and specific capacity for 500 cycles at 7 C-rate for the N-PSi@C electrode and (e) the corresponding Coulombic efficiency plot. (f) Cross-section SEM images of control (left) and N-PSi@C (right) electrodes. The images on top are higher magnification. (g) Comparison of literature data for volumetric capacity vs various C-rates.

stress generated at their roots upon alloying with Li.27 As a result, the N-PSi@C anode shows stable fast charging performance and long-term cycling stability (at 7 C-rate and over 500 cycles). We also assembled the anode with commercial LiCoO2 (LCO) and LiFePO4 (LFP) cathodes to construct full cells, which showed high volumetric energy and power densities after long-term cycling.

nanowires, where the connected nanowires with diameter of ∼300 nm exhibit a coral-like structure. A transmission electron microscopy (TEM) image in Figure 1d shows the detailed structure of a junction between two connected nanowires. Besides the porous structure, several dense cores with thickness around 20 nm are also observed in the nanowires. These cores could help to support the framework during initial cycles, and after tens of cycles, they were found to turn into porous structures as explained later this article (in the part about electrochemical measurements).31 The porous character of the silicon nanowires was characterized by high-resolution TEM (HRTEM) (Figures 1e,f, and S1). A thin carbon layer with thickness of ∼2 nm was observed on the surface of the porous SiNWs. The carbon coating was formed through a PECVD process using methane as the carbon source, as such a gas-phase deposition can yield a conformal carbon coating both on the surface of materials and also inside the pores.22 The elemental maps in Figure 1g show a uniform carbon

RESULTS AND DISCUSSION To fabricate the N-PSi@C composites, a plasma-enhanced chemical vapor deposition (PECVD) reactor with a planar type heating element was first used to synthesize the interconnected SiNWs; then an acid etching process was used to render the wires porous, followed by coating the structure with thin carbon layers through another PECVD system in a tubular configuration (see Scheme S1 and more details in the Methods section). Scanning electron microscopy (SEM) images in Figure 1b,c show the interconnected C

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Figure 3. Analysis of charge-transfer kinetics. (a) Electrochemical impedance spectra of N-PSi@C and control electrodes at different stages of cycling and (b) calculated charge-transfer resistance after the 1st, 100th, and 300th cycles. CV curves of (c) N-PSi@C and (d) control electrodes at different scan rates from 0.1 to 1.0 mV s−1 along with b values plotted against the battery voltage of each electrode for anodic scan.

electrodes at C/20 (175 mA g−1) differ in their initial Coulombic efficiency (ICE), in which the N-PSi@C and the control electrode have 93.8% and 88.9% of the ICEs with a charge (delithiation) capacity of 3188 and 2983 mAh g−1, respectively, as shown in Figure 2a. Such a capacity loss probably originates from the freshly exposed cross-section of damaged nanowires during slurry formation and electronic contact loss from the current collector. The N-PSi@C electrode has a stable charge/discharge process at various Crates, where its capacity retention approximates 100% at 1 Crate after 200 cycles (Figure 2b), even though the two electrodes have a similar cycling stability at the slow rate of 0.2 C due to the interconnected structure (Figure S5). Fast Li-ion storage capability usually requires efficient electron transport and Li-ion diffusion deep into the Si structure at the same time. In this regard, the control electrode has an insufficient electron pathway across its composite structure, thus suffering from high polarization and eventual lithium metal plating on the electrode at a higher C-rate, as typically observed in anodes with poor kinetics (Figure 2c).6,7 The N-PSi@C electrode shows a smaller over-potential, even at a 7 C-rate that still delivers a reversible capacity of 1275 mAh g−1 after 500 cycles, as shown in Figure 2d. Its Coulombic efficiency per cycle reaches 99.5% at the second cycle, and the average Coulombic efficiency is 99.9% for further cycles, evidently owing to the highly interconnected structure directly grown onto the current collectors and the thoroughly coated conductive carbon sheath (Figure 2e). Without the use of binders and conductive carbons, the volumetric capacity of the N-PSi@C electrode is 1208 mAh cm−3 at a 7 C-rate based on the ∼10 μm thick electrodes, as

coating on the nanowire. We then used Raman spectroscopy to investigate the composition of the N-PSi@C (Figure 1h). Besides the signal from silicon, carbon was observed with D, G, 2D, and D + D′ bands after the PECVD process, characteristic of an amorphous carbon. Energy-dispersive X-ray spectroscopy (EDS) was used to investigate the elemental distribution of the sample as shown in Figure S2. Only silicon, carbon, and oxygen were observed in the spectrum with weight percentage of 83.1%, 12.1%, and 4.8%, respectively. Note that the Cu signal was from the Cu-based TEM grid. The X-ray diffraction pattern in Figure S3 shows the crystalline silicon on the SS foil (JCPDS card no. 00-027-1402). No obvious diffraction peak was observed for the thin carbon coatings, and the gold signal (JCPDS card no. 04-0784) was obtained due to the ultrathin gold layer deposited on the stainless steel (SS) foils prior to the silicon growth. The N-PSi@C synthesized on SS foils could be directly used as electrodes in the LIBs without using additional binders and conductive additives. As a control experiment, the nanowires were scraped from the foils and mixed with super-P carbon black and polymeric binders to make a composite electrode using a traditional slurry process. The Li storage behavior was first evaluated by cyclic voltammetry (CV), Figure S4, which shows the typical activation process of silicon anodes at the initial cycles, along with the SEI layer formation steps at a voltage higher than 1.0 V. Both anodic and cathodic sweeps and SEI peak currents were quickly saturated in the N-PSi@C electrode, while the curves of the control sample were saturated much more slowly, and the large portion of the reduction current contributed to thick SEI formation. Accordingly, galvanostatic charge/discharge curves of the two D

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Figure 4. Full-cell validation of N-PSi@C anode. (a) Voltage profiles at 0.05 C-rate with a potential window of 3−4.3 V and (b) capacity retention for 200 cycles at 1 C-rate of a full cell consisting of N-PSi@C or control anodes paired with LCO cathodes. (c) Voltage profiles at 0.05 C-rate for the first cycle and at 0.2−10 C-rate for subsequent cycles with a potential window of 1.5−3.9 V and (d) rate capability test followed by cycling stability test at 5 C-rate of a full cell consisting of N-PSi@C anode and LFP cathode. (e) Ragone plot of the two different full cells compared to other fast-charging Si-anode-based full cells.

attempts to investigate such current deconvolution in alloyingtype anode materials because of the small chance of having capacitive current in those anodes. Zhang et al. reported kinetic analysis on crystalline Si for a sodium-ion battery anode with a variation of CV scan rates, and they noted that diffusioncontrolled charge was dominant at the peak voltages.38 In agreement with this observation, both electrodes have b values of approximately 0.5 at 0.25 V, which is associated with first order reaction of amorphous Si (a-Si) with Li to form a-Li2.0Si, suggesting an alloying reaction between Si and Li.39 In the subsequent deep discharge process (a-Li2.0Si to a-Li3.5Si) at 0.1 V, the N-PSi@C electrode still shows dominant diffusioncontrolled current with a b-value of 0.5−0.6. Based on the above results, the relative ratio of each contribution is calculated at different scan rates. The capacitive current gradually develops with increasing scan rates and reaches 49% and 32%, respectively, for the N-PSi@C and control electrodes at 1.0 mV s−1 (Figure S7). Generally, the diffusion-controlled lithium storage corresponds to the reactions that take place in the bulk phase with the ions deeply diffused into the electrode material, while the pseudocapacity is generated by the faster redox charge-transfer reactions that occurred at the active sites on the electrode surface. Because the N-PSi@C electrode showed a higher ratio of capacitive current to the diffusioncontrolled current than the control sample, larger lithium storage occurred on the material surface or in the near-surface region, which rationalizes the observed faster charging. Enhanced charge transfer and stable cycling over hundreds of cycles suggest that the interconnected porous SiNWs form a structure that accommodates large expansion of silicon upon lithium insertion without fracture and contact loss from the current collector. The evolution of the morphology of the NPSi@C materials was monitored at different stages of cycling

shown in Figure 2f. Unlike the various designs for fast-charging anodes,11,34−37 which either presented scant loading levels or high-mass passivation layers or volume loss with void spaces, this highly interconnected porous SiNWs configuration maintains a robust electrode structure even at high-massloading and fast-charging conditions, at least for a half cell configuration (Figure 2g). We studied the electrochemical performance of the interconnected porous SiNWs without the carbon sheath (Figure S6). We found that the interconnections and the carbon sheath both contribute to the fast-charging capability of the electrode, with the interconnections playing a (relatively) more-important role. We suggest that it is a synergistic coupling of the interconnections, porous structure, and conformal carbon sheath that enables the fast charge storage. Panels a and b of Figure 3 compare the interfacial resistance of the two electrodes during the cycling progress by electrochemical impedance spectroscopy (EIS) analysis. According to the EIS spectra, no obvious impedance increase was observed in the N-PSi@C electrode for 300 cycles, while the control electrode had charge-transfer resistance that is 3 times higher with a continuously increasing impedance that is indicative of thick SEI formation. When current responses at different scan rates (0.1−1.0 mV s−1) in CV measurements are recorded, the relative contributions from either the capacitive (k1v) or the diffusion-controlled (k2v1/2) terms can be considered, based on the following eq (Figure 3c,d): i = k1v + k 2v1/2 = avb

(1)

where k1, k2, and a are constants; v is the scan rate; and b is a value ranging from 0.5 (fully diffusion-controlled system) to 1 (fully capacitive system). However, there have been few E

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ACS Nano by ex situ HRTEM and SEM analysis (Figure S8). Figure S8a illustrates the overall structural changes, going from the interconnected frame to a fully merged 3D network. In particular, the porous wire structure was retained with volumeaccommodating mesopores until 10 cycles (Figure S8b−e). Afterward, such pores grew to have a ridged surface morphology and the adjacent nanowires in the network became electrochemically welded through a lithium-assisted fusion, as previously observed in typical high-volume-change anodes such as Ge, Sn, and Si (Figures S8f,g and S9).40−42 The deformation of the N-PSi@C structure thus enters a final stage of becoming a robust 3D network. Although the clustering of nanowires produces a local island-like architecture over the electrode, with a domain size less than 20 μm, the restructured active wires still showed good connection to the current collector in Figure S10a. In terms of electrode stability, fast activation of active materials into this invariant and robust structure is highly desirable for the long-term cycle life of batteries, and this is realized for this N-PSi@C electrode within 50 cycles. The interior open spaces between the wires and the interconnected structure prevent severe volume change in the vertical direction because the thickness is unchanged during cycling (Figure S10b). Recent designs also suffer from large volume changes of at least 30% at the electrode level despite void spaces or porous structure.43,44 In contrast, our porous network was uniformly passivated with dense and uniform SEI layers that improve the Coulombic efficiency per cycle and rate capability (Figure S10c−e). A TEM image of the N-PSi@C electrode after 500 cycles is in Figure S10f and shows that the porous structure was retained after long-term cycling. Although the control electrode showed a similar deformation pattern as the N-PSi@C after cycling, the lack of neighboring nanowires and partially broken interconnections led to slow activation of the active materials (Figure S11). Moreover, the scraped interconnected nanowires formed larger aggregations than that of the N-PSi@C samples due to longer interwire distances caused by integrating the conductive carbons and the rigid polymeric cross-linked network. This led to a large change of thickness of 80% after 100 cycles and the formation of a defective network with moss-like interfaces and thick and porous SEI layers, as shown in Figure S12. The nature of the SEI layers on the two electrodes also differs in their composition of organic and inorganic components (Figure S13). The predominant components of the SEI layers formed on the N-PSi@C electrode are lithium fluoride (LiF)based inorganic species with less polymeric and organic derivatives. It makes the SEI layers robust and prevents the continuous electrolyte decomposition in high-capacity anodes that enable the stable electrolyte architecture.45 However, ethylene dicarbonate portions dominate in the control electrode. With the in situ formation of a porous network in a short period as well as stabilized interfaces, the N-PSi@C electrode yields a long-lasting and fast-charging battery. Practical feasibility of the proposed electrode architecture was evaluated by pairing the N-PSi@C anode with two different cathodes, namely conventional LCO (LiCoO2) and LFP (LiFePO4); all this is summarized in Figures 4 and S14 and Table S1. The areal capacity loading of the full cell reaches ∼2.7 mAh cm−2 for both configurations with an N/P ratio (capacity ratio of negative to positive electrodes) of ∼1.1 (Figure 4a). The high ICE and rapid increase in Coulombic efficiency of the N-PSi@C anode minimize the initial capacity

loss when fabricating the full cell. Pairing with the LCO cathode, stable capacity retention of 91% after 300 cycles at 1 C-rate was achieved, while the control electrode showed severe degradation (Figure 4b). Because continuous lithium consumption occurred in the control electrode by thick SEI formation, lithium solely supplied from the LCO cathode depleted quickly. In addition, the maximum current density of the N-PSi@C∥LCO full cell is limited to a 1 C-rate because of sluggish Li-ion diffusion in the LCO cathode and possible lithium plating at a higher C-rate.46 Thus, a N-PSi@C∥LFP full cell was constructed to further test the fast-charging ability in a near-practical system. The charge/discharge cycling of the full cell was monitored at different C-rates and the long-term stability was assessed at a 5 C-rate (Figure 4c,d). The capacity retention at 5 and 10 C-rate were 71% and 58%, respectively, along with further cycling at the 5 C-rate showing a capacity retention of 88% after 150 cycles. The EIS analysis in Figure S15 further confirms reduced internal resistances in the full cell configurations compared with the control full cell that had a commercial graphite anode and an LCO cathode at a similar capacity loading. No obvious delamination or cracking of the cathodes was observed after 200 cycles (Figure S16). The calculated volumetric energy and power density of the N-PSi@ C∥LFP full cell were thoroughly compared with recent promising achievements in Figure 4e.7,47,48 The highest volumetric energy density from LCO pairing is 1621 Wh L−1 and the power density from LFP pairing reaches 7762 W L−1. If cathode materials are further developed with higher capacity and faster Li-ion-transfer kinetics, the overall battery specification can be further improved to advance the realization of Si-based anodes for fast-charging capability with high volumetric energy density.

CONCLUSIONS An electrode for fast-charging anodes with an interconnected structure consisting of porous Si nanowires sheathed by a conformal carbon coating layer and without polymeric binders and conductive carbons was made, characterized, and tested. This structure facilitated uniform electronic conduction over the electrode and fast Li-ion diffusion that enables fastcharging for this binder-free system; a highly reversible capacity over 1200 mAh g−1 with stable retention over 500 cycles at 7 C-rate was obtained. For the binder-free and interconnected structure, the volumetric capacity significantly improved with an electrode thickness of 10 μm that was maintained even after repeated charge/discharge cycles. A favorable morphological change into the 3D porous network was observed after a short period, providing an activation process that made the electrode robust and resilient to large volume expansion of the Si anode. When paired with LCO or LFP cathodes or both, the N-PSi@C-based full cells achieved high volumetric energy (1621 Wh L−1) and power density (7762 W L−1), which is, to the best of our knowledge, beyond the previous Si-based anodes and offers promise for meeting fast-charging requirements. METHODS Preparation of the Network of Interconnected Porous Silicon Nanowires Coated with Thin Carbon (N-PSi@C). The highly interconnected silicon nanowires were grown on stainless steel (SS, ∼ 5 μm thick, 23 mg cm−2) foils in a plasma-enhanced chemical vapor deposition (PECVD) reactor (A-Tech-South Korea) with an embedded planar type heating element by using silane gas (SiH4) as F

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ACS Nano the silicon source.27 Briefly, a 200 nm thick Au thin film was deposited on the SS foil as the seed for the growth of silicon nanowires by using a thermal evaporator at 2.5 × 10−6 Torr pressure. The samples were then heated at 420 °C under a H2 plasma at a pressure of 600 mTorr for 10 min to break the Au thin film into nanoparticles. SiH4 and PH3 diluted in H2 (H2/SiH4/PH3: 270:30:6 standard cubical centimeters per min, sccm) was introduced to the PECVD chamber at 420 °C to grow at a pressure of 600 mTorr for 60 min, then a gas mixture of H2/SiH4 (120:30 sccm) was maintained for 50 more min to grow the interconnected silicon nanowires before cooling down naturally. The inclusion of PH3 yields n-doped silicon nanowires, which was found to improve the electrical conductance of the interconnected framework and also facilitates the subsequent chemical perforation process. The as-synthesized samples were immersed into 2 vol % aqueous hydrofluoric acid (HF) solution containing 5 mM AgNO3 for 15 min to make the silicon nanowires porous. Following the pore creation, the samples were immersed in 10 wt % aqueous HNO3 for 20 min to remove Ag and then a 2 vol % HF solution for 10 min to remove SiO2. They were finally washed in deionized water and dried in a vacuum. To coat thin carbon layers on the porous silicon nanowires, another tubular PECVD reactor was used. The samples were loaded into a quartz tube in the furnace and heated to 600 °C in a flow of 50 sccm Ar and 10 sccm H2 during 30 min. Next, the plasma generator was turned on, and 20 sccm CH4 was introduced into the chamber for 10 min. After that, the plasma generator was turned off and the supply of CH4 was stopped. The sample was fast-cooled to room temperature in the flow of 50 sccm Ar and 10 sccm H2 and unloaded from the furnace and stored in the desiccator in vacuum. Characterization. SEM (Verios 460, FEI) analysis was used to characterize the structure of the samples at an acceleration voltage of 5 kV and current of 0.8 nA. The dimensions and internal structure of N-PSi@C were determined by TEM (JEOL-2100) and HRTEM (JEOL-2100F) with an acceleration voltage of 200 kV. Raman spectroscopy was performed with a Wi-Tec micro-Raman instrument using a 532 nm laser at ambient temperature. XRD (Smartlab, Rigaku) between 10° and 90° was performed using Cu−Kα radiation (λ = 1.5418 Å). X-ray photoelectron spectroscopy (XPS, ThermoFisher, K-α) was done to obtain qualitative information on the SEI layers after cycling. For electrochemical measurements, the N-PSi@C on SS foils were directly used as working electrodes with lithium foil as the counter/reference electrode and a Celgard 2400 membrane as the separator. As a control, the PSi@C nanowires were scraped from the SS foils and then ground with super-P and poly(acrylic acid) (PAA)/carboxy methyl cellulose (CMC) binder in water solvent in a weight ratio of 90:5:2.5:2.5. Thus, the interconnection between PSi@ C nanowires that was observed in N-PSi@C was severely damaged in the samples of PSi@C. For the preparation of cathodes, the active materials (LCO and LFP, respectively) were ground and mixed together with super-P and polyvinylidene fluoride (PVdF) in the weight ratio of 95:2.5:2.5. The areal capacity loading of both cathodes was ∼2.6 mAh cm−2. The electrolyte was 1.3 M LiPF6 in 3:7 v/v ethylene carbonate (EC) and diethyl carbonate (DEC) with 10 wt % fluorinated ethylene carbonate additives included. For galvanostatic cycling tests, coin-type half/full cells were fabricated using CR 2032type cells (Welcos) in an argon-filled glovebox. Galvanostatic charge/ discharge test results were made with a battery cycler (WBCS3000K8, Wonatech Co., Ltd.). CV measurements were obtained at 0.1−1.0 mV s−1 from 0 to 1.5 V (Biologic VMP3). EIS measurements were carried out between 100 kHz−0.1 Hz with an amplitude of 10 mV at a fully lithiated state (∼0.01 V). The charge-transfer resistance was calculated based on the nominal electrode area instead of the interface area between the active material and the electrolyte because the accurate measurement of the interface area is challenging and the comparison between our sample and the control sample assembled in the same type of cells makes more sense. Unless otherwise noted, all the SEM and TEM images of the cycled electrodes were taken after removal of the SEI layer. To remove the SEI layers, the residual electrolytes were first removed by immersing the electrode in the excess ethylene carbonate for 10 h. Then, a drop of 1 mM acetic acid

was used to remove the SEI layer for a few seconds and immediately diluted in distilled water to prevent further damage of the SS substrate.49

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.8b09034. A scheme showing the preparation of the anode material; figures showing HRTEM images, EDS and XRD results, cycle retention results, SEM and TEM images; XPS analysis of the electrodes, and electrochemical properties and morphology of the cathodes; a table showing a summary of the electrode configuration and full-cell energy calculation (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail: *E-mail: *E-mail: *E-mail:

[email protected]. [email protected]. [email protected]. ruoffl[email protected] and rsruoff@ibs.re.kr.

ORCID

Bin Wang: 0000-0001-9576-2646 Didier Pribat: 0000-0002-5539-2051 Xianglong Li: 0000-0002-6200-1178 Linjie Zhi: 0000-0003-2042-2780 Soojin Park: 0000-0003-3878-6515 Author Contributions ○

B.W. and J.R. made equal major contributions to this work.

Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work was supported by the Institute for Basic Science (IBS-R019-D1) and the Center for Advanced Soft-Electronics funded by the Ministry of Science, ICT and Future Planning as Global Frontier Project (CASE-2015M3A6A5072945) as well as the National Natural Science Foundation of China (grant no. 51425302) and Youth Innovation Promotion Association CAS (no. 2016033). REFERENCES (1) Takami, N.; Satoh, A.; Hara, M.; Ohsaki, T. Structural and Kinetic Characterization of Lithium Intercalation Into Carbon Anodes for Secondary Lithium Batteries. J. Electrochem. Soc. 1995, 142, 371−378. (2) Levi, M. D.; Aurbach, D. Diffusion Coefficients of Lithium Ions during Intercalation into Graphite Derived from the Simultaneous Measurements and Modeling of Electrochemical Impedance and Potentiostatic Intermittent Titration Characteristics of Thin Graphite Electrodes. J. Phys. Chem. B 1997, 101, 4641−4647. (3) Yu, P.; Popov, B. N.; Ritter, J. A.; White, R. E. Determination of the Lithium Ion Diffusion Coefficient in Graphite. J. Electrochem. Soc. 1999, 146, 8−14. (4) Li, H.; Huang, X.; Chen, L.; Wu, Z.; Liang, Y. A High Capacity Nano -Si Composite Anode Material for Lithium Rechargeable Batteries. Electrochem. Solid-State Lett. 1999, 2, 547−549. (5) Chan, C. K.; Peng, H.; Liu, G.; McIlwrath, K.; Zhang, X. F.; Huggins, R. A.; Cui, Y. High-Performance Lithium Battery Anodes Using Silicon Nanowires. Nat. Nanotechnol. 2008, 3, 31−35. G

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DOI: 10.1021/acsnano.8b09034 ACS Nano XXXX, XXX, XXX−XXX

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DOI: 10.1021/acsnano.8b09034 ACS Nano XXXX, XXX, XXX−XXX