Decorating in situ ultrasmall tin particles on crumpled N-doped graphene for lithium-ion batteries with a long life cycle

Journal of Power Sources 328 (2016) 482e491

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Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour

Decorating in situ ultrasmall tin particles on crumpled N-doped graphene for lithium-ion batteries with a long life cycle Lianjun Liu, Xingkang Huang, Xiaoru Guo, Shun Mao, Junhong Chen* University of Wisconsin-Milwaukee, Mechanical Engineering Department, Milwaukee, WI 53211, United States

h i g h l i g h t s

g r a p h i c a l a b s t r a c t


Ultrasmall Sn nanoparticles are decorated in situ on crumpled Ndoped graphene.
Sn@NG shows high dispersion of Sn and large pore volume.
The effects of Sn loading step, treatment temperature and N-doping are studied.
Sn@NG demonstrates excellent longterm cycle stability at high current rates.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 2 May 2016 Received in revised form 14 July 2016 Accepted 8 August 2016

The practical application of Sn, a promising anode material for lithium-ion batteries, is hindered primarily by its huge volume change (up to 260%) upon lithiation. To tackle this obstacle, here we report a facile one-pot method, i.e., pyrolysis of a mixture of GO, SnCl4, and cyanamide at elevated temperatures to create in situ a novel mesoporous structure of Sn@N-doped graphene (Sn@NG). In the constructed architecture, the ultrasmall Sn nanoparticles (2e3 nm) are uniformly embedded in the NG network while the crumpled NG provides good electronic conductivity, abundant defects, high surface area, and large mesopore volume. Due to the combination of these merits, Sn@NG exhibits extremely long-term cycling stability, even at high rates, retaining a capacity of 568 mAh g1 at 1 A g1 (90% retention) and 535 mAh g1 at 2 A g1 (91.6% retention) after 1000 and 900 cycles, respectively. This performance is superior to that of Sn@G (without N-doping) and Sn//NG prepared using a two-step process with large particle sizes (>30 nm) and uneven dispersion of Sn. The findings from this work will shed light on the design of efficient and stable Sn and other metal-based materials for energy storage and conversion. © 2016 Elsevier B.V. All rights reserved.

Keywords: Sn N-doped graphene Mesoporous Nanoparticles Lithium storage

1. Introduction To apply rechargeable lithium-ion batteries (LIBs) in electric vehicles and renewable energy storage, alternative anode materials with high energy and power density and long life cycles are highly desired, since existing commercial graphite anodes have a

* Corresponding author. E-mail address: [email protected] (J. Chen). http://dx.doi.org/10.1016/j.jpowsour.2016.08.033 0378-7753/© 2016 Elsevier B.V. All rights reserved.

low capacity (372 mAh g1) and very limited energy output [1,2]. Alternatively, metallic tin (Sn) has been explored extensively as a promising anode candidate for LIBs because of its high theoretical capacity (993 mAh g1), good electronic conductivity, moderate operating voltage, and low cost [3e10]. However, the practical application of Sn anodes for LIBs is hindered by large volume change (up to 260%) during the lithium ion insertion/extraction into/from Sn [11e13], which causes severe pulverization problems, loss of contact with the current collector, and particle

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aggregation, and thus results in rapid capacity fading and poor cyclability. To tackle the obstacles presented by Sn anodes, in addition to forming SnM (M ¼ Co, Sb, Ni, Cu, etc.) alloys [7,14e17], two other effective strategies are (1) constructing small Sn particles (<10 nm) with a narrow particle size distribution to mitigate mechanical stress and retard particle pulverization [6,18,19], and (2) uniformly dispersing Sn nanoparticles (NPs) in a porous carbon matrix to accommodate the volume change [10,13,20e27]. Integrating both of these design principles to synthesize nanosized Sn decorated on porous carbon may provide advanced Sn-based anodes for highly reversible lithium storage. For example, Zhu et al. directly carbonized divalent Sn complex to synthesize ultrasmall Sn NPs (~5 nm) finely embedded in an N-doped porous carbon network, which delivered a capacity of 722 mAh g1 at a current density of 0.2 A g1 up to 200 cycles [19]. Until now, various nanostructured Sn/C materials have been explored extensively, such as Sn embedded in or grown on the porous carbon [4,12,28], hollow carbon spheres [29e31], graphene or carbon/graphene nanosheets [9,10,13,20,22,25,27,32e34], carbon nanofibers (CNFs) [35e37], and carbon nanotubes (CNTs) [3,5,38]. Among the previously described Sn/carbon anodes, Sn/graphene composites are of particular interest because of their outstanding electrical conductivity, superior mechanical flexibility, large surface area, and the high thermal/chemical stability of graphene [32,39]. However, the rate capability and long-term cycle stability at high current densities are not superior to or even worse than those of the other Sn/carbon. For example, flatted Sn sheets were confined in the space of pre-seeded graphene through multisteps (impregnation, filtration, annealing, melting, and quenching) [40], but the capacity of this “sandwich” layer structure decreased from 1000 mAh g1 to 650 mAh g1 after 100 cycles at 0.1 A g1 (65% retention). Another Sn/carbon/graphene 3D network, prepared by a combination of hydrothermal, freeze-drying, and thermal treatment methods, only delivered specific capacities of 600, 498, 440 and 344 mA h g1 at 0.1, 0.3, 0.5, and 1 A g1, respectively [27]. This inferior performance is likely attributed to (1) the irreversible aggregation or restacking of graphene sheets, (2) poorly controlled particle size and dispersion uniformity of Sn, and (3) the aggregation and unsatisfied confinement of Sn within graphene (partial Sn still decorated on the surface). Therefore, to achieve high electrochemical performance, it is highly desirable to develop a facile, general, and cost-effective approach to engineer Sn/ graphene-based composites with small particle size, uniform distribution, and porous network structure. In this work, we report for the first time the facile one-pot synthesis of ultrasmall Sn NPs (2e3 nm) finely decorated in crumpled N-doped graphene nanosheets (Sn@NG) by directly thermal treating the mixture of graphene oxide (GO), SnCl4 (Sn source), and cyanamide (N source). We hypothesized that during pyrolysis the N-doping effect will induce the transformation of GO to a porous graphene network with high surface area, large pore volume, and abundant defects, while Sn4þ cations are reduced in situ to Sn nanoparticles that are homogenously encapsulated on the folded graphene layers and the mesopores. In addition, partial N species in the precursor could be reserved in the graphene sheets, generating crumpled N-doped graphene that not only enhances the electrical conductivity of Sn, but also contributes to the capacity as an active component. This unique structure of Sn@NG has ultrasmall and uniform Sn particles to shorten the ion diffusion length, reduce the lattice strain, and provide rich space to buffer Sn volume expansion, thereby warranting excellent electrochemical performance for LIBs with a long lifetime.

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2. Experimental section 2.1. Material synthesis GO was prepared from natural graphite by a modified Hummers' method, as reported in our previous work [41]. Sn@NG was synthesized by a facile one-pot method. Briefly, 0.2 g GO, 0.6 g SnCl4, and 6 g cyanamide were dissolved in 100 ml deionized water by stirring for 1 h. Then, the solution was evaporated at 80  C to obtain the complex, which was ground in a mortar. The obtained powder was calcined at a desired temperature (700  C, 800  C, and 900  C) for 0.5 h in an Ar atmosphere. The resulting samples were denoted as Sn@NG700, Sn@NG800, and Sn@NG900. For comparison, NGsupported Sn particles (Sn//NG) were prepared using a two-step wet-impregnation method. In the first step, NG was synthesized in the same procedure as the Sn@NG900, except that the Sn precursor was added. In the second step, a desired amount of NG and SnCl4 was dissolved in an H2O/ethanol (1:1) solution. After evaporation at 80  C, the mixture was calcined at 700  C for 0.5 h to obtain the Sn//NG700. 2.2. Material characterization The crystal structures of the electrode materials were identified by X-ray diffraction (XRD, Bruker D8-Advance X-ray powder diffractometer). Raman spectroscopy was conducted with a Renishaw Raman spectrometer (Renishaw Inc., Wotton-underEdge, UK; Inc., 1000B) with a HeNe laser. The specific surface area and pore volume were analyzed by nitrogen adsorption-desorption at 77 K using the Brunauer-Emmett-Teller (BET) method (Micromeritics, ASAP 2020). The valence state of the C, O, N, and Sn elements were identified by X-ray photoelectron spectroscopy (XPS), VG ESCA 2000, with Mg Ka as the X-ray source, and the C1s peak at 284.6 eV as an internal standard. Scanning electron microscopy (SEM) (Hitachi S4800) was used to obtain the surface morphology. The surface dispersion of the C, O, N, and Sn elements was analyzed by X-ray elemental mapping. The lattice structure of Sn@NG was visualized by phase-contrast, high-resolution transmission electron microscopy (HRTEM) carried out with 300 keV electrons in a Hitachi H9000NAR instrument with 0.18 nm point and 0.11 nm lattice resolution. Amplitude contrast TEM images were used to obtain the sizes and morphology. The metallic Sn content of the electrode materials was measured on a thermal gravimetric analyzer (TGA-DAT-2660 SDT) at a heating rate of 10  C min1 from 25 to 800  C in air. 2.3. Electrochemical measurements For the electrode fabrication, the active material powder was mixed with carbon black and Poly(vinylydene fluoride) dissolved in n-methyl pyrolidinone, NMP, 8 wt%, in a weight ratio of 75:10:15. The slurry was mixed to obtain a homogeneous black paste, which was then coated onto copper foils. The as-coated copper foils were dried under vacuum at 90  C for 12 h. The mass loading of active materials on each electrode is about 1.06 mg/cm2. The working electrode and Li metal foil counter electrode were assembled into coin cells using a Celgard 2400 as the separator and a solution of 1 M LiPF6 in ethylene carbonate (EC)/diethyl carbonate (DEC) as the electrolyte. The cells were constructed in an Ar-filled glove box. The charge/discharge cycles were measured on a LAND CT2001A electrochemical workstation at a cutoff voltage of 0.01e2.5 V or 0.005e3.0 V under various current densities of 0.1 A g1, 0.2 A g1, 0.5 A g1, 1 A g1, 2 A g1, 3 A g1, and 5 A g1. The specific capacity was calculated based on the total mass of the active materials. CV measurements were conducted at 0.1 mV/s within 0e2.5 V on a

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CHI660D electrochemical workstation. Electrochemical impedance spectroscopy (EIS) measurements were performed using a CHI660D electrochemical workstation by employing an AC voltage of 5 mV amplitude in the frequency range 0.1e100 kHz. 3. Results and discussion 3.1. Preparation and morphology Fig. 1 schematically shows the procedure for synthesizing the Sn@NG, beginning with SnCl4, cyanamide (CA), and flat GO sheets. During the drying process, positively charged Sn4þ cations could be adsorbed on the negatively charged GO surface through the electrostatic interactions [42], while CA could cover the surface of GOlike thin films through either the chemical bonds (oxygen-containing groups in the GO reacting with amine groups in CA), hydrogen bond, or electrostatic interactions [22,41]. Such interactions may lead to the formation of a sandwich-like structure and the uniform dispersion of Sn4þ onto the GO sheets (Fig. S1). During the subsequent thermal treatment process in an Ar atmosphere, excessive CA could self-condense and completely decomþ þ pose into nitrogen-containing species (e.g., C2Nþ 2 , C3N2 , and C3N3 ) that provide nitrogen sources as the dopants [41]. When these plyometric species were removed from the surface, the GO shrank spontaneously to form the crumpled layers, leading to the formation of the mesoporous N-doped structure. Meanwhile, the decomposition of CA-Sn4þ-GO at moderate temperatures (e.g., 400  C) will generate SnO2 intermediates (Fig. S2), which are further reduced into metallic Sn via the reaction of SnO2 with carbon at high temperatures (700e900  C). The crumpled graphene layers could spontaneously wrap the Sn NPs produced in situ and thus generate the final composites (denoted as Sn@NG). As confirmed by the SEM images, Sn@NG700 (i.e., Sn@NG treated at 700  C) shows distinct crumpled networks and a porous structure, and it maintains the 2D geometry (Fig. 2a-b); however, no Sn-based particles are observed, even at a high magnification (Fig. 2c), suggesting that Sn particles are small-sized and highly dispersed. The elemental mapping images in Fig. 2d-h further demonstrate the good dispersion of Sn NPs and the successful doping of abundant N species into the graphene lattice. On the other hand, when synthesizing the hybrid in a two-step method, i.e., wet-impregnating the as-prepared NG by SnCl4 solution and thermal treatment at 700  C sequentially, the resulting Sn//NG700 only has larger particles (>30 nm) that were unevenly distributed on the surface of NG (Fig. S3). The digital photo in Fig. S4 further proves that Sn@NG700 is fluffy in a large volume while Sn//NG700 looks very dense, again highlighting the merits of the one-step method to prepare Sn@NG composites with a small size, good dispersion, and a porous structure. The difference in the two methods is likely because in the one-step process, the formation of mesopores and the confinement of Sn particles in situ occur

simultaneously by treating the complex of GO, SnCl4, and CA; in the two-step process, however, it is difficult to fully decorate Sn in the small mesopores of NG, as the majority of Sn particles tend to aggregate on its surface. Noticeably, when treated at higher temperatures (800  C and 900  C), the obtained Sn@NG800 and Sn@NG900 retained a layer morphology and porous structure very similar to Sn@NG700 (Fig. S5), but temperatures that were too high induced the sintering of partial Sn particles into big chunks over 1 mm (see the Sn elemental mapping image in Fig. S6). To clarify the N-doping effect on the morphology, we prepared Sn@G700 using the same procedure as for Sn@NG700 but without adding CA. Fig. 3 shows the SEM and elemental mapping images for Sn@G700. Clearly, the Sn@G700 has large, relatively flat and stacked sheets (more than 5 mm, Fig. 3a-b) on the surface or at the edge of which Sn nanoparticles aggregate to form big chunks (>100 nm, Fig. 3c). Some Sn particles are prone to separating from graphene (Fig. 3b right bottom corner and the mapping images). Obviously, the morphology of graphene and the dispersion and size of Sn particles are completely different from those of Sn@NG700. This comparison demonstrates the important role of N-doping, which improves the dispersion uniformity and size distribution of Sn particles and creates a porous graphene structure (as confirmed by N2 sorption described later). TEM and HRTEM were conducted to identify the morphology, particle size, dispersion, and lattice structure of Sn@NG700. The TEM images in Fig. 4a-c clearly show that Sn@NG700 has a crumpled structure, in which the wrinkles resulted from the crumpling of graphene sheets rather than the stacking of graphene. The highmagnification TEM image (Fig. 4d) shows that Sn NPs (black spots) with an average size of 3 nm are either uniformly dispersed on the surface of NG sheets or wrapped in the wrinkles. The HRTEM images (Fig. 4e-f) confirmed that the ultrasmall Sn particles embedded in the NG sheets have a narrow size distribution (2e3 nm). The sheets have an expanded lattice spacing of 0.393 nm around the edges, corresponding with the (002) planes of graphene, while the particles have a lattice spacing of 0.291 nm, consistent with the (200) plane of metallic Sn crystal [20,28,38]. The TEM/HRTEM analyses well support the above SEM observations. It should be noted that the particle size of Sn revealed by HRTEM observations (2e3 nm) is smaller than its average crystallite size (31.9 nm) calculated from the XRD result using Sherrer's Equation. One possible reason is that some Sn nanoparticles still tend to aggregate. The SEM and elemental mapping images in Fig. S7 show that some big particles with a size of 300e500 nm are embedded in NG, even though the majority of Sn nanocrystals are ultrafine nanoparticles. Since the XRD result gives us the average crystallite size of the overall bulk phase particles, the presence of some large particles (300e500 nm) could account for the larger average crystallite size of Sn from the XRD. In addition, metallic Sn nanocrystals often have good crystallinity that can lead to sharp diffraction peaks, even when the particle size of Sn is relatively

Fig. 1. Scheme for the synthesis procedure of Sn nanoparticles in situ decorated on N-doped graphene (Sn@NG).

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Fig. 2. (a-c), SEM images and (d-h) elemental mapping images for Sn@NG700.

Fig. 3. (a-c), SEM images and (d-h) elemental mapping for Sn@G700.

small. Similar phenomenon had also been observed in the literature that ultrasmall Sn particles (5 nm) embedded in N-doped carbon exhibited sharp diffraction peaks in XRD [19]. The present strategy for synthesizing this unique Sn NPdecorated, crumpled N-doped graphene nanostructure has several advantages. First, the synthetic procedure is quite simple and facile: a one-pot process by directly annealing the solid sources. Second, this method is highly effective and environmentally friendly: it takes less than 12 h to complete the entire process without any post-treatments (washing and drying). Third, ultrasmall Sn NPs and the porous structure of NG can be produced simultaneously in situ without using any surfactants, templates, expensive Sn precursor, organic solvents, or hydrogen, thereby holding great potential for large-scale production. Finally, abundant N can be uniformly doped into the lattice of the graphene sheets, which significantly provides more active sites and creates more mesopores in the graphene architecture (see the description in the next section).

3.2. Structure, textual property and chemical state The crystal structure, the quality of graphene, and the textural

properties of the Sn@NG samples were characterized by X-ray diffraction (XRD), Raman, and N2 sorption, respectively. As shown in Fig. 5a, Sn@NG700 and Sn@NG800 show very strong diffraction peaks, which are readily indexed to crystalline tin (JCPDS card No. 04-0673) [13,28,43]. In contrast, Sn@NG900 exhibits weaker peaks for Sn crystals and a broad peak at around 26 arising from the (002) plane of graphene sheets. This difference indicates that higher temperatures may facilitate the encapsulation of Sn nanoparticles by the crumpled graphene due to the greater exposure of defects/vacancies and the large pore volume. This is supported by the Raman spectra of the three samples (Fig. 5b), in which Sn@NG900 has a ratio of ID/IG that is nearly 100% higher than that of Sn@NG700, probably because the N-doping level is enhanced at higher temperatures [41]. Here, the D band at 1335 cm1 is related to edge defects, carbon disorders, and vacancies, while the G band at 1596 cm1 is attributed to SP2-carbon [13,27,32]. In addition, the ratio of ID/IG on Sn@G700 (1.13) is much lower than that on Sn@NG700 (1.76), indicating that N-doping could facilitate the creation of defects in graphene. In terms of the porosity, we measured the NG, Sn@G, and Sn@NG samples to understand the temperature and N-doping effects. Both the NG and Sn@NG materials displayed typical type-IV

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Fig. 4. (a-c) low-magnification TEM, (d) high-magnification TEM images, and (e, f) HRTEM images for Sn@NG700.

Fig. 5. (a) XRD patterns, (b) Raman spectra, (c) N2 adsorption-desorption isotherms, BET surface area, pore size, and pore volume, and (d) pore size distribution for NG and Sn@NG calcined at different temperatures, i.e., Sn@NG700, Sn@NG800, and Sn@NG900.

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isotherms with a type H3 hysteresis loop (Fig. 5c), suggesting the existence of slit-shaped mesopores caused by the space of the thin graphene sheets. The NG and Sn@NG also have similar pore size distributions, with a peak at 2.7 nm and a broad tail extending more than 50 nm (Fig. 5d). On the other hand, introducing Sn onto the NG resulted in a remarkable decrease of surface area and pore volume of the NG (inset Table in Fig. 5c). For example, NG has an ultrahigh surface area of 710 m2 g1 and pore volume of 2.22 cm3 g1 compared with 128 m2 g1 and 0.38 cm3 g1 for Sn@NG700. The changes in the textural properties of NG clearly demonstrated that Sn nanoparticles in situ were confined in the small mesopores (2.7 nm) of NG. Due to the confinement effect arising from the mesopores of NG, the sintering of Sn nanoparticles could be prevented, even when treated at 700  C, thus leading to the formation of the ultrasmall and highly dispersed Sn particles. It should be noted that Sn@NG800 and Sn@NG900 have a comparable, but slightly higher BET surface area and pore volume than Sn@NG700. Without N-doping, Sn@G700 shows a mesoporous structure and slightly lower BET surface area (105 m2 g1) (Fig. S8), but its mesopore size is located in the broad range of 3 nme7 nm and its pore volume (0.14 cm3 g1) is less than half that of the N-doped Sn@NG. The combination of SEM, Raman, and N2 sorption analyses suggest that higher temperatures and N-doping create more defects, higher surface area, and larger pore volume, but temperatures that are too high could cause the partial aggregation of Sn nanoparticles. X-ray photoelectron spectroscopy (XPS) was conducted to investigate the chemical states and surface composition of Sn@NG700. The atomic percentage of doped nitrogen is about 12.4 at.%, nearly twice as high as that reported in the literature (4 at.%-5 at.%) [19,20] and 50% higher than that obtained by a traditional ammonia-mediated CVD method (8 at.%) [44]. In addition, the weight percentage of Sn (calculated from XPS) is about 21.7 wt%, while the TGA shows the total mass loading of Sn is ~50 wt% calculated based on SnO2 (Fig. S9). This difference suggests the majority of Sn particles are covered by graphene layers. The above analyses again verify that our one-pot method is an effective strategy for developing a novel structure of Sn NPs wrapped with crumpled graphene with a high N-doping content. As shown in Fig. 6a, the C1s spectrum can be fitted with three components: the main peak at 284.6 eV is assigned to sp2 hybridized C atoms in graphene, while the other peak at 285.5 eV and the small shoulder 287.7 eV should be attributed to sp2 C atoms bonded to N and sp3 C bonded to oxygen-containing groups, respectively [12,13,22,41]. Similarly, in Fig. 6b, the high-resolution N 1s peaks at 398.4 eV, 399.7 eV, and 401.2 eV represent the pyridinic, pyrrolic, and graphitic types of N atoms doped into the graphene structure, respectively [22,41,45]. Although the XRD patterns prove that Sn@NG700 has crystalline Sn in the bulk phase, the strong Sn 3d5/2 (486.8 eV) and Sn 3d3/2 (495.4 eV) peaks in Fig. 6c suggest the surface of Sn@NG700 is mainly dominated by SnO2, likely because the sample is oxidized by air when exposed in an ambient environment. The O 1s peaks shown in Fig. 6d further confirm the presence of metal-bonded and adsorbed-oxygen-containing species on the surface. Noticeably, since GO cannot be fully reduced while XPS can only provide surface composition (less than 10 nm) rather than the bulk phase, the real ratio of Sn to SnO2 or the weight percentage of SnO2 is unclear. Our future work will focus on controlling the conditions to reduce SnO2/graphene and quantitatively studying the effect of Sn/SnO2 ratio on the electrochemical performance. 3.3. Electrochemical performance The electrochemical performance of the Sn@NG composites was measured with coin-type cells using lithium metal as both

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reference and counter electrodes (see the experimental section for details on the coin cell fabrication). All specific capacities reported in this paper were calculated on the basis of the total mass of Sn@NG. Fig. 7a compares the first discharge/charge voltage profiles of Sn@NG anodes calcined at different temperatures between 0.01 V and 2.5 V at a current density of 0.1 A g1. The three Sn@NG samples display similar voltage profiles, corresponding with the characteristics typical of a Sn electrode [8,23,38]. Noticeably, the Sn@NG700 shows the highest first discharge/charge capacities (2028 mAh g1/1054 mAh g1) corresponding with a Coulombic efficiency (CE) of 52.1% compared with Sn@NG800 (1619 mAh g1/ 786 mAh g1, CE ¼ 49.2%) and Sn@NG900 (1404 mAh g1/ 687 mAh g1, CE ¼ 49.2%). The low CE in the initial stage could be derived from the formation of solid-electrolyte interphase (SEI) film, the decomposition of electrolyte, and irreversible reduction of SnO2 [19,20,33]. Surprisingly, the initial charge capacity of Sn@NG700 (1054 mAh g1) is higher than the theoretical value of 860.5 mAh g1 (calculated by 1 727 mAh g  50% þ 993 mAh g1  50%). Such a phenomenon has also been observed in many other literature reported previously [46e52]. Although the exact mechanism of this high initial charge capacity is not very clear, the extra capacity should mainly result from the unique structure of Sn@NG itself. First, N-doping has induced the formation of more disorders (defects or vacancies), folded mesoporous structure with fewer layers, and expanded (002) interlayer spacing in NG, which can provide more lithium insertion active sites such as edge-type sites and nanopores [46,48e50]. For example, N-doped graphene analogous carbon particles exhibited a highly revisable capacity of 2132 mAh g1 at 0.1 A g1 and 785 mAh g1 at 5 A g1 after 1000 cycles [50]. On the other hand, there are some small SnO2 species that are strongly bonded with NG (see XPS analysis), which will be electrochemically active, since electrons can be rapidly transported to SnO2 via conducting NG networks. As is well known, SnO2 has a high theoretical capacity of 1494 mAh g1 if going through conversion reaction mechanism with lithium [52]. In this case, the reversible reaction between Liþ and SnO2 may enhance the overall capacity [51e53]. Zhou et al. once reported that SnO2 nanocrystals binding with Ndoped graphene sheets showed a high capacity of 1346 mAh g1 at 0.5 A g1 after 500 cycles, mainly due to the unchanged small size and enhanced lithium electrochemical activity of SnO2 nanocrystals [52]. In addition, the Sn@NG700 always displays superior rate capability over Sn@NG800, Sn@NG900, and Sn@G700. As shown in Fig. 7b, the specific reversible capacity of the samples decreases moderately when increasing the current densities from 0.1 to 5 A g1. At a low-current density of 0.2 A g1, the specific reversible capacity of Sn@NG700 reaches ~720 mAh g1, more than 100% higher than those of Sn@NG900 and Sn@G700 (~340 mAh g1). Even at a higher current density of 3 A g1, Sn@NG700 delivers a capacity of ~415 mAh g1. However, Sn@G700 has a capacity of only ~80 mAh g1 at 3 A g1. Besides, the capacity of Sn@NG700 could recover to the original values when switching to a lower current density again after high rate cycling. The outstanding rate performance of Sn@NG700 should benefit from the combination of good electrical conductivity offered by NG sheets, the short diffusion path for both electrons and ions provided by the ultrasmall Sn particles, and the crumpled mesoporous NG networks. Clearly, the novel structure of Sn@NG700 could preserve the integrity of the electrode and thus tolerate varying charge and discharge currents, which is very important for high-power applications of rechargeable batteries. For further comparison, we examined the long-term cycling performance (1000 cycles) of Sn@NG at a high current density (1 A g1), as shown in Fig. 7c. Clearly, Sn@NG700 exhibits a much

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Fig. 6. XPS spectra for Sn@NG700, (a) C 1s, (b) N 1s, (c) Sn 3d, and (d) O 1s.

higher capacity and better cycling stability than Sn//NG700 and Sn@G700, suggesting that the Sn loading step and N-doping are very important factors for governing the cycling performance. A high CE of >99% also was achieved after the initial several cycles, demonstrating a high reversibility of Sn@NG700. After 1000 cycles, Sn@NG700 still delivered a capacity of 568 mAh g1 (~90% retention efficiency), more than twice higher than that of Sn//NG700 (180 mAh g1, 30% retention efficiency) and Sn@G700 (150 mAh g1, 56% retention efficiency). This result is also superior to the traditional graphite anode (372 mAh g1), Sn/graphene and Sn/C composites reported previously (see Table S1). Since the mass loading of Sn is similar on Sn@NG700 and Sn// NG700, the much better cyclability of Sn@NG700 is attributed to the following merits: (1) The ultrasmall Sn NPs (2e3 nm) could significantly facilitate ion transport, minimize the strain generated during the lithiation/delithiation process, and then suppress the fracture of Sn NPs [18,19]; (2) Due to the high dispersion and close contact of Sn NPs with NG, the emerged stress could be evenly distributed in the whole composite, thereby preventing local cracking; and (3) The crumpled porous NG structure with small mesopores (3 nm) and large pore volume not only prevents Sn NPs from aggregation, but also provides sufficient space to allow for the volume expansion of Sn [10,27]. More interestingly, Fig. 7c shows that Sn@NG800 and Sn@NG900 have capacities lower than that for Sn//NG700 before 250 cycles, but demonstrate better cycling stability and finally deliver higher capacities than that for Sn//NG700. For example, the capacity of Sn@NG800 shows an upward trend that increases from the initial 370 mAh g1 to 413 mAh g1 after 1000 cycles at 1 A g1.

This outstanding cyclability again highlights the unique morphology and structure of Sn@NG prepared using the one-step method. There are two possible reasons for the increased capacity: (1) High-temperature treatment may confine more Sn NPs in the pores of the crumpled NG, which might not be fully activated during the initial charge/discharge process. As the cycling proceeded, the residual Sn NPs could further participate in an electrochemical reaction with Liþ ions. (2) Sn@NG800 and Sn@NG900 have a higher surface area and pore volume, and contain more vacancies and disorders, which could enhance extra lithium storage. A similar phenomenon in raising the capacity also was observed on the other Sn/CNTs, Sn/hollow carbon, and Sn/graphene composites reported in the literature [9,28,54]. To further understand the outstanding performance of Sn@NG700, its cyclic voltammograms (CVs) were conducted in the initial three cycles at a scan rate of 0.1 mV s1 between 0 and 2.5 V, as shown in Fig. 7d. During the first cathodic scan, the broad peaks from 0.8 to 0.2 V is ascribed to both the reduction of SnO2 (see XPS in Fig. 6) and the formation of SEI film from the decomposition of the electrolyte, thereby explaining the capacity loss during the first cycle [18,19]. In the anodic sweep, oxidation peaks between 0.5 and 0.9 V are assigned to the dealloying reaction of LixSn, i.e., from Li22Sn5 to Li7Sn3, LiSn, Li2Sn5 and Sn, respectively [28]. The broad peak at 1.2 V is ascribed to the conversion of LixC to C, confirming that NG participated in the reaction with Liþ and contributed to the capacity. All these peaks are reproducible and stable after the first scan, implying the good reversibility of the electrochemical reaction of the Sn@NG700 electrode. The impendence spectra of the three Sn@NG composites also were collected to explore the

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Fig. 7. (a) First charge-discharge curves of Sn@NG anodes at a current density of 0.1 A g1 ,(b) rate cycle performance of the Sn@NG and Sn@G composites at charge/discharge rates from 0.1 A g1 to 5 A g1 for 50 cycles, (c) long-term cycle performance of the Sn@NG, Sn//NG700, and Sn@G electrodes at current densities of 0.1 and 0.2 A g1 for the initial five cycles and then 1 A g1 for the next 1000 cycles; all the above performances are recorded within the range of 0.01e2.5 V, (d) cyclic voltammogram of Sn@NG700 composite collected at a scan rate of 0.2 mV within the voltage range of 0e2.5 V, and (e) EIS plots for the Sn@NG composites.

Fig. 8. Rate cycle performance of the Sn@NG700 composite at charge/discharge rates from 0.1 A g1 to 5 A g1 for 900 cycles within the range of 0.005e3.0 V.

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difference between them. As shown in Fig. 7e, the Nyquist plots for all the three samples displayed a single semicircle in the lowfrequency range, corresponding with the charge-transfer resistance (Rct) in the electrodes [55,56]. Sn@NG700 shows a semicircle diameter much smaller than that of Sn@NG800 and Sn@NG900, indicating that Sn@NG700 has a lower Rct than other two hybrids. Since Sn@NG700 has a lower surface area and pore volume than Sn@NG800 and Sn@NG800, the lower resistance of Sn@NG700 largely results from the smaller size and more homogeneous dispersion of Sn NPs. Considering NG as an active anode material, the LIBs performance of Sn@NG700 was tested with an extended voltage window of 0.005 Ve3.0 V Fig. 8 shows its rate cycle capability and long-term cycle performance at a high-current density. One can easily see that Sn@NG700 has a durable and stable rate capacity and long life cycle at different charge/discharge rates. In the first rate cycle, Sn@NG700 shows the average reversible capacities of 1,016, 840, 690, 548, 456, and 340 mAh g1 for 0.2, 0.5, 1, 2, 3, and 5 A g1, respectively. When the current rate returns from 5 A g1 to 0.2 A g1 after the second rate cycle, a capacity of 1041 mAh g1 is still recoverable without any losses. More impressively, when switching to a high-current rate of 2 A g1 again, the capacity can still reach as high as 584 mAh g1 and retain about 91.6% (535 mAh g1), even after 900 cycles. In contrast, the Sn//NG700 prepared by a two-step method, with a larger particle size and uneven distribution of Sn, shows a poor rate capability and suffers from fast capacity fading at a current density of 2 A g1 (see Fig. S10). This specific capacity and excellent long-term cycling stability of Sn@NG700, achieved at such a high rate of 2 A g1, are superior to most previously reported Sn nanostructures, Sn/carbon, and Sn/graphene composites (Table S1). In fact, most of the Snbased LIBs anodes exhibit a cycle life of less than 200 cycles at a current density between 0.05 and 2 A g1, and their capacity retention is also modest (69%e86%). Only very few Sn-based LIB anodes reported recently, including 3D Sn-graphene and Sn/CNT/ carbon box [28,32], can endure hundreds of charge/discharge cycles at high rates with good capacity retention. But, unlike our simple one-pot method, these two hybrids were prepared by a relatively complex, template-assisted CVD method with multiple steps and post-treatment. 4. Conclusion Sn@NG composite with ultrasmall Sn NPs (2e3 nm) decorated on a crumpled N-doped graphene network was prepared by pyrolysis of GO, SnCl4, and cyanamide mixture and evaluated as an anode material for rechargeable LIBs. From a synthetic perspective, this one-pot method is very attractive, because Sn@NG can be constructed in a straightforward, surfactant-free, and template-free process without any post-treatment. More importantly, the formation of Sn NPs, N-doping, and the creation of mesopores occur simultaneously. The treatment temperature and Sn loading step are very important factors governing the particle size and distribution of Sn NPs. Beyond the synthesis, benefiting from the combination of ultrasmall Sn NPs, uniform dispersion, high-level N doping (12 at.%), and mesoporous NG structure, Sn@NG delivered high specific capacity (1042 mAh g1 at 0.2 A g1), good rate capability, and long cycling stability. Even at high current rates of 1 A g1 and 2 A g1, reversible capacities of 568 mAh g1 and 535 mAh g1 were maintained after 1000 and 900 cycles, respectively. Acknowledgements Financial support for this work was provided by the University of Wisconsin System Applied Research Grant.

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