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Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries
Accepted Manuscript Double conductivity-improved porous Sn/Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries Qiang Liu, Jiajia Ye, Zizhong Chen, Qin Hao, Caixia Xu, Jiagang Hou PII: DOI: Reference:
S0021-9797(18)31375-4 https://doi.org/10.1016/j.jcis.2018.11.060 YJCIS 24325
To appear in:
Journal of Colloid and Interface Science
Received Date: Revised Date: Accepted Date:
7 October 2018 14 November 2018 15 November 2018
Please cite this article as: Q. Liu, J. Ye, Z. Chen, Q. Hao, C. Xu, J. Hou, Double conductivity-improved porous Sn/ Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries, Journal of Colloid and Interface Science (2018), doi: https://doi.org/10.1016/j.jcis.2018.11.060
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Double conductivity-improved porous Sn/Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries Qiang Liua, Jiajia Yea, Zizhong Chena, Qin Haoa, Caixia Xua*, and Jiagang Houb* a Institute for Advanced Interdisciplinary Research, University of Jinan, Jinan 250022, Shandong Province, China b Qilu University of Technology (Shandong Academy of Sciences), Jinan 250353, Shandong Province, China Fax: +86-531-82767046; Tel: +86-531-82767046 E-mail: [email protected]; [email protected] Abstract Carbon encapsulated porous Sn/Sn4P3 (Sn/Sn4P3@C) composite is conveniently prepared by one-step electrochemical dealloying of Sn80P20 alloy in mild conditions followed by growing one carbon layer. Controllable dealloying of the Sn80P20 alloy results in the formation of bicontinuous spongy Sn4P3 nanostructure with a part of residued metallic Sn atoms embedded in the porous skeleton. A uniform carbon layer is deposited on the nanoporous Sn/Sn4P3 to prevent the nanostructure’s pulverizing and agglomerating during lithium ion insertion/extraction. Upon double conductivity modification from metallic Sn matrix and carbon layer, the as-made composite displays superior lithium-storage performances with much higher specific capacity as well as better cycling stability compared with pure porous Sn4P3. It offers a specific capacity of 837 mA h g-1 after 100 cycles at a rate of 100 mA g-1. Even after 700 cycles at the higher 1
rate of 1000 mA g -1, the specific capacity still maintains as high as 589 mA h g-1. The Sn/Sn4P3 @C material possesses promising application potential as an alternative anode in the lithium storage fields. Keywords: porous, Sn4P3, lithium-ion batteries, anode, electrochemical dealloying 1. Introduction Lithium-ion batteries (LIBs) represent the most attractive commercialized device of rechargeable energy storage concerning the high energy density, long-term and stable cycling life, and environment harmless operation [1-6]. At present, the mainly used carbon anode materials in LIBs cannot meet the ever-growing demands of power and energy output for electric cars and smart grid due to their lower specific capacities (372 mA h g-1 for graphite). Consequently, it is stimulated to develop potential anode materials with higher lithium-storage performances. Transition metal phosphides have attracted great attentions owing to excellent electrochemical storage performances, which make them quitely absorbing as anode materials for promising applications in the energy storage fields[7,8], such as CoP [9], MoP [10], Sn4P3 [11], Cu3P [12], and FeP2 [13]. In addition, metal phosphide can cut down the risk of lithium dendrite as well as enhance the inherent security at the time of overcharge process because the main potential voltage of the metal phosphide is much larger than that of the lithium deposition voltage. For example, Wu et al. fabricated a yolk-shelled anode material composed of embedded Ni2P nanoparticles and interconnected porous graphene shell through an assembly and self-template strategy, which showed much enhanced cycling stability and rate capability during lithium ion 2
charge-discharge processes [14]. Yang et al. synthesized cobalt phosphide nanowires anchored on reduced graphene oxide through a hydrothermal method followed by a subsequent annealing step, which exhibited a high reversible specific capacity as well as excellent rate capability for lithium storage [15]. Among various metal phosphide anodes, tin phosphide (Sn4P3) holds attractive potential as one of the alternative electrode materials for lithium storage on account of its superiorities of higher specific capacity, better cycling stability, as well as lower cost [16]. Sn4P3 possesses high theoretical specific capacity (1255 mA h g -1), which is almost 3 times of that of graphite material [17]. Nevertheless, for Sn4P3 the relatively low electrical conductivity and the particle pulverization from the large volume expansion during lithium ion insertion/extraction also generates the irreversible structure collapse and damage, which may lead to the repeated construction of unstable solid electrolyte interface (SEI) and thus the low cycling stability. These drawbacks restrict the practical application of Sn4P3 as an anode in LIBs. In order to improve the conductivity of Sn4P3 as well as alleviate its volume expansion, a number of strategies have been developed, including the fabrication of various nanoscaled structure and the mixing with conductive materials [18,19]. The nano-sized composites can efficiently restrain the volume change and contraction by buffering the strain induced by lithium ion insertion/deinsertion reactions as well as accelerate the entrance of ions to the inside of the electrode matrix by furnishing shorter ion transfer distances. For instance, Liu et al. synthesized uniform yolk-shell Sn4P3@C nanospheres through a top-down phosphorized method, which showed outstanding rate 3
capability in Na-ion batteries compared with bulk materials [20]. Liu et al. synthesized Sn4P3 nanoparticles through solvothermal method at 180oC for 10 h, which showed high reversible capacity, remarkable rate capability, and excellent cycling stability in LIBs [21]. The great advances in the development of Sn4P3 anode in rechargeable batteries have encouraged the further exploration of high performance Sn4P3-based electrode material. It should be noticed that among a variety of nanostructures, the porous structure has been regarded as the preferable architecture for Li+ insertion/extraction due to the unique combination of continuous conductive network and abundant pore voids, which is favourable for providing interpenetrated ion transport pathway, buffering the volume expansion as well as alleviating the structural strain compared with particle-type anodes [18,22-24]. Consequently, the design of porous Sn4P3 materials has been fascinated as another effective method to improve their energy storage performances. In previous studies, sol-gel method [25], chemical vapor deposition [26], and solvothermal methods [27] etc. have been employed to prepare porous materials. It should be mentioned that dealloying technique has been confirmed to be a well scalable as well as widely applicable method for the fabrication of porous materials by rationally designing the composition of source alloys. Hao et al. fabricated three-dimensional nanoporous Si/Cu composite by the way of alloy refining combined with facile dealloying of SiCuAl alloy at room temperature [28]. In the present work, we aimed to fabricate nanoporous (NP) Sn/Sn4P3 composite by dealloying of Sn80P20 alloy considering that the Sn is highly conductive and also has 4
high lithium storage capacity itself. In the following, one uniform protective carbon layer was coated on the porous Sn/Sn4P3 surface. The double modification of conductivity for Sn4P3 is straightforwardly achieved based on the metallic Sn in the porous architecture as well the encapsulated carbon layer. In addition, the rich porosity in the composite and protective carbon layer is advantageous to buffer the volume change of the anode during repeated charge and discharge cycles. NP-Sn/Sn4P3@C anode delivers greatly enhanced reversible capacity compared with pure NP-Sn4P3 electrode. Moreover, NP-Sn/Sn4P3@C composite performs outstanding rate capability and cycling stability as a promising alternative anode material for LIBs. 2. Experimental 2.1 Sample preparation For a typical preparation of target anode, the commercial Sn80P20 (at.%) alloy was used as the precursor source. Sn80P20 alloy ribbons with the thickness of about 50 µm are fabricated by utilizing a single roller melt spinning instrument under a protective argon atmosphere. NP-Sn/Sn4P3 composite is fabricated by etching Sn80P20 alloy foils in 0.1 M H2SO4 solution for 1200 s under the applied potential of -0.15 V with the Ag/AgCl electrode as the reference electrode and a platinum plate as the counter electrode. Pure porous Sn4P3 was also prepared by further etching the electrochemically dealloyed sample above in 0.1 M HCl solution for 10 h. The NP-Sn/Sn4P3 sample was encapsulated with the dopamine hydrochloride by adding into the HCl/Trisbuffer solution (pH=8.5) containing a predesigned concentration of dopamine by polymerizing gently on the surface of Sn/Sn4P3 with the aid of vigorous 5
stirring for 24 h. The products were washed with deionized water (18.2 MΩ cm) by several times and dried in an oven at 60oC overnight. The dopamine hydrochloride-coated NP-Sn/Sn4P3 was acquired by repeatedly washing with ultra-pure water and then ethanol for several times to totally eliminate the unreacted dopamine monomer, followed by drying at 70oC for 10 h. Furthermore, the content of the carbon on the NP-Sn/Sn4P3 could be tailored by adjusting the dopamine concentration. Finally, the target NP-Sn/Sn4P3@C composite was achieved by sintering the sample above at 400oC for 2 h and then at 500oC for 3 h with a heating rate of 10oC min-1 in the nitrogen atmosphere [23,29,30]. 2.2 Sample characterization The crystal structures of the fabricated samples were identified by powder X-ray diffraction (XRD) using Bruker D8 advanced X-ray diffractometer with Cu Kα radiation at a step rate of 0.04o s-1 in the scan range of 20 to 90o (2θ). The morphology of the product was analysed by using scanning electron microscope (SEM) (JEOL JSM-6700) equipped with an Oxford INCA X-sight energy dispersive X-ray spectrometer for the compositional analysis. The elemental mapping was acquired using a FEI QUANTA FEG250 scanning electron microscope equipped with an INCA Energy XMAX-50 X-ray spectroscopy analyzer. X-ray photoelectron spectroscopy (XPS) was used to analyse the surface electron states of the as-made composite on an ESCALAB 250 X-ray photoelectron spectrometer with the Monochromated Al Kα X-ray as the excitation source and C 1s (284.6 eV) as the reference line. Transmission electron microscopy (TEM) (JEM-2100) was employed to verify the nanostructure of 6
the samples. A Raman spectrometer (Raman, Horiba) using an excitation laser beam with a wavelength of 532 nm was employed to collect Raman spectrum. The pore size distribution of the NP-Sn/Sn4P3@C composite was measured with KuBo-X1000 instrument using the Brunauer-Emmett-Teller (BET) method. 2.3 Electrode preparation and electrochemical tests The electrochemical performance versus Li was gauged using coin-type half-cells (size: 2032). To make the working anode, the Sn/Sn4P3@C composite, acetylene black, and sodium alginate were mixed in a weight ratio of 7:2:1 in ultra-pure water. The slurry was milled for 2 h, and then coated onto a copper foil current collector and dried in vacuum at 70oC for 10 h. A celgard 2300 microporous membrane was used as the separator. A 1 mol/L LiPF6 mixed solution (in ethylene carbonate, dimethyl carbonate, ethylene methyl carbonate with a v/v of 1:1:1 and 3 wt.% vinylene carbonate.) as the electrolyte. The galvanostatic charge-discharge experiments were tested using a NEWARE BTS 5 V-5 mA computer-controlled galvanostat (Shenzhen, China) in the voltage window from 0.01 to 3.0 V at different current rates. The cyclic voltammetry tests were performed on CHI 760 E electrochemical workstation (Shanghai CH Instruments Co, China) with a scan rate of 0.1 mV s-1 in the range of 0-3.0 V. Electrochemistry impedance spectroscopy (EIS) was measured with an amplitude of 10 mV in the frequency range from 0.01 to 100 kHz by Princeton Applied Research spectrometer. 3.
Results and discussion
7
Fig. 1. XRD patterns of (a) pure Sn 4P3, Sn/Sn4P3, and the Sn/Sn 4P3@C samples. The standard patterns of Sn (JCPDS 65-2631) and Sn 4P3 (JCPDS 20-1294) are also included for comparison. XPS data of (b) Sn 3d, (c) P 2p, and (d) C 1s for the fresh Sn/Sn 4P3@C. Selectively etching the more reactive element from the source alloy has been a powerful and versatile method to produce the porous structures [31]. The commercially purchased Sn80P20 alloy is used as the source alloy. As is shown in Fig. S1a, according to the atomic ratio between the Sn and P based on the energy dispersive X-ray spectroscopy (EDS) analysis of the alloy foils, precursor alloy composition is around Sn80P20. To obtain the facile and controllable fabrication of the nanostructured Sn/Sn4P3 composite, H2SO4 solution was utilized as the electrolyte to electrochemically remove Sn atoms to avoid the loss of P atoms. Upon electrochemical dealloying, the resulting Sn and P component is Sn61P39 as shown in Fig. S1b, indicating that a part of Sn atoms 8
were residued. According to the atomic ratio between the Sn and P based on the EDS analysis, the content of metallic Sn in Sn/Sn4P3 composite is calculated around 40.9 at.%. XRD technique was utilized to examine the phase structure of the Sn80P20 source alloy and the dealloyed sample. From Fig. 1a, it is clear that the Sn80P20 alloy consists of pure Sn and Sn4P3 phase. Upon etching the Sn80P20 alloy, the peak intensity of pure Sn decreased. The diffraction peaks around 28.9, 31.5, and 45.7 o (2θ) can be readily identified as the Sn4P3 species. The diffraction peaks at 30.6, 32.0, and 44.9o (2θ) can be attributed to the metallic Sn phase. After coating the carbon layer, the phase structure of pure Sn and Sn4P3 was well maintained, indicating the good controllability for the target structure during the whole preparation process of Sn/Sn4P3@C sample. From XRD pattern, it is clear that there is only pure Sn4P3 existed, indicating that the Sn atoms after two-step dealloying were almost removed completely. XPS was employed to further characterize the chemical state of the carbon coated dealloyed sample. Fig. 1b-d presents the core level spectra of Sn 3d, P 2p, and C 1s for the target composite. It can be seen in Fig. 1b that the peaks of P 2p3/2 and P 2p1/2 located at 128.1 and 129.1 eV are in great coincide with the previously recorded data for the Sn-P bonds of Sn4P3 phase [32,33]. The peak at 133.7 eV can be assigned to the formation of P-O bonds [34]. Similarly, from Fig. 1c it can be seen that the 3d5/2 and 3d3/2 peaks for spin-orbital splitting photoelectrons of Sn situated at the binding energies of 487.1 eV and 495.3 eV, respectively. These values are all in agreement with the values of the Sn-P bonds of Sn4P3 phase [35]. The peaks at 484.8 eV and 493.4 eV 9
can be abscribed to Sn0, also indicating the existence of the metallic Sn phase [36]. As shown in Fig. 1d, the single peak of C 1s located at 284.7 eV, which is in accordance with the binding energy of carbon. This also proved that the carbon was successfully coated [37]. XPS was also used to analyze the initial chemical states of Sn and P elements for Sn80P20 alloy and Sn/Sn4P3 sample after electrochemical corrosion. As shown in Fig. S2, by careful comparison with the XPS data of Sn/Sn4P3@C, it can be found that the chemical states of Sn and P elements are well maintained during the whole electrochemical dealloying and carbonization process. This is consistent with the XRD observations for the phase structure. We adopted the simple method to measure the carbon content by weighting the mass change of the Sn/Sn4P3 sample before and after carbonization process. The mass percentage of carbon in the composite was estimated to be ~8.1 wt.%. As shown in Fig. S3, raman spectroscopy was further used to identify the carbon in the Sn/Sn4P3@C, The two peaks at 1355 cm-1 as well as 1588 cm-1 correspond to the disordered (D) band as well as the graphitic (G) band of carbon [38].
Fig. 2. Schematic diagram for the preparation process of Sn/Sn4P3@C composite. Fig. 2 depicts the preparation process of the porous core-shelled Sn/Sn4P3@C material. After the tin atoms are electrochemically dissolved from Sn80P20 in H2SO4 solution, the nanostructured Sn/Sn4P3 framework can be successfully produced. The thin poly-dopamine layer is deposited on the Sn/Sn4P3 composite surface by 10
polymerization as the carbon source and then followed by carbonization to obtain the final anode sample.
Fig. 3. (a, b) SEM images of the Sn/Sn4P3 sample after electrochemical corrosion. (c, d) SEM images of the Sn/Sn4P3@C after covering carbon layer. The resulted structure after dealloying the sample alloy was characterized by SEM. As shown in Fig. S4, the Sn80P20 precursor alloy foil has smooth and compact microstructure. It can be shown in Fig. 3a that the interconnected backbone composed of several large channels distributed on the surface after dealloying the Sn80P20 alloy in 0.1 M H2SO4 solution by electrochemical corrosion. High-magnification SEM image (Fig. 3b) illustrates that a great deal of pores penetrated through the integral nanoscaled skeleton with the typical ligament size about 300 nm. Based on observations above, it is clear that upon dealloying Sn80P20 alloy, the uniform nanoporous architecture can be successfully obtained with interconnected skeleton and hollow channels. As shown in 11
Fig. 3c, upon coating the carbon layer, the original porous structure was well preserved without the surfaces pores blocked. From Fig. 3d, it is clearly observed that the ligaments underwent certain coarsening to 450 nm due to the high temperature carbonization.
Fig. 4. Elemental mappings of C (b), P (c), and Sn (d) in the Sn/Sn 4P3@C sample matching to Fig. 4a. The elemental distributions of the dealloyed sample were investigated by element mapping. It can be illustrated in Fig. 4a-d that the product mostly contained the Sn, P, and C elements, which uniformly dispersed in the whole configuration, further demonstrating that the uniform nanocomposite of metal phosphide has been successfully enwrapped by one homogeneous carbon layer.
12
Fig. 5. (a, b, c) TEM images of the Sn/Sn4P3@C sample. The detailed microstructure of the Sn/Sn4P3@C sample was further characterized by TEM. It can be illustrated in Fig. 5a that the black framework indicated the formation of interconnected ligaments, while the inner white region suggested the existence of embedded pores. From the high magnification TEM image (Fig. 5b&c), it can be obviously observed that the carbon layer was uniformly deposited on the surface of porous Sn/Sn4P3 with the typical thickness around 10 nm. The high-resolution TEM image was not obtained because the strong electron beam led to the sample resolved under the state of high applied voltage. Based on the experimental observations above, 13
electrochemical corrosion of the commercial Sn80P20 alloy followed by carbonization can successfully generate the uniform carbon encapsulated porous Sn4P3 nanocomposite with metallic Sn well distributed in the porous architecture. On the basis of the BET results (Fig. S5), the pore voids of Sn/Sn4P3@C have the wider size distribution with the relatively small pores examined. It should be mentioned that by the intrinsic restriction of BET technique the large pores around several hundred nanometers cann’t be measured. Combining with the BET results as well as the SEM and TEM observations, it is conclusive that there are multiscale pore voids distributed in the Sn/Sn4P3@C sample, which can provide the sufficient space for the volume change during the ion insertion/extraction. The lithium storage performances of the porous Sn/Sn4P3@C were examined in half-coin cells configuration utilizing the lithium foil as both reference and counter electrodes. Fig. 6a illustrated the cycle voltamemetric (CV) curves of the as-made anode for the initial three cycles in the voltage range of 0.01 to 3.0 V (vs Li+/Li) at 0.1 mV s-1. In the initial cathodic process, the broad peak appeared at 1.66 V and a reduction peak appeared at 0.59 V, which was characterized by the lithium insertion process as well as the conversion of Sn4P3 into metallic Sn and the formation of Li3P (Eq. (1)). The reduction peak located at 0.32 V was regarded as the reversible lithium alloying reactions with tin (Eq. (2)) as well as the formation of a SEI layer. The two obvious reduction peaks at 0.59 and 0.32 V could be assigned to the lithium ion and electron transportation, thus resulting in the rapid and adequate disintegration of the active composite. The anodic peaks at 0.53 and 0.66 V in the initial charging process 14
could be attributed to the transformation of the Lix Sn alloy and Li3P into Sn4P3 (Eq. (3) and Eq. (4)). After the first cycle, it was clear that the cathodic peaks shift obviously, indicating the configurational change in the electrode material after the Li+ insertion in the initial cycle and the formation of SEI layer. However, the oxidation peaks had no obvious shift, implying the great reversibility owing to the protective carbon layer and porous architecture. On account of these points of discussion, the reaction mechanism of Sn4P3 during the initial discharge was shown to be as follows: Sn4P3 + 9Li+ + 9e-1 → 4Sn + 3Li3P (1) 2Sn + 5Li+ + 5e-1 → Li5Sn2 (2) During the first charging process: Li5Sn2 → 2Sn + 5Li+ + 5e-1 (3). Fig. 6b illustrated the initial three galvanostatic charge-discharge curves of Sn/Sn4P3 @C electrode at 100 mA g -1. The discharge capacities of Sn/Sn4P3@C in the first three cycles were as high as 1332.2, 1036.9, and 964.8 mA h g-1. The initially high discharge capacity of the anode could be attributed to the formation of SEI film as well as the large surface area of Sn/Sn4P3@C sample, which furnished abundant surface reaction sites. The extensive quantity of SEI films also consumed a part of Li ions, and gave rise to irreversible insertion/extraction Li+ in the following cycles. Furthermore, there were two plateau regions during the initial discharge. The former presumably resulted from the reaction between Li and P atoms, and the latter corresponded to the alloying reaction between Sn and Li. For comparison, Fig. S6 showed the initial three galvanostatic charge-discharge curves of Sn/Sn4P3@C electrode at 2000 mA g-1. The 15
trend of degradation was more serious than that of 100 mA g-1 due to the more severe polarization under the high rate. 4Sn + 3Li3P → Sn4P3 + 9Li+ + 9e-1 (4)
Fig. 6. (a) CV curves of Sn/Sn4P3@C. (b) The galvanostatic charge-discharge curves at 100 mA g-1 for the first three cycles. Cycling performances and CEs of Sn4P3 and Sn/Sn4P3 @C anodes at 100 (c) and 1000 mA g-1 (d), respectively. To further clarify the enhanced lithium storage performances for the Sn/Sn4P3@C anode, Fig. 6c illustrated the galvanostatic cycling performance of the as-made anode at 100 mA g-1 with the profile of the pure porous Sn4P3 attached for comparison. As demonstrated in Fig. 6c, the Sn/Sn4P3@C anode demonstrated an initial discharge capacity of 1332 mA h g-1 with the slight degradation in the following cycles. And then, the capacity of Sn/Sn4P3@C began to rise, which was possibly because that the carbon layer protected the porous Sn4P3 with the growth of a stable polymeric SEI film preventing
the
collapse
of the
framework 16
structure
during
lithium
ion
insertion/extraction. For the pure Sn4P3, the capacity decayed rapidly with a linear mode. The excellent electrochemical storage performance of Sn/Sn4P3@C composite in LIBs was also demonstrated by evaluating the capacity retention. The capacity retention of the Sn/Sn4P3@C nanocomposite was 62.8%, which was assigned to the volume expansion as well as the decomposition of electrolyte during the reaction between Sn/Sn4P3@C and Li+. In addition, the capacity retention of the Sn/Sn4P3@C anode was dramatically higher than that of the pure Sn4P3 (9.3%). The cycling performance of the Sn/Sn4P3@C material was also evaluated at high rate of 1000 mA g-1. Fig. 6d showed that the initial capacity of Sn/Sn4P3@C anode was 1261 mA h g-1, which was higher than that of pure Sn4P3 electrode (1209 mA h g-1). It was fascinating to find that the capacity of Sn/Sn4P3@C anode illustrated a slight rising tendency from the 100th cycle, which was normally observed for transition metal phosphide anodes. It is generally ascribed to the growth of a stable polymeric SEI film, which can prevent the collapse of the framework structure during lithium ion insertion/extraction [39]. Upon 700 cycles the Sn/Sn4P3 @C anode offered a high discharge capacity of 589 mA h g-1, which was much higher than the pure Sn4P3 (114 mA h g-1). The corresponding initial capacity retention of the Sn/Sn4P3@C nanocomposite (~46.7%) at the higher rate of 1000 mA g-1 was much higher than that of Sn4P3 (9.4%). It was noted that the CE of the Sn/Sn4P3@C composite suddenly increased to 97.1% in the 5th cycle and further sustained between 97 and 100% in the following several hundred cycles. Compared with the Sn4P3, the as-made Sn/Sn4P3@C anode showed much higher capacity as well the longer cycling life for LIBs. The higher 17
reversible capacity and superior cycling stability of the Sn/Sn4P3@C could be assigned to the introduction of the carbon layer and the metallic Sn matrix. As described in Fig. 7a, the reversible capacities around 701, 596, 594, 565, and 526 mA h g-1 were achieved at the current densities of 100, 300, 500, 1000, and 2000 mA g -1, and the specific capacity rebounded back to 698 mA h g-1 when the current density was reduced back to 100 mA g-1. Respectively, which were dramatically higher than that of the Sn4P3 electrode under the identical situation. In particular, the Sn/Sn4P3@C electrode demonstrated the good cycling stability and high specific capacity when the current density reached 2000 mA g-1, illustrating its superior high rate capability at the larger current density. Upon long term cycling, the capacity of pure Sn4P3 decreased to relatively inferior value, in contrast, the capacity of the Sn/Sn4P3@C anode entirely recovered. These observations validated that the Sn/Sn4P3@C electrode had the higher rate capability as well as remarkably improved discharge capacity. The EIS measurement for the two anodes was gauged by utilizing fresh cells to explore the possible contribution of doping Sn and covering carbon layer. As indicated by the Nyquist profiles in Fig. 7b, both samples displayed a depressed semicircle in the charge transfer impedance of cells, and a straight line in the low-frequency range indicating the mass transfer of Li+. It was detected that in the diameter of the semicircle for the two samples are obviously different. For clear comparison, the EIS results have been fitted to be a semicircle and a line as shown in Fig. 7b inset. The small diameter for Sn/Sn4P3 @C indicated that the charge transfer of the Sn4P3 became more facile owing to the adulteration of excellent conductive metallic Sn and carbon layer, thus resulting 18
in a superior electrochemical performance.
Fig. 7. (a) Rate performance of Sn4P3 and Sn/Sn4P3@C electrodes. (b) The nyquist plots of the fresh Sn4P3 and Sn/Sn4P3 @C samples. Inset is the fitting EIS. To further recommend the lithium storage performances of Sn/Sn4P3@C nanocomposite, we made a comparison of the Sn/Sn4P3@C anode with other Sn4P3-based anode materials in Table 1 [17,21,40]. It was noted that the Sn/Sn4P3@C electrode displayed the highest reversible capacity of 837 and 589 mA h g-1 at the current densities of 100 and 1000 mA g -1, respectively, even after 100 and 700 cycles. We would highlight that the as-prepared Sn/Sn4P3@C material has the excellent discharge capacity and remarkable stable cyclability among all of the formerly recorded Sn4P3-based anode materials. The excellent discharge capacity and the higher capacity retention of the Sn/Sn4P3@C demonstrated its potential application in the lithium storage fields. 19
Table 1 Comparison of the electrochemical performances of various Sn4P3 materials with this work. Anode
Current
Cycle number
Remaining
density
capacity
(mA g-1)
(mA h g-1)
Sn4P3 hollow
100
20
261
microspheres Sn4P3
Electrolytes
Ref
EC/DMC/DEC
[40]
=1:1:1 50
60
481
EC/DMC/DEC
[21]
=1:1:1 Sn4P3
100
100
416
EC/DMC/DEC
[21]
=1:1:1 Sn4P3
500
100
210
EC/DMC/DEC
[21]
=1:1:1 Sn4P3/graphite
100
100
651
EC/EMC
[17]
=3:7 Sn/Sn4P3@C
100
100
837
EC/DMC/EMC =1:1:1+3%(wt)VC
Sn/Sn4P3@C
1000
700
589
EC/DMC/EMC =1:1:1+3%(wt)VC
This work This work
EC: Ethylene carbonate; DMC: Dimethyl carbonate; EMC: Ethylene methyl carbonate; DC: Diethyl carbonate; VC: vinylene carbonate
4. Conclusions Porous Sn/Sn4P3 composite with protected carbon layer anchored on the network matrix was easily prepared by facile dealloying of Sn80P20 alloy under mild conditions followed by carbonization. Owing to the incorporation of well-conductive Sn and cladding of carbon layer, the Sn/Sn4P3@C nanocomposite certified much reinforced cycling stability and high rate capability, particularly at large current rates. With the superiorities of unique performances as well as scalable fabrication, the as-prepared Sn/Sn4P3 @C composite holds excellent application potential for lithium storage. Acknowledgements This work was supported by the National Natural Science Foundation of 20
China (51772133) and Shandong Province (ZR2017JL022), and China postdoctoral science foundation (2017M622117), and the program for Taishan Scholar of Shandong province (ts201712048). References [1] J.T. Jiang, K. Zhu, Y.Z. Fang, H.Z. Wang, K. Ye, J. Yan, G.L. Wang, K. Cheng, L.M. Zhou, D.X. Cao, Coralloidal carbon-encapsulated CoP nanoparticles generated on biomass carbon as a high-rate and stable electrode material for lithium-ion batteries, J. Colloid Interface Sci. 530 (2018) 579-585. [2] B.B. Wang, X. Zhang, X.J. Liu, G. Wang, H. Wang, J.T. Bai, Rational design of Fe3O4@C yolk-shell nanorods constituting a stable anode for high-performance Li/Na-ion batteries, J. Colloid Interface Sci. 528 (2018) 225-236. [3] Y. Yang, G.Y. Huang, H.Y. Sun, M. Ahmad, Q.Y. Mou, H.M. Zhang, Preparation and electrochemical properties of mesoporous NiCo 2O4 double-hemisphere used as anode for lithium-ion battery, J. Colloid Interface Sci. 529 (2018) 357-365. [4] J. Lee, Y. Wu, Z.B. Peng, Hetero-nanostructured materials for high-power lithium ion batteries, J. Colloid Interface Sci. 529 (2018) 505-519. [5] X.Z. Hu, Y. Jin, B. Zhu, Z, Liu, D.F. Xu, Y.D. Guan, M.Y. Sun, F.L. Liu, Tuning density of Si nanoparticles on graphene sheets in graphene-Si aerogels for stable lithium ion batteries, J. Colloid Interface Sci. 532 (2018) 738-745. [6] Y.Z. Yan, H.L. Tang, J.S. Li, F. Wu, T.B. Wu, R. Wang, D. Liu, M. Pan, Z.Z. 21
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surfaces
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platinum,
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and
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[39] Y.L. Xu, B. Peng, F.M. Mulder, A high-rate and ultrastable sodium ion anode based on a novel Sn4P3-P@Graphene nanocomposite, Adv. Energy Mater. (2017) 1701847. [40] S.L. Liu, H.Z. Zhang, L.Q. Xu, L. B. Ma, Synthesis of hollow spherical tin phosphides (Sn 4P3) and their high adsorptive and electrochemical performance, J. Cryst. Growth. 438 (2016) 31-37.
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Graphical Abstract:
Porous
core-shelled
Sn/Sn4P3@C
composite
is
easily
fabricated
by
electrochemically etching SnP alloy followed by covering the uniform carbon layer, which shows superior lithium-storage performances with higher specific capacity, much enhanced rate capability, as well as better cycling stability compared with pure porous Sn4P3. The Sn/Sn4P3 @C material holds promising application potential as an alternative anode in the lithium storage fields.
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S0021-9797(18)31375-4 https://doi.org/10.1016/j.jcis.2018.11.060 YJCIS 24325
To appear in:
Journal of Colloid and Interface Science
Received Date: Revised Date: Accepted Date:
7 October 2018 14 November 2018 15 November 2018
Please cite this article as: Q. Liu, J. Ye, Z. Chen, Q. Hao, C. Xu, J. Hou, Double conductivity-improved porous Sn/ Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries, Journal of Colloid and Interface Science (2018), doi: https://doi.org/10.1016/j.jcis.2018.11.060
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Double conductivity-improved porous Sn/Sn4P3@carbon nanocomposite as high performance anode in Lithium-ion batteries Qiang Liua, Jiajia Yea, Zizhong Chena, Qin Haoa, Caixia Xua*, and Jiagang Houb* a Institute for Advanced Interdisciplinary Research, University of Jinan, Jinan 250022, Shandong Province, China b Qilu University of Technology (Shandong Academy of Sciences), Jinan 250353, Shandong Province, China Fax: +86-531-82767046; Tel: +86-531-82767046 E-mail: [email protected]; [email protected] Abstract Carbon encapsulated porous Sn/Sn4P3 (Sn/Sn4P3@C) composite is conveniently prepared by one-step electrochemical dealloying of Sn80P20 alloy in mild conditions followed by growing one carbon layer. Controllable dealloying of the Sn80P20 alloy results in the formation of bicontinuous spongy Sn4P3 nanostructure with a part of residued metallic Sn atoms embedded in the porous skeleton. A uniform carbon layer is deposited on the nanoporous Sn/Sn4P3 to prevent the nanostructure’s pulverizing and agglomerating during lithium ion insertion/extraction. Upon double conductivity modification from metallic Sn matrix and carbon layer, the as-made composite displays superior lithium-storage performances with much higher specific capacity as well as better cycling stability compared with pure porous Sn4P3. It offers a specific capacity of 837 mA h g-1 after 100 cycles at a rate of 100 mA g-1. Even after 700 cycles at the higher 1
rate of 1000 mA g -1, the specific capacity still maintains as high as 589 mA h g-1. The Sn/Sn4P3 @C material possesses promising application potential as an alternative anode in the lithium storage fields. Keywords: porous, Sn4P3, lithium-ion batteries, anode, electrochemical dealloying 1. Introduction Lithium-ion batteries (LIBs) represent the most attractive commercialized device of rechargeable energy storage concerning the high energy density, long-term and stable cycling life, and environment harmless operation [1-6]. At present, the mainly used carbon anode materials in LIBs cannot meet the ever-growing demands of power and energy output for electric cars and smart grid due to their lower specific capacities (372 mA h g-1 for graphite). Consequently, it is stimulated to develop potential anode materials with higher lithium-storage performances. Transition metal phosphides have attracted great attentions owing to excellent electrochemical storage performances, which make them quitely absorbing as anode materials for promising applications in the energy storage fields[7,8], such as CoP [9], MoP [10], Sn4P3 [11], Cu3P [12], and FeP2 [13]. In addition, metal phosphide can cut down the risk of lithium dendrite as well as enhance the inherent security at the time of overcharge process because the main potential voltage of the metal phosphide is much larger than that of the lithium deposition voltage. For example, Wu et al. fabricated a yolk-shelled anode material composed of embedded Ni2P nanoparticles and interconnected porous graphene shell through an assembly and self-template strategy, which showed much enhanced cycling stability and rate capability during lithium ion 2
charge-discharge processes [14]. Yang et al. synthesized cobalt phosphide nanowires anchored on reduced graphene oxide through a hydrothermal method followed by a subsequent annealing step, which exhibited a high reversible specific capacity as well as excellent rate capability for lithium storage [15]. Among various metal phosphide anodes, tin phosphide (Sn4P3) holds attractive potential as one of the alternative electrode materials for lithium storage on account of its superiorities of higher specific capacity, better cycling stability, as well as lower cost [16]. Sn4P3 possesses high theoretical specific capacity (1255 mA h g -1), which is almost 3 times of that of graphite material [17]. Nevertheless, for Sn4P3 the relatively low electrical conductivity and the particle pulverization from the large volume expansion during lithium ion insertion/extraction also generates the irreversible structure collapse and damage, which may lead to the repeated construction of unstable solid electrolyte interface (SEI) and thus the low cycling stability. These drawbacks restrict the practical application of Sn4P3 as an anode in LIBs. In order to improve the conductivity of Sn4P3 as well as alleviate its volume expansion, a number of strategies have been developed, including the fabrication of various nanoscaled structure and the mixing with conductive materials [18,19]. The nano-sized composites can efficiently restrain the volume change and contraction by buffering the strain induced by lithium ion insertion/deinsertion reactions as well as accelerate the entrance of ions to the inside of the electrode matrix by furnishing shorter ion transfer distances. For instance, Liu et al. synthesized uniform yolk-shell Sn4P3@C nanospheres through a top-down phosphorized method, which showed outstanding rate 3
capability in Na-ion batteries compared with bulk materials [20]. Liu et al. synthesized Sn4P3 nanoparticles through solvothermal method at 180oC for 10 h, which showed high reversible capacity, remarkable rate capability, and excellent cycling stability in LIBs [21]. The great advances in the development of Sn4P3 anode in rechargeable batteries have encouraged the further exploration of high performance Sn4P3-based electrode material. It should be noticed that among a variety of nanostructures, the porous structure has been regarded as the preferable architecture for Li+ insertion/extraction due to the unique combination of continuous conductive network and abundant pore voids, which is favourable for providing interpenetrated ion transport pathway, buffering the volume expansion as well as alleviating the structural strain compared with particle-type anodes [18,22-24]. Consequently, the design of porous Sn4P3 materials has been fascinated as another effective method to improve their energy storage performances. In previous studies, sol-gel method [25], chemical vapor deposition [26], and solvothermal methods [27] etc. have been employed to prepare porous materials. It should be mentioned that dealloying technique has been confirmed to be a well scalable as well as widely applicable method for the fabrication of porous materials by rationally designing the composition of source alloys. Hao et al. fabricated three-dimensional nanoporous Si/Cu composite by the way of alloy refining combined with facile dealloying of SiCuAl alloy at room temperature [28]. In the present work, we aimed to fabricate nanoporous (NP) Sn/Sn4P3 composite by dealloying of Sn80P20 alloy considering that the Sn is highly conductive and also has 4
high lithium storage capacity itself. In the following, one uniform protective carbon layer was coated on the porous Sn/Sn4P3 surface. The double modification of conductivity for Sn4P3 is straightforwardly achieved based on the metallic Sn in the porous architecture as well the encapsulated carbon layer. In addition, the rich porosity in the composite and protective carbon layer is advantageous to buffer the volume change of the anode during repeated charge and discharge cycles. NP-Sn/Sn4P3@C anode delivers greatly enhanced reversible capacity compared with pure NP-Sn4P3 electrode. Moreover, NP-Sn/Sn4P3@C composite performs outstanding rate capability and cycling stability as a promising alternative anode material for LIBs. 2. Experimental 2.1 Sample preparation For a typical preparation of target anode, the commercial Sn80P20 (at.%) alloy was used as the precursor source. Sn80P20 alloy ribbons with the thickness of about 50 µm are fabricated by utilizing a single roller melt spinning instrument under a protective argon atmosphere. NP-Sn/Sn4P3 composite is fabricated by etching Sn80P20 alloy foils in 0.1 M H2SO4 solution for 1200 s under the applied potential of -0.15 V with the Ag/AgCl electrode as the reference electrode and a platinum plate as the counter electrode. Pure porous Sn4P3 was also prepared by further etching the electrochemically dealloyed sample above in 0.1 M HCl solution for 10 h. The NP-Sn/Sn4P3 sample was encapsulated with the dopamine hydrochloride by adding into the HCl/Trisbuffer solution (pH=8.5) containing a predesigned concentration of dopamine by polymerizing gently on the surface of Sn/Sn4P3 with the aid of vigorous 5
stirring for 24 h. The products were washed with deionized water (18.2 MΩ cm) by several times and dried in an oven at 60oC overnight. The dopamine hydrochloride-coated NP-Sn/Sn4P3 was acquired by repeatedly washing with ultra-pure water and then ethanol for several times to totally eliminate the unreacted dopamine monomer, followed by drying at 70oC for 10 h. Furthermore, the content of the carbon on the NP-Sn/Sn4P3 could be tailored by adjusting the dopamine concentration. Finally, the target NP-Sn/Sn4P3@C composite was achieved by sintering the sample above at 400oC for 2 h and then at 500oC for 3 h with a heating rate of 10oC min-1 in the nitrogen atmosphere [23,29,30]. 2.2 Sample characterization The crystal structures of the fabricated samples were identified by powder X-ray diffraction (XRD) using Bruker D8 advanced X-ray diffractometer with Cu Kα radiation at a step rate of 0.04o s-1 in the scan range of 20 to 90o (2θ). The morphology of the product was analysed by using scanning electron microscope (SEM) (JEOL JSM-6700) equipped with an Oxford INCA X-sight energy dispersive X-ray spectrometer for the compositional analysis. The elemental mapping was acquired using a FEI QUANTA FEG250 scanning electron microscope equipped with an INCA Energy XMAX-50 X-ray spectroscopy analyzer. X-ray photoelectron spectroscopy (XPS) was used to analyse the surface electron states of the as-made composite on an ESCALAB 250 X-ray photoelectron spectrometer with the Monochromated Al Kα X-ray as the excitation source and C 1s (284.6 eV) as the reference line. Transmission electron microscopy (TEM) (JEM-2100) was employed to verify the nanostructure of 6
the samples. A Raman spectrometer (Raman, Horiba) using an excitation laser beam with a wavelength of 532 nm was employed to collect Raman spectrum. The pore size distribution of the NP-Sn/Sn4P3@C composite was measured with KuBo-X1000 instrument using the Brunauer-Emmett-Teller (BET) method. 2.3 Electrode preparation and electrochemical tests The electrochemical performance versus Li was gauged using coin-type half-cells (size: 2032). To make the working anode, the Sn/Sn4P3@C composite, acetylene black, and sodium alginate were mixed in a weight ratio of 7:2:1 in ultra-pure water. The slurry was milled for 2 h, and then coated onto a copper foil current collector and dried in vacuum at 70oC for 10 h. A celgard 2300 microporous membrane was used as the separator. A 1 mol/L LiPF6 mixed solution (in ethylene carbonate, dimethyl carbonate, ethylene methyl carbonate with a v/v of 1:1:1 and 3 wt.% vinylene carbonate.) as the electrolyte. The galvanostatic charge-discharge experiments were tested using a NEWARE BTS 5 V-5 mA computer-controlled galvanostat (Shenzhen, China) in the voltage window from 0.01 to 3.0 V at different current rates. The cyclic voltammetry tests were performed on CHI 760 E electrochemical workstation (Shanghai CH Instruments Co, China) with a scan rate of 0.1 mV s-1 in the range of 0-3.0 V. Electrochemistry impedance spectroscopy (EIS) was measured with an amplitude of 10 mV in the frequency range from 0.01 to 100 kHz by Princeton Applied Research spectrometer. 3.
Results and discussion
7
Fig. 1. XRD patterns of (a) pure Sn 4P3, Sn/Sn4P3, and the Sn/Sn 4P3@C samples. The standard patterns of Sn (JCPDS 65-2631) and Sn 4P3 (JCPDS 20-1294) are also included for comparison. XPS data of (b) Sn 3d, (c) P 2p, and (d) C 1s for the fresh Sn/Sn 4P3@C. Selectively etching the more reactive element from the source alloy has been a powerful and versatile method to produce the porous structures [31]. The commercially purchased Sn80P20 alloy is used as the source alloy. As is shown in Fig. S1a, according to the atomic ratio between the Sn and P based on the energy dispersive X-ray spectroscopy (EDS) analysis of the alloy foils, precursor alloy composition is around Sn80P20. To obtain the facile and controllable fabrication of the nanostructured Sn/Sn4P3 composite, H2SO4 solution was utilized as the electrolyte to electrochemically remove Sn atoms to avoid the loss of P atoms. Upon electrochemical dealloying, the resulting Sn and P component is Sn61P39 as shown in Fig. S1b, indicating that a part of Sn atoms 8
were residued. According to the atomic ratio between the Sn and P based on the EDS analysis, the content of metallic Sn in Sn/Sn4P3 composite is calculated around 40.9 at.%. XRD technique was utilized to examine the phase structure of the Sn80P20 source alloy and the dealloyed sample. From Fig. 1a, it is clear that the Sn80P20 alloy consists of pure Sn and Sn4P3 phase. Upon etching the Sn80P20 alloy, the peak intensity of pure Sn decreased. The diffraction peaks around 28.9, 31.5, and 45.7 o (2θ) can be readily identified as the Sn4P3 species. The diffraction peaks at 30.6, 32.0, and 44.9o (2θ) can be attributed to the metallic Sn phase. After coating the carbon layer, the phase structure of pure Sn and Sn4P3 was well maintained, indicating the good controllability for the target structure during the whole preparation process of Sn/Sn4P3@C sample. From XRD pattern, it is clear that there is only pure Sn4P3 existed, indicating that the Sn atoms after two-step dealloying were almost removed completely. XPS was employed to further characterize the chemical state of the carbon coated dealloyed sample. Fig. 1b-d presents the core level spectra of Sn 3d, P 2p, and C 1s for the target composite. It can be seen in Fig. 1b that the peaks of P 2p3/2 and P 2p1/2 located at 128.1 and 129.1 eV are in great coincide with the previously recorded data for the Sn-P bonds of Sn4P3 phase [32,33]. The peak at 133.7 eV can be assigned to the formation of P-O bonds [34]. Similarly, from Fig. 1c it can be seen that the 3d5/2 and 3d3/2 peaks for spin-orbital splitting photoelectrons of Sn situated at the binding energies of 487.1 eV and 495.3 eV, respectively. These values are all in agreement with the values of the Sn-P bonds of Sn4P3 phase [35]. The peaks at 484.8 eV and 493.4 eV 9
can be abscribed to Sn0, also indicating the existence of the metallic Sn phase [36]. As shown in Fig. 1d, the single peak of C 1s located at 284.7 eV, which is in accordance with the binding energy of carbon. This also proved that the carbon was successfully coated [37]. XPS was also used to analyze the initial chemical states of Sn and P elements for Sn80P20 alloy and Sn/Sn4P3 sample after electrochemical corrosion. As shown in Fig. S2, by careful comparison with the XPS data of Sn/Sn4P3@C, it can be found that the chemical states of Sn and P elements are well maintained during the whole electrochemical dealloying and carbonization process. This is consistent with the XRD observations for the phase structure. We adopted the simple method to measure the carbon content by weighting the mass change of the Sn/Sn4P3 sample before and after carbonization process. The mass percentage of carbon in the composite was estimated to be ~8.1 wt.%. As shown in Fig. S3, raman spectroscopy was further used to identify the carbon in the Sn/Sn4P3@C, The two peaks at 1355 cm-1 as well as 1588 cm-1 correspond to the disordered (D) band as well as the graphitic (G) band of carbon [38].
Fig. 2. Schematic diagram for the preparation process of Sn/Sn4P3@C composite. Fig. 2 depicts the preparation process of the porous core-shelled Sn/Sn4P3@C material. After the tin atoms are electrochemically dissolved from Sn80P20 in H2SO4 solution, the nanostructured Sn/Sn4P3 framework can be successfully produced. The thin poly-dopamine layer is deposited on the Sn/Sn4P3 composite surface by 10
polymerization as the carbon source and then followed by carbonization to obtain the final anode sample.
Fig. 3. (a, b) SEM images of the Sn/Sn4P3 sample after electrochemical corrosion. (c, d) SEM images of the Sn/Sn4P3@C after covering carbon layer. The resulted structure after dealloying the sample alloy was characterized by SEM. As shown in Fig. S4, the Sn80P20 precursor alloy foil has smooth and compact microstructure. It can be shown in Fig. 3a that the interconnected backbone composed of several large channels distributed on the surface after dealloying the Sn80P20 alloy in 0.1 M H2SO4 solution by electrochemical corrosion. High-magnification SEM image (Fig. 3b) illustrates that a great deal of pores penetrated through the integral nanoscaled skeleton with the typical ligament size about 300 nm. Based on observations above, it is clear that upon dealloying Sn80P20 alloy, the uniform nanoporous architecture can be successfully obtained with interconnected skeleton and hollow channels. As shown in 11
Fig. 3c, upon coating the carbon layer, the original porous structure was well preserved without the surfaces pores blocked. From Fig. 3d, it is clearly observed that the ligaments underwent certain coarsening to 450 nm due to the high temperature carbonization.
Fig. 4. Elemental mappings of C (b), P (c), and Sn (d) in the Sn/Sn 4P3@C sample matching to Fig. 4a. The elemental distributions of the dealloyed sample were investigated by element mapping. It can be illustrated in Fig. 4a-d that the product mostly contained the Sn, P, and C elements, which uniformly dispersed in the whole configuration, further demonstrating that the uniform nanocomposite of metal phosphide has been successfully enwrapped by one homogeneous carbon layer.
12
Fig. 5. (a, b, c) TEM images of the Sn/Sn4P3@C sample. The detailed microstructure of the Sn/Sn4P3@C sample was further characterized by TEM. It can be illustrated in Fig. 5a that the black framework indicated the formation of interconnected ligaments, while the inner white region suggested the existence of embedded pores. From the high magnification TEM image (Fig. 5b&c), it can be obviously observed that the carbon layer was uniformly deposited on the surface of porous Sn/Sn4P3 with the typical thickness around 10 nm. The high-resolution TEM image was not obtained because the strong electron beam led to the sample resolved under the state of high applied voltage. Based on the experimental observations above, 13
electrochemical corrosion of the commercial Sn80P20 alloy followed by carbonization can successfully generate the uniform carbon encapsulated porous Sn4P3 nanocomposite with metallic Sn well distributed in the porous architecture. On the basis of the BET results (Fig. S5), the pore voids of Sn/Sn4P3@C have the wider size distribution with the relatively small pores examined. It should be mentioned that by the intrinsic restriction of BET technique the large pores around several hundred nanometers cann’t be measured. Combining with the BET results as well as the SEM and TEM observations, it is conclusive that there are multiscale pore voids distributed in the Sn/Sn4P3@C sample, which can provide the sufficient space for the volume change during the ion insertion/extraction. The lithium storage performances of the porous Sn/Sn4P3@C were examined in half-coin cells configuration utilizing the lithium foil as both reference and counter electrodes. Fig. 6a illustrated the cycle voltamemetric (CV) curves of the as-made anode for the initial three cycles in the voltage range of 0.01 to 3.0 V (vs Li+/Li) at 0.1 mV s-1. In the initial cathodic process, the broad peak appeared at 1.66 V and a reduction peak appeared at 0.59 V, which was characterized by the lithium insertion process as well as the conversion of Sn4P3 into metallic Sn and the formation of Li3P (Eq. (1)). The reduction peak located at 0.32 V was regarded as the reversible lithium alloying reactions with tin (Eq. (2)) as well as the formation of a SEI layer. The two obvious reduction peaks at 0.59 and 0.32 V could be assigned to the lithium ion and electron transportation, thus resulting in the rapid and adequate disintegration of the active composite. The anodic peaks at 0.53 and 0.66 V in the initial charging process 14
could be attributed to the transformation of the Lix Sn alloy and Li3P into Sn4P3 (Eq. (3) and Eq. (4)). After the first cycle, it was clear that the cathodic peaks shift obviously, indicating the configurational change in the electrode material after the Li+ insertion in the initial cycle and the formation of SEI layer. However, the oxidation peaks had no obvious shift, implying the great reversibility owing to the protective carbon layer and porous architecture. On account of these points of discussion, the reaction mechanism of Sn4P3 during the initial discharge was shown to be as follows: Sn4P3 + 9Li+ + 9e-1 → 4Sn + 3Li3P (1) 2Sn + 5Li+ + 5e-1 → Li5Sn2 (2) During the first charging process: Li5Sn2 → 2Sn + 5Li+ + 5e-1 (3). Fig. 6b illustrated the initial three galvanostatic charge-discharge curves of Sn/Sn4P3 @C electrode at 100 mA g -1. The discharge capacities of Sn/Sn4P3@C in the first three cycles were as high as 1332.2, 1036.9, and 964.8 mA h g-1. The initially high discharge capacity of the anode could be attributed to the formation of SEI film as well as the large surface area of Sn/Sn4P3@C sample, which furnished abundant surface reaction sites. The extensive quantity of SEI films also consumed a part of Li ions, and gave rise to irreversible insertion/extraction Li+ in the following cycles. Furthermore, there were two plateau regions during the initial discharge. The former presumably resulted from the reaction between Li and P atoms, and the latter corresponded to the alloying reaction between Sn and Li. For comparison, Fig. S6 showed the initial three galvanostatic charge-discharge curves of Sn/Sn4P3@C electrode at 2000 mA g-1. The 15
trend of degradation was more serious than that of 100 mA g-1 due to the more severe polarization under the high rate. 4Sn + 3Li3P → Sn4P3 + 9Li+ + 9e-1 (4)
Fig. 6. (a) CV curves of Sn/Sn4P3@C. (b) The galvanostatic charge-discharge curves at 100 mA g-1 for the first three cycles. Cycling performances and CEs of Sn4P3 and Sn/Sn4P3 @C anodes at 100 (c) and 1000 mA g-1 (d), respectively. To further clarify the enhanced lithium storage performances for the Sn/Sn4P3@C anode, Fig. 6c illustrated the galvanostatic cycling performance of the as-made anode at 100 mA g-1 with the profile of the pure porous Sn4P3 attached for comparison. As demonstrated in Fig. 6c, the Sn/Sn4P3@C anode demonstrated an initial discharge capacity of 1332 mA h g-1 with the slight degradation in the following cycles. And then, the capacity of Sn/Sn4P3@C began to rise, which was possibly because that the carbon layer protected the porous Sn4P3 with the growth of a stable polymeric SEI film preventing
the
collapse
of the
framework 16
structure
during
lithium
ion
insertion/extraction. For the pure Sn4P3, the capacity decayed rapidly with a linear mode. The excellent electrochemical storage performance of Sn/Sn4P3@C composite in LIBs was also demonstrated by evaluating the capacity retention. The capacity retention of the Sn/Sn4P3@C nanocomposite was 62.8%, which was assigned to the volume expansion as well as the decomposition of electrolyte during the reaction between Sn/Sn4P3@C and Li+. In addition, the capacity retention of the Sn/Sn4P3@C anode was dramatically higher than that of the pure Sn4P3 (9.3%). The cycling performance of the Sn/Sn4P3@C material was also evaluated at high rate of 1000 mA g-1. Fig. 6d showed that the initial capacity of Sn/Sn4P3@C anode was 1261 mA h g-1, which was higher than that of pure Sn4P3 electrode (1209 mA h g-1). It was fascinating to find that the capacity of Sn/Sn4P3@C anode illustrated a slight rising tendency from the 100th cycle, which was normally observed for transition metal phosphide anodes. It is generally ascribed to the growth of a stable polymeric SEI film, which can prevent the collapse of the framework structure during lithium ion insertion/extraction [39]. Upon 700 cycles the Sn/Sn4P3 @C anode offered a high discharge capacity of 589 mA h g-1, which was much higher than the pure Sn4P3 (114 mA h g-1). The corresponding initial capacity retention of the Sn/Sn4P3@C nanocomposite (~46.7%) at the higher rate of 1000 mA g-1 was much higher than that of Sn4P3 (9.4%). It was noted that the CE of the Sn/Sn4P3@C composite suddenly increased to 97.1% in the 5th cycle and further sustained between 97 and 100% in the following several hundred cycles. Compared with the Sn4P3, the as-made Sn/Sn4P3@C anode showed much higher capacity as well the longer cycling life for LIBs. The higher 17
reversible capacity and superior cycling stability of the Sn/Sn4P3@C could be assigned to the introduction of the carbon layer and the metallic Sn matrix. As described in Fig. 7a, the reversible capacities around 701, 596, 594, 565, and 526 mA h g-1 were achieved at the current densities of 100, 300, 500, 1000, and 2000 mA g -1, and the specific capacity rebounded back to 698 mA h g-1 when the current density was reduced back to 100 mA g-1. Respectively, which were dramatically higher than that of the Sn4P3 electrode under the identical situation. In particular, the Sn/Sn4P3@C electrode demonstrated the good cycling stability and high specific capacity when the current density reached 2000 mA g-1, illustrating its superior high rate capability at the larger current density. Upon long term cycling, the capacity of pure Sn4P3 decreased to relatively inferior value, in contrast, the capacity of the Sn/Sn4P3@C anode entirely recovered. These observations validated that the Sn/Sn4P3@C electrode had the higher rate capability as well as remarkably improved discharge capacity. The EIS measurement for the two anodes was gauged by utilizing fresh cells to explore the possible contribution of doping Sn and covering carbon layer. As indicated by the Nyquist profiles in Fig. 7b, both samples displayed a depressed semicircle in the charge transfer impedance of cells, and a straight line in the low-frequency range indicating the mass transfer of Li+. It was detected that in the diameter of the semicircle for the two samples are obviously different. For clear comparison, the EIS results have been fitted to be a semicircle and a line as shown in Fig. 7b inset. The small diameter for Sn/Sn4P3 @C indicated that the charge transfer of the Sn4P3 became more facile owing to the adulteration of excellent conductive metallic Sn and carbon layer, thus resulting 18
in a superior electrochemical performance.
Fig. 7. (a) Rate performance of Sn4P3 and Sn/Sn4P3@C electrodes. (b) The nyquist plots of the fresh Sn4P3 and Sn/Sn4P3 @C samples. Inset is the fitting EIS. To further recommend the lithium storage performances of Sn/Sn4P3@C nanocomposite, we made a comparison of the Sn/Sn4P3@C anode with other Sn4P3-based anode materials in Table 1 [17,21,40]. It was noted that the Sn/Sn4P3@C electrode displayed the highest reversible capacity of 837 and 589 mA h g-1 at the current densities of 100 and 1000 mA g -1, respectively, even after 100 and 700 cycles. We would highlight that the as-prepared Sn/Sn4P3@C material has the excellent discharge capacity and remarkable stable cyclability among all of the formerly recorded Sn4P3-based anode materials. The excellent discharge capacity and the higher capacity retention of the Sn/Sn4P3@C demonstrated its potential application in the lithium storage fields. 19
Table 1 Comparison of the electrochemical performances of various Sn4P3 materials with this work. Anode
Current
Cycle number
Remaining
density
capacity
(mA g-1)
(mA h g-1)
Sn4P3 hollow
100
20
261
microspheres Sn4P3
Electrolytes
Ref
EC/DMC/DEC
[40]
=1:1:1 50
60
481
EC/DMC/DEC
[21]
=1:1:1 Sn4P3
100
100
416
EC/DMC/DEC
[21]
=1:1:1 Sn4P3
500
100
210
EC/DMC/DEC
[21]
=1:1:1 Sn4P3/graphite
100
100
651
EC/EMC
[17]
=3:7 Sn/Sn4P3@C
100
100
837
EC/DMC/EMC =1:1:1+3%(wt)VC
Sn/Sn4P3@C
1000
700
589
EC/DMC/EMC =1:1:1+3%(wt)VC
This work This work
EC: Ethylene carbonate; DMC: Dimethyl carbonate; EMC: Ethylene methyl carbonate; DC: Diethyl carbonate; VC: vinylene carbonate
4. Conclusions Porous Sn/Sn4P3 composite with protected carbon layer anchored on the network matrix was easily prepared by facile dealloying of Sn80P20 alloy under mild conditions followed by carbonization. Owing to the incorporation of well-conductive Sn and cladding of carbon layer, the Sn/Sn4P3@C nanocomposite certified much reinforced cycling stability and high rate capability, particularly at large current rates. With the superiorities of unique performances as well as scalable fabrication, the as-prepared Sn/Sn4P3 @C composite holds excellent application potential for lithium storage. Acknowledgements This work was supported by the National Natural Science Foundation of 20
China (51772133) and Shandong Province (ZR2017JL022), and China postdoctoral science foundation (2017M622117), and the program for Taishan Scholar of Shandong province (ts201712048). References [1] J.T. Jiang, K. Zhu, Y.Z. Fang, H.Z. Wang, K. Ye, J. Yan, G.L. Wang, K. Cheng, L.M. Zhou, D.X. Cao, Coralloidal carbon-encapsulated CoP nanoparticles generated on biomass carbon as a high-rate and stable electrode material for lithium-ion batteries, J. Colloid Interface Sci. 530 (2018) 579-585. [2] B.B. Wang, X. Zhang, X.J. Liu, G. Wang, H. Wang, J.T. Bai, Rational design of Fe3O4@C yolk-shell nanorods constituting a stable anode for high-performance Li/Na-ion batteries, J. Colloid Interface Sci. 528 (2018) 225-236. [3] Y. Yang, G.Y. Huang, H.Y. Sun, M. Ahmad, Q.Y. Mou, H.M. Zhang, Preparation and electrochemical properties of mesoporous NiCo 2O4 double-hemisphere used as anode for lithium-ion battery, J. Colloid Interface Sci. 529 (2018) 357-365. [4] J. Lee, Y. Wu, Z.B. Peng, Hetero-nanostructured materials for high-power lithium ion batteries, J. Colloid Interface Sci. 529 (2018) 505-519. [5] X.Z. Hu, Y. Jin, B. Zhu, Z, Liu, D.F. Xu, Y.D. Guan, M.Y. Sun, F.L. Liu, Tuning density of Si nanoparticles on graphene sheets in graphene-Si aerogels for stable lithium ion batteries, J. Colloid Interface Sci. 532 (2018) 738-745. [6] Y.Z. Yan, H.L. Tang, J.S. Li, F. Wu, T.B. Wu, R. Wang, D. Liu, M. Pan, Z.Z. 21
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27
Graphical Abstract:
Porous
core-shelled
Sn/Sn4P3@C
composite
is
easily
fabricated
by
electrochemically etching SnP alloy followed by covering the uniform carbon layer, which shows superior lithium-storage performances with higher specific capacity, much enhanced rate capability, as well as better cycling stability compared with pure porous Sn4P3. The Sn/Sn4P3 @C material holds promising application potential as an alternative anode in the lithium storage fields.
28