Putting Nanoarmors on Yolk–Shell Si@C Nanoparticles: A Reliable

Putting Nanoarmors on Yolk–Shell Si@C Nanoparticles: A Reliable Engineering Way To Build Better Si-Based Anodes ... Publication Date (Web): June 27,...
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Research Article Cite This: ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

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Putting Nanoarmors on Yolk−Shell Si@C Nanoparticles: A Reliable Engineering Way To Build Better Si-Based Anodes for Li-Ion Batteries Jian Jiang,*,†,§,⊥ Han Zhang,†,§,⊥ Jianhui Zhu,‡ Linpo Li,†,§ Yani Liu,†,§ Ting Meng,†,§ Lai Ma,†,§ Maowen Xu,†,§ Jinping Liu,*,∥ and Chang Ming Li*,†,§

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†

Institute for Clean Energy & Advanced Materials, Faculty of Materials and Energy, Southwest University, Chongqing 400715, P. R. China ‡ School of Physical Science and Technology, Southwest University, No. 2 Tiansheng Road, BeiBei District, Chongqing 400715, P. R. China § Chongqing Key Laboratory for Advanced Materials and Technologies of Clean Energies, Chongqing 400715, P. R. China ∥ School of Chemistry, Chemical Engineering and Life Sciences, Wuhan University of Technology, Wuhan, Hubei 430070, China S Supporting Information *

ABSTRACT: Practical utilization of silicon (Si) for Li-ion batteries (LIBs) still remains sluggish because of its formidable kinetic problems of huge volume expansions over 300%, instable solid electrolyte interphase (SEI), and unsatisfactory electrical conductivity. Though using a yolk−shell (Y−S) Si@C nanodesign indeed helps to mitigate active changes, optimize SEI properties, and lower intrinsic charge-transfer impedances, the total anodic behaviors in reversibility, rate capabilities, and long-lasting cyclability are still far from perfect. To settle the above issues, we herein propose a reliable and effective way by putting tough and malleable Ni nanoarmors on Y−S Si@C nanoparticles (Si@C⊆Ni). The unique functionalized configurations endow such hybrid systems with superb reversible capacity retention (almost no capacity decay emerges in 600 cycles, retaining a reversible capacity beyond ∼1307 mA h g−1), prominent cyclic stability, and rate behaviors. To justify their potential usage, full cells of (−)Si@C⊆Ni//LiFePO4(+) are further constructed, delivering impressive specific energy and power densities (max. values: ∼423 W h kg−1/∼497.8 W kg−1). This paradigm work may offer a highly feasible engineering protocol to push forward Si anode performances for next-generation LIBs. KEYWORDS: Si@C, ⊆Ni, Ni nanoarmors, anode, full cells, LIBs

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INTRODUCTION Silicon (Si) has long attracted worldwide attention as a “shining star” anode for Li-ion batteries (LIBs) on account of its overwhelming merits of (i) great abundance in nature, (ii) safe working potential range (0.4−0.5 V vs Li/Li+, less Li dendritic crystal formation), and (iii) high theoretical gravimetric/volumetric capacity up to 4200 mA h g−1/9786 mA h cm−3.1−3 Nevertheless, as every coin gets two sides, such high-capacity anodes suffer from dramatic volume changes (>300%) in lithiation/delithiation, giving rise to active collapse/pulverization and further disconnection with current collectors.4−6 In parallel, undesired irreversible consumption of Li+ and electrolyte would proceed because of continuous reformation of solid electrolyte interface (SEI) layer on damaged Si surfaces.7 Besides the critical volume expansion issues, Si anodes are also subjected to sluggish electrochemical kinetics led by their intrinsically poor electrical conductivity.8,9 All of these factors would deteriorate Si electrode behaviors (e.g., capacity retention, long-time cyclic lifespan, reversibility, rate capability, etc.) in a synergistic way, making them far from © 2018 American Chemical Society

satisfactory to meet essential commercial criterions of LIBs.11,12 To tackle these formidable issues, tremendous efforts have been devoted to promoting Si anode behaviors using manifold state-of-the-art nanostructures (e.g., Si nanowires/nanospheres/nanotubes, ordered mesoporous Si nanoparticles, etc.) together with functionalized electrode configurations.6 Particularly note that the intriguing yolk−shell (Y−S) Si@C nanostructures hold a great promise in practical applications because the void space in each hybrid unit provides sufficient room to accommodate Si volume expansions.13−15 Moreover, the outer uniform and intact C shell avoids the direct contact of electrolyte to inner Si nanoactives, preventing them from negative side reactions and hence helping to build a more electrochemically stable and durable SEI layer.16−18 However, this goal realization has to be based on an ideal fundamental Received: May 10, 2018 Accepted: June 27, 2018 Published: June 27, 2018 24157

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

Research Article

ACS Applied Materials & Interfaces

(80 mL) and held at 95 °C for 6 h in an oven. Later, the as-formed Si@C⊆Ni(OH)2 intermediate was collected and washed with deionized water several times. The preparation of ultimate Si@ C⊆Ni hybrids was achieved in a horizontal quartz-tube furnace system. Si@C⊆Ni(OH)2 powders (0.3 g; predispersed in 10 mL of distilled water, dropped onto a ceramic boat, and dried at 60 °C) were placed in the center of a quartz tube. Then, 1.5 mL of ethylene glycol (EG) loaded in an alumina boat was put at the upstream zone of the quartz tube (the distance from the alumina boat to the quartz tube center: 15 cm). Prior to heating, the tube reactor was sealed and flushed with Ar flow (200 sccm) for 15 min. The furnace was then heated to 350 ± 10 °C at a heating rate of ∼10 °C min−1 under a constant Ar flow of 80 sccm, kept for 15 min, and allowed to cool down naturally. Characterization Techniques and Battery Testing. Morphological and crystalline structures of the samples were characterized by a JEOL JSM-7800F field emission scanning electron microscope with energy-dispersive X-ray spectroscopy (EDS) and a JEM 2010F highresolution transmission electron microscope. X-ray powder diffraction (XRD) patterns were recorded on a Bruker D8 ADVANCE diffractometer using Cu Kα radiation. X-ray photoelectron spectroscopy (XPS; Thermo Electron, VG ESCALAB 250 spectrometer) measurement was conducted using Al Kα radiation with the fixed analyzer transmission mode. The Raman spectrum was recorded by using an Ar laser Raman spectroscope (Witec, CRM200) with a 532 nm excitation laser. Thermogravimetric (TG) analysis was performed on an SDT600 apparatus at a heating rate of 10 °C min−1 in air. To estimate the content of C, Si, and Ni in Si@C⊆Ni hybrid products, 1 g of fresh Y−S Si@C⊆Ni hybrid powders was put into a 100 mL glass beaker containing 60 mL of 3 M hydrochloric acid solution. To guarantee the complete removal of Ni, the mixture solution was magnetically stirred for at least ∼30 min at a low rotation rate. Later, the samples were washed with deionized water several times and carefully harvested using the vacuum filtration method. According to mass differences, the metallic Ni statistically took up ∼10.2% in the Si@C⊆Ni hybrids. The above collected powders were further analyzed by TG measurement in the air condition in a temperature range of ∼20−600 °C. The electrode mass was measured on a microbalance with an accuracy of 0.01 mg (OHAUS EX125DZH, USA). The hybrid anode of Si@C⊆Ni was made by the conventional slurry-coating method, where Si@C⊆Ni nanopowders, poly(vinylidene fluoride) binder, and carbon black were mixed in a mass ratio of 80:10:10 and homogenized in N-methyl-2-pyrrolidone. The as-formed homogeneous slurry was then pasted onto a Cu film (thickness: 0.4 mm) and dried at 120 °C for 12 h under a vacuum condition, and then the anode films were compressed under a pressure of ∼0.5−1 MPa. The mass loading on each current collector was controlled at a level of ∼2−3.5 mg/cm2. The electrode mass matching is conducted by referring to the electrochemical behaviors of the Si@C⊆Ni anode in half-cell testing and the theoretical capacity of the LiFePO4 cathode. Prior to full-cell assembly, the Si@C⊆Ni anode is attached to a piece of Li metal foil and then immersed in the electrolyte for the prelithiation treatment (lasting time: 3−5 h). Later, this anode is peeled off from the Li foil carefully and assembled with a LiFePO4 cathode to constitute a full cell. The specific capacity calculations of full-cell LIBs are based on the total mass of both anode and cathode (to ensure the cyclic stability of (−)Si@C⊆Ni// LiFePO4(+) full-cell LIBs, the best mass ratios are estimated around 1:6.2). Electrode testing was performed using CR-2032 coin-type cells in a potential window of 0.005−3 V. Cells were assembled in an Arfilled glovebox (Mikrouna Super; H2O < 0.1 ppm, O2 < 0.1 ppm) with Li foil as the counter and reference electrodes. A solution of 1 M LiPF6 dissolved in a 1:1 (v/v) mixture of ethylene carbonate and diethyl carbonate was used as the electrolyte. In full-cell testing, the commercialized LiFePO4 nanopowders were employed as LIB cathodes (the electrode fabrication was similar to that of the anode except that the slurry was pressed onto an Al foil; active mass loading: ∼8−12 mg cm−2). Galvanostatic CD tests were performed by a professional battery tester (Neware, Shenzhen). Electrochemical impedance spectroscopy (EIS) and cyclic voltammetry (CV) scans

prerequisite that C layers are “extremely” robust enough to tolerate Si electrode stress and can always keep an intact state to impede the electrolyte access.19,20 The actual individual use of the Y−S Si@C configuration, though indeed benefiting Si reaction kinetics in reversible capacity, cyclic stability, Coulombic efficiency (CE), etc., still fails to address the above issues once and for all because of C shell fatigue and damages in repeated charge/discharge (CD) procedures.22−24 The mechanical property of Si@C hybrid anodes, thereby, requires to be strengthened so as to quicken their practical utilization pace.25 Making the Y−S Si@C nanoparticles dressed in malleable “nanoarmors” is conceived as a smart and potential route to upgrade the overall anodic behaviors. We herein propose a facile strategy to make Y−S Si@C nanoparticles firmly trapped in a durable and ultrathin Ni matrix (denoted as Si@C⊆Ni) for LIB applications. This functionalized electrode design may bring about several advantages. First, tenacious Ni nanoarmors conformally worn on every Si@C hybrid nanounit can well cooperate with C interlayers, exerting positive effects to suppress Si volume expansions and adverse aggregation effects. Second, the single layer of Ni film is ultrathin, permeable (Li+ can readily pass across the metallic film matrix via structural defects and reach inner deep regions of the whole anode), and lightweight (total weight ratio of involved Ni: ∼10.2%, calculated from mass differences between the pristine Si@C and the ultimate Si@C⊆Ni), ensuring the high gravimetric energy density for LIBs. Third, conducting Ni nanoframeworks tightly surrounding Si@C nanoparticles can also function as nanosized current collectors, capable of accelerating electron transfer from redox reaction sites (along special nanocables) to the external circuit. Thanks to this unique Y−S electrode construction,27−29 such hybrid anodes show superb reversible capacity retention (the reversible capacity always stays above ∼1307 mA h g−1 without obvious capacity fading in a total of 600 cycles) and outstanding cyclic stability and rate capabilities. To evidence their potential utilization, we further make full-cell LIBs of (−)Si@C⊆Ni//LiFePO4(+), which also exhibit remarkable specific energy and power densities. Our success in this hybrid electrode engineering may open up the possibility of making high-performance Si-based anodes for advanced LIBs and set up a smart platform for the design of high-level nanoparticle ⊆ metal matrix hybrid systems that might be available for other practical fields.

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EXPERIMENTAL SECTION

All reagents and solvents involved were of analytical grade and used without any additional purification. Synthesis of Y−S Si@C Nanoparticles. Typically, 0.5 g of battery-grade Si nanopowders (size: ∼50 nm and purity: 99.9%, purchased from Shanghai Gerxin Nanotechnology Co. Ltd) and 0.25 g of dopamine (DA) molecules (Adamas) were successively dispersed into 400 mL of Tris-buffer (pH 8.5) solution and treated by ultrasonication for 1 h. Afterward, the Si@PDA intermediates were then collected by vacuum filtration, washed with deionized water several times, and dried at 60 °C in a vacuum oven. Next, 0.5 g of Si@ PDA and 1.5 g of NH4F (Sigma-Aldrich) were mixed by grinding in an agate mortar for 20 min. The ultimate Y−S Si@C hybrid products were synthesized via heating the above mixture under an Ar atmosphere at 500 °C for 1 h. Synthesis of Si@C⊆Ni Nanoparticles. Typically, 0.6 g of Y−S Si@C nanoparticles, 0.5 g of hexamethylenetetramine (SigmaAldrich), 0.25 g of Ni(NO3)2·6H2O (Sigma-Aldrich), and 50 mL of distilled water were mixed and treated by ultrasonication for 15 min. The resulting suspension was then transferred into a sealed container 24158

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

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Figure 1. (a) General schematic displaying the entire evolution of Si@C⊆Ni. SEM images of samples at distinct synthesis stages: (b) Si nanoparticles, (c) Y−S Si@C intermediates, and (d) Si@C⊆Ni hybrids. (e−h) TEM observations on the ultimate Si@C⊆Ni hybrids.

Figure 2. (a) EDS detection and (b−e) elemental mappings of Si@C⊆Ni hybrids. (f) XRD patterns of evolved samples. (g−i) High-resolution XPS spectra of Si@C⊆Ni hybrids: (g) Si 2p, (h) C 1s, and (i) Ni 2p3/2. were evaluated by an electrochemical workstation (CorrTest CS310). Before battery testing, all cells were aged for 8 h. The specific energy and power densities (E and P) of the full cells were calculated by the following equations: E = U ·Q

(1)

P = E /Δt

(2)

scanning electron microscopy (SEM) image is displayed in Figure 1b). Step I is the controllable synthesis of Y−S Si@C nanohybrids. Si nanoparticles are evenly dispersed into a DAinvolved Tris-buffer solution, where DA molecule polymerization would proceed, making each Si nanoparticle enclosed with PDA layers. The Y−S Si@C hybrids are then generated by a heat treatment of Si@PDA/NH4F powder mixtures in an Ar atmosphere. NH4F salt was added because its thermal decomposition product (e.g., HF gas) can gently etch inner Si nanocores, leading to the scalable yield of Si@C hybrid products with interior hollow architectures.30,31 Figure 1c definitely shows their geometric morphology; nearly all Si nanoparticles (typical diameter: ∼40 nm) are perfectly confined in nanosized C capsules. In step II, the as-formed Si@C samples are evenly packaged within nickel nitrate hydroxide (NNH) layers via a simple solution method

where U is the operating voltage for cell discharging (V), Q is the specific cell capacity (A h·kg−1), and Δt represents the discharging time (h).

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RESULTS AND DISCUSSION The schematic in Figure 1a shows the overall synthesis flow that involves two major procedures. To shorten the gap between the laboratory study and industrial developments, we purposely select the commercial battery-grade Si nanoparticles as starting materials (central size: ∼50 nm; their typical 24159

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

Research Article

ACS Applied Materials & Interfaces

Figure 3. (a) CV curves, (b) long-term cyclic performance, (c) CD profiles, and (d) rate capabilities of Si@C⊆Ni hybrid anodes. (e) SEM observation and (f−i) EDS elemental mappings on cycled electrodes after 600 cycles.

chemical bonds of Si−Si and Si−C, respectively. The remaining one at ∼102.1 eV should result from the Si−O bond of SiOx species that are inevitably yielded on highreactivity Si nanoparticles when exposed to air.10,15 Nevertheless, the relatively high intensity/ratio of the Si−Si bond discloses that the elemental Si takes up the dominant component. For the XPS spectrum of C 1s (Figure 2h), peaks fitted at ∼284.7, ∼285.8, and ∼288.2 eV are successively assigned to C shells with chemical bonds/functional groups of C−C, C−OR, and COOR, respectively.25,26 The Ni 2p3/2 XPS spectrum (Figure 2i) comprises a strong peak at a low BE of 852.6 eV along with a wide satellite at a high value of 858.7 eV, both of which are fingerprint signals in the literature for metallic Ni0.26 The Raman spectrum (Figure S1) reveals that there are two sets of peak signals. Peaks at the wavenumbers of ∼542.8 and ∼980.7 cm−1 are mainly induced by Si-related vibration modes, whereas the ones located at ∼1347 and ∼1578 cm−1 match well with the D and G vibration modes of graphite, respectively. The relatively high integrated intensity and IG/ID ratio (1.45:1) suggest the good graphitization degree of C interlayers (less disordered or defective/amorphous C is contained). To further determine the content of C and Si in Si@C⊆Ni hybrid products, the powder samples are in turn subjected to the acid-immersion treatment and TG measurement (Figure S2) given the fact that Si can be highly stable up to 600 °C when heated in air atmospheres.22 As reflected, the total Si content is thereby estimated to be ∼70.6% while the C content is ∼19.2%. The unique hierarchical structural features of Si@C⊆Ni may help to facilitate electron transport, stabilize the interfacial properties, and withstand Si volume variations, making this hybrid system intriguing for LIB applications. The anodic characteristics of Si@C⊆Ni are initially evaluated by CV at a scan rate of 0.1 mV s−1 between 0.005 and 3 V (vs Li/Li+). As recorded (Figure 3a), the intensive reduction peak emerging at ∼0.69 V in the first cathodic scan (but absent in later CV cycles) is related to the SEI layer formation on the C shell surface, whereas the strong reduction peak lying below ∼0.25 V is highly associated with Li ion insertion into C and Si (Si meantime alters from the crystalline phase into an amorphous state). The following anodic scan involves two oxidation peaks at the potentials of ∼0.36 and ∼0.58 V in remaining CV scans; the gradual increase in current intensity indicates that more Si substances become activated. Long-term galvanostatic cyclic test on Si@C⊆Ni anodes has been conducted (Figure 3b; for

followed by a mild in situ reduction process, resulting in the ultimate formation of Si@C⊆Ni (total size: ∼60 nm; see Figure 1d). The choice of EG as the reductant is driven by the fact that it enables a gentle reduction reaction (via alcoholysis or polyol process) to alter the Ni2+ oxidation state and structures on the atomic scale whilst still keeping the framework topology on the mesoscale.32 In transmission electron microscopy (TEM) observations on Si@C⊆Ni, the sharp contrast in Figure 1e implies that Si@C hybrids are uniformly embedded in interconnected Ni nanoframeworks, highly consistent with our previous SEM analysis. Zoom-in TEM observation (Figure 1f) on broken regions of Si@C⊆Ni further verifies the hierarchical structures where Si nano yolks are well packaged in a double-shell architecture. Highresolution TEM images (Figure 1g,h) uncover that both Ni films and C interlayers (thickness: ∼6−10 nm) are continuous and integrated, much favorable to tolerate mechanical stresses induced by inner Si actives; a lattice distance of ∼0.19 nm (Figure 1h) is definitely identified as the (220) face of Si. The above TEM results fully confirm the successful realization of Si@C nanoparticles dressed in Ni nanoarmors. EDS coupled with elemental mapping technique is used to characterize the final samples. All detectable elements involve C, Ni, Si, and O (see the inset in Figure 2a; EDS elemental mappings (Figure 2b−e) further affirm their uniform distribution). The slight elemental O signal may chiefly stem from either residual EG [chemical formula: (CH2OH)2] molecules or their derivatives yielded in heating procedures. Figure 2f displays the XRD patterns of samples in distinct evolution stages. For all cases, three sharp diffraction peaks lying at 2θ of 28.3°, 47.0°, and 55.8° are indexed to the (111), (220), and (311) facets of Si (JCPDS no. 27-1402), respectively. While in the green pattern, peak signals present at the positions of 44.4° and 51.8° match well with the standard diffraction patterns of metallic Ni (JCPDS no. 040850), solidly evidencing the in situ transformation of NNH layers into Ni. Besides, the presence of a weak peak at ∼25.3° in both green and blue patterns should be correlated with the (002) face of PDA-derived hard carbon (JCPDS no. 65-6212); no noticeable diffraction peaks from possible impurities (e.g., SiOx and NiO) are detected. The components and valence state information of Si@C⊆Ni hybrids are further determined by XPS (Figure 2g−i). In the high-resolution XPS spectrum of Si 2p, the deconvoluted peaks at binding energies (BEs) of ∼99.1 and ∼100.8 eV correspond to Si nanoparticles in the 24160

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

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Figure 4. (a) CV curves, (b) cyclic records under varied current densities (inset: CD plots), and (c) Ragone plot of (−)Si@C⊆Ni//LiFePO4(+). (d) Optical images showing that a single full cell can drive a high-power blue LED light.

believe that such a unique hierarchically functionalized electrode architecture plays a dominant role. On the one hand, the inside Y−S Si@C nanostructures provide rich redox reaction sites for Li storage, help to digest Si volume changes, and suppress nanoparticle aggregations upon deep cycling. On the other hand, the incorporation of Ni nanoarmors intimately surrounding Si@C nanohybrids promotes the anodic mechanical properties and electrical conductivity, impedes the electrolyte entrance, and thereby facilitates the SEI film recovery at crushed/cracked places by obeying the “selfhealing” mechanism.8 The rate behavior of Si@C⊆Ni electrodes at programmed current rates is studied (Figure 3d). Upon an increase in the current density, the electrode exhibits reversible capacities of ∼1370 mA h g−1 (0.2 A g−1), ∼1164 mA h g−1 (0.6 A g−1), ∼939 mA h g−1 (1.2 A g−1), ∼718 mA h g−1 (2.4 A g−1), and ∼449 mA h g−1 (4.8 A g−1). Even at a high rate up to 9.6 A g−1 (48 times the initial value of 0.2 A g−1), the half-cell capacity value (fully charged in few minutes) still exceeds ∼300 mA h g−1. Typical SEM images of cycled electrodes are also present in Figure 3e (more additional SEM observations are available in Figure S5). Though undergoing a long and fatigue cyclic period, such hybrid anodes still remain highly dispersive and own average dimensions below ∼300 nm in the absence of bulky (size: >5 μm) aggregation formation, signifying the superior mechanical/structural stability of Si@C⊆Ni when compared to the case of Y−S Si@C counterparts (see the comparative SEM observations of cycled Si@C anodes in Figure S6) for Li storage. EDS elemental mapping records (Figure 3f−i) on a selected region (see the SEM image in Figure 3f) distinctly uncover that the cycled electrode still preserves integrated configurations in which Si nanoparticles are firmly wrapped by a Ni/C double-shell architecture. The TEM images (Figure S7) illustrate that the morphology for cycled electrodes deforms a lot (typical size: 160 nm, with a rise by 167%);

comparative study, cyclic tests of Y−S Si@C and bare Si nanoparticles are performed as well). At a low rate of ∼0.3 A g−1, the original discharge/charge capacities of the Si@C⊆Ni electrode are recorded around ∼2327 and ∼1382 mA h g−1, respectively, with an initial CE close to ∼60% (the CE records for Y−S Si@C and bare Si nanoparticles are also provided in Figure S3). Afterward, its reversible specific capacity is stabilized at a high value of ∼1345 mA h g−1 (fluctuation range: ∼1309−1387 mA h g−1) until the end of the cyclic test (CE always remains above ∼99.3%). By sharp contrast, both Y−S Si@C and bare Si nanoparticles suffer from evident capacity degradations in long-time operation. The gravimetric capacity for Y−S Si@C anodes only retains a small value of ∼465 mA h g−1 after 600 cycles, whereas for bare Si electrodes, there is almost no capacity retained even after a short cyclic period (∼100 cycles). The corresponding CD profiles of Si@ C⊆Ni are shown in Figure 3c. In the first discharge process, the voltage decreases steeply from the open-circuit voltage to ∼0.52 V wherein a long and flat sloping voltage plateau sets in and continues until a specific capacity of ∼2327 mA h g−1 is achieved. This may arise from SEI film formation and accords well with our CV records aforementioned. In subsequent charge profiles, a strong polarization plateau begins at ∼0.5 V, which is attributed to the dealloying of Li4.4Si as well as Li+ extraction from C. The EIS spectra of anodes with distinct electrode configurations before and after 600 cycles are measured (Figure S4) so as to gain deep insights into electrode kinetics. A semicircle in the high-frequency range representing the SEI surface/charge-transfer (RSEI+ct) resistance appears in all electrode cases, with fitted values of ∼396 Ω (for Si@C⊆Ni), ∼489 Ω (for Si@C), and ∼642 Ω (for pure Si). This comparison highly indicates the preferable chargetransfer properties for our configured systems. To account for the excellent anodic properties in prolonged lifespan, good cyclic stability, and large reversible capacity of Si@C⊆Ni, we 24161

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

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ACS Applied Materials & Interfaces

function as efficient current collectors. The rationally designed Si@C⊆Ni hybrid configurations endow Si nanoactives with attractive electrochemical performances including superior reversible capacity retention and cyclic stability (nearly no capacity decay appears within a total of 600 cycles with a large reversible capacity) together with excellent rate behaviors. Full cells of (−)Si@C⊆Ni//LiFePO4(+) are further constructed, which are capable of exhibiting great specific energy/powder densities, outstanding rate capability, and cyclic endurance as well. This research aims to paving the way for the development of reliable and applicable electrode configurations proper to nano-Si anodes. The simplicity and scalability of our synthesis strategy would make Si species promising for practical applications in future LIBs.

the outer double-shell architecture evolves into gel-like thin shells after a long-term cyclic period, and most of the Si nanocores turn amorphous and vary into several divergent fragments. However, a spherelike hybrid constitution is kept still, and the overall shelly films remain intact in the absence of evident splits/breaks. To confirm their great potential in power-supply usage, Si@ C⊆Ni hybrid anodes are paired with LiFePO4 cathodes (theoretical capacity: ∼170 mA h g−1) to constitute (−)Si@ C⊆Ni//LiFePO4(+) full-cell LIBs. Figure 4a shows the CV plots of the as-assembled full cells at a scan rate of 0.5 mV s−1 in a voltage window of ∼2.5−4.2 V (the inset schematic showing the full-cell constitution). The first charge process includes three redox peaks at the positions of ∼3.16, ∼3.64, and ∼3.87 V. The former two peaks disappearing in subsequent CV cycles are greatly associated with electrolyte oxidations and SEI formation, whereas the last one stems from Li+ (deintercalated from LiFePO4 crystals) insertion into Si actives driven by an external electric field. In the reverse discharge stage, a broad peak ranging from ∼2.88 to 3.69 V (peak summit: ∼3.3 V) corresponds to Li+ extraction from LixSi and return to LiFePO4 crystals. The following CV plots are nearly overlapped with each other, revealing the excellent operation stability of (−)Si@C⊆Ni//LiFePO4(+). The rate capability is furthermore estimated under varied current densities from ∼0.15 to ∼4.8 A g−1 (Figure 4b). Such cells can output reversible discharge capacities (calculated based on the total mass of anode and cathode) as large as ∼169.8 (0.15 A g−1), ∼147.4 (0.3 A g−1), ∼130.7 (0.6 A g−1), ∼114.6 (1.2 A g−1), ∼94.8 (2.4 A g−1), and ∼69.8 mA h g−1 (4.8 A g−1). Even when the current density jumps immediately from 4.8 back to the pristine 0.15 A g−1, a max. specific capacity of ∼168.5 mA h g−1 is achieved (capacity recovery ratio close to ∼100%), justifying their great cyclic reversibility for practical applications. To check whether the reversible Li storage properties are changed in continual cycling, comparisons of CD profiles between the 2nd and 90th cycles are carried out (see the inset in Figure 4b). Except for a slight capacity decrease, no remarkable variations in cell operation parameters (e.g., CE, voltage profiles/plateaus, etc.) are noted. To highlight the superiority of (−)Si@C⊆Ni//LiFePO4(+) in energy supply, the relationship of energy density versus power density is plotted (see the Ragone plot in Figure 4c). At a low power density of ∼64.65 W kg−1, the configured cells can deliver a max. energy density as large as ∼423 W h kg−1, much higher than the values of commercial LIBs or even comparable to those of Li−S batteries.6,9,10,12,20,21 Even if running at a power density up to ∼497.8 W kg−1, the cell energy density can stay at a high level of ∼197.8 W h kg−1. Optical images in Figure 4d reveal that a high-power blue light-emitting diode light (operating voltage: >2.5 V and rated power: 3−5 W) is easily driven by an individual fully charged (−)Si@C⊆Ni//LiFePO4(+) with a long working duration time over 5 min, exhibiting its great potential in portable energy supply applications.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b07737.

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Raman spectrum, EIS and TG plots, SEM/TEM observations, and cyclic CE curve for distinct electrode samples (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. Phone: +86-23-68254842. Fax: +86-23-68254969 (J.J.). *E-mail: [email protected] (J.L.). *E-mail: [email protected] (C.M.L.). ORCID

Jian Jiang: 0000-0001-7883-6631 Chang Ming Li: 0000-0002-4041-2574 Author Contributions ⊥

J.J. and H.Z. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge financial support from the Fundamental Research Funds for the Central Universities (XDJK2018C005, SWU115027, and SWU115029), National Natural Science Foundation of China (11604267), and Chongqing Natural Science Foundation (cstc2016jcyjA0477). This project is also supported by the Program for Innovation Team Building at Institutions of Higher Education in Chongqing (CXTDX201601011) and Graduate Student Research Innovation (CYS16048).

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REFERENCES

(1) Li, Y.; Yan, K.; Lee, H.-W.; Lu, Z.; Liu, N.; Cui, Y. Growth of Conformal Graphene Cages on Micrometre-Sized Silicon Particles as Stable Battery Anodes. Nat. Energy 2016, 1, 15029−15036. (2) Casimir, A.; Zhang, H.; Ogoke, O.; Amine, J. C.; Lu, J.; Wu, G. Silicon-Based Anodes for Lithium-Ion Batteries: Effectiveness of Materials Synthesis and Electrode Preparation. Nano Energy 2016, 27, 359−376. (3) Su, H.; Barragan, A. A.; Geng, L.; Long, D.; Ling, L.; Bozhilov, K. N.; Mangolini, L.; Guo, J. Colloidal Synthesis of Silicon-Carbon Composite Material for Lithium-Ion Batteries. Angew. Chem. 2017, 56, 10780−10785.

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CONCLUSIONS In summary, we have demonstrated a highly feasible engineering protocol to push forward Si anode performances for advanced LIBs. Y−S Si@C nanoparticles are trapped in a durable and ultrathin Ni matrix, a tough and malleable nanoarmor that can restrain volume expansions/negative aggregation effects of Si and cooperate with C layers to 24162

DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163

Research Article

ACS Applied Materials & Interfaces (4) Peng, H.-J.; Huang, J.-Q.; Cheng, X.-B.; Zhang, Q. Review on High-Loading and High-Energy Lithium-Sulfur Batteries. Adv. Energy Mater. 2017, 7, 1700260−1700313. (5) Jin, Y.; Zhu, B.; Lu, Z.; Liu, N.; Zhu, J. Challenges and Recent Progress in the Development of Si Anodes for Lithium-Ion Battery. Adv. Energy Mater. 2017, 7, 1700715−1700731. (6) Zhang, L.; Rajagopalan, R.; Guo, H.; Hu, X.; Dou, S.; Liu, H. Lithium-Ion Batteries: A Green and Facile Way to Prepare GranadillaLike Silicon-Based Anode Materials for Li-Ion Batteries. Adv. Funct. Mater. 2016, 26, 440−446. (7) Liu, N.; Lu, Z.; Zhao, J.; McDowell, M. T.; Lee, H.-W.; Zhao, W.; Cui, Y. A Pomegranate-Inspired Nanoscale Design for LargeVolume-Change Lithium Battery Anodes. Nat. Nanotechnol. 2014, 9, 187−192. (8) Jin, Y.; Li, S.; Kushima, A.; Zheng, X.; Sun, Y.; Xie, J.; Sun, J.; Xue, W.; Zhou, G.; Wu, J.; Shi, F.; Zhang, R.; Zhu, Z.; So, K.; Cui, Y.; Li, J. Self-Healing SEI Enables Full-Cell Cycling of A Silicon-Majority Anode with A Coulombic Efficiency Exceeding 99.9%. Energy Environ. Sci. 2017, 10, 580−592. (9) Lin, D.; Lu, Z.; Hsu, P.-C.; Lee, H. R.; Liu, N.; Zhao, J.; Wang, H.; Liu, C.; Cui, Y. A high tap density secondary silicon particle anode fabricated by scalable mechanical pressing for lithium-ion batteries. Energy Environ. Sci. 2015, 8, 2371−2376. (10) Liu, X.; Huang, J.-Q.; Zhang, Q.; Mai, L. Nanostructured Metal Oxides and Sulfides for Lithium-Sulfur Batteries. Adv. Mater. 2017, 29, 1601759−1601783. (11) Wu, H.; Chan, G.; Choi, J. W.; Ryu, I.; Yao, Y.; McDowell, M. T.; Lee, S. W.; Jackson, A.; Yang, Y.; Hu, L.; Cui, Y. Stable Cycling of Double-Walled Silicon Nanotube Battery Anodes through SolidElectrolyte Interphase Control. Nat. Nanotechnol. 2012, 7, 310−315. (12) Zhou, L.; Zhuang, Z.; Zhao, H.; Lin, M.; Zhao, D.; Mai, L. Intricate Hollow Structures: Controlled Synthesis and Applications in Energy Storage and Conversion. Adv. Mater. 2017, 29, 1602914− 1602942. (13) Wu, H.; Yu, G.; Pan, L.; Liu, N.; Mcdowell, M. T.; Bao, Z.; Cui, Y. Stable Li-Ion Battery Anodes by In-Situ Polymerization of Conducting Hydrogel to Conformally Coat Silicon Nanoparticles. Nat. Commun. 2013, 4, 1943−1948. (14) Zuo, X.; Zhu, J.; Müller-Buschbaum, P.; Cheng, Y.-J. Silicon Based Lithium-Ion Battery Anodes: A Chronicle Perspective Review. Nano Energy 2017, 31, 113−143. (15) Zheng, G.; Lee, S. W.; Liang, Z.; Lee, H.-W.; Yan, K.; Yao, H.; Wang, H.; Li, W.; Chu, S.; Cui, Y. Interconnected Hollow Carbon Nanospheres for Stable Lithium Metal Anodes. Nat. Nanotechnol. 2014, 9, 618−623. (16) Qian, J.; Henderson, W. A.; Xu, W.; Bhattacharya, P.; Engelhard, M.; Borodin, O.; Zhang, J.-G. High Rate and Stable Cycling of Lithium Metal Anode. Nat. Commun. 2015, 6, 6362−6370. (17) Liu, N.; Wu, H.; McDowell, M. T.; Yao, Y.; Wang, C.; Cui, Y. A Yolk-Shell Design for Stabilized and Scalable Li-Ion Battery Alloy Anodes. Nano Lett. 2012, 12, 3315−3321. (18) Wei, Q.; Xiong, F.; Tan, S.; Huang, L.; Lan, E. H.; Dunn, B.; Mai, L. Porous One-Dimensional Nanomaterials: Design, Fabrication and Applications in Electrochemical Energy Storage. Adv. Mater. 2017, 29, 1602300−1602338. (19) Zhou, X.; Yin, Y.-X.; Wan, L.-J.; Guo, Y.-G. Self-Assembled Nanocomposite of Silicon Nanoparticles Encapsulated in Graphene through Electrostatic Attraction for Lithium-Ion Batteries. Adv. Energy Mater. 2012, 2, 1086−1090. (20) Zhang, L.; Rajagopalan, R.; Guo, H.; Hu, X.; Dou, S.; Liu, H. A Green and Facile Way to Prepare Granadilla-Like Silicon-Based Anode Materials for Li-Ion Batteries. Adv. Funct. Mater. 2016, 26, 440−446. (21) Bruce, P. G.; Freunberger, S. A.; Hardwick, L. J.; Tarascon, J.M. Li-O2 and Li-S batteries with high energy storage. Nat. Mater. 2012, 11, 19−29. (22) Huang, X.; Sui, X.; Yang, H.; Ren, R.; Wu, Y.; Guo, X.; Chen, J. HF-Free Synthesis of Si/C Yolk/Shell Anodes for Lithium-Ion Batteries. J. Mater. Chem. A. 2018, 6, 2593−2599.

(23) Su, X.; Wu, Q.; Li, J.; Xiao, X.; Lott, A.; Lu, W.; Sheldon, B. W.; Wu, J. Silicon-Based Nanomaterials for Lithium-Ion Batteries: A Review. Adv. Energy Mater. 2014, 4, 1300882−1300904. (24) Peled, E.; Patolsky, F.; Golodnitsky, D.; Freedman, K.; Davidi, G.; Schneier, D. Tissue-like Silicon Nanowires-Based Three-Dimensional Anodes for High-Capacity Lithium Ion Batteries. Nano Lett. 2015, 15, 3907−3916. (25) Sun, Y.; Liu, N.; Cui, Y. Promises and Challenges of Nanomaterials for Lithium-Based Rechargeable Batteries. Nat. Energy 2016, 1, 16071−16082. (26) Hou, S.; Xu, X.; Wang, M.; Xu, Y.; Lu, T.; Yao, Y.; Pan, L. Carbon-incorporated Janus-type Ni2P/Ni hollow spheres for high performance hybrid supercapacitors. J. Mater. Chem. A. 2017, 5, 19054−19061. (27) Guan, C.; Sumboja, A.; Wu, H.; Ren, W.; Liu, X.; Zhang, H.; Liu, Z.; Cheng, C.; Pennycook, S. J.; Wang, J. Hollow Co3 O4 Nanosphere Embedded in Carbon Arrays for Stable and Flexible Solid-State Zinc-Air Batteries. Adv. Mater. 2017, 29, 1704117− 1704125. (28) Guan, C.; Liu, X.; Elshahawy, A. M.; Zhang, H.; Wu, H.; Pennycook, S. J.; Wang, J. Metal-organic framework derived hollow CoS2 nanotube arrays: an efficient bifunctional electrocatalyst for overall water splitting. Nanoscale Horiz. 2017, 2, 342−348. (29) Guan, C.; Liu, X.; Ren, W.; Li, X.; Cheng, C.; Wang, J. Rational Design of Metal-Organic Framework Derived Hollow NiCo2 O4 Arrays for Flexible Supercapacitor and Electrocatalysis. Adv. Energy Mater. 2017, 7, 1602391−1602398. (30) Yang, L. Y.; Li, H. Z.; Liu, J.; Sun, Z. Q.; Tang, S. S.; Lei, M. Dual Yolk-Shell Structure of Carbon and Silica-Coated Silicon for High-Performance Lithium-Ion Batteries. Sci. Rep. 2015, 5, 10908− 10916. (31) Xie, J.; Tong, L.; Su, L.; Xu, Y.; Wang, L.; Wang, Y. Core-Shell Yolk-Shell Si@C@Void@C Nanohybrids as Advanced Lithium Ion Battery Anodes with Good Electronic Conductivity and Corrosion Resistance. J. Power Sources 2017, 342, 529−536. (32) Tü y sü z , H.; Liu, Y.; Weidenthaler, C.; Schü t h, F. Pseudomorphic Transformation of Highly Ordered Mesoporous Co3O4 to CoO via Reduction with Glycerol. J. Am. Chem. Soc. 2008, 130, 14108−14110.

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DOI: 10.1021/acsami.8b07737 ACS Appl. Mater. Interfaces 2018, 10, 24157−24163