In-situ embedding cobalt-doped copper sulfide within ultrathin carbon nanosheets for superior lithium storage performance

Journal of Colloid and Interface Science 566 (2020) 1–10

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In-situ embedding cobalt-doped copper sulfide within ultrathin carbon nanosheets for superior lithium storage performance Huilin Qing a, Ruirui Wang a, Ziliang Chen a, Mingming Li b, Lilei Zhang b, Yong-Ning Zhou a, Renbing Wu a,c a

Department of Materials Science, Fudan University, Shanghai 200433, PR China Yantai Chungway New Energy Technology Co., Ltd., Yantai 264000, PR China c The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, PR China b

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

a r t i c l e

i n f o

Article history: Received 6 November 2019 Revised 17 January 2020 Accepted 18 January 2020 Available online 20 January 2020 Keywords: Cobalt doping Metal-organic frameworks Two-dimensional structure Ultrathin nanosheet Copper sulfides Lithium-ion batteries

a b s t r a c t Construction of well-defined hybrid composites consisting of transition metal sulfides and twodimensional (2D) carbon nanosheets as high-performance anodes for lithium-ion batteries (LIBs) is of great significance but remains challenging. Herein, we have developed a novel strategy to in-situ fabricate a nanohybrid composites consisting of cobalt-doped copper sulfides nanoparticles embedded in 2D carbon nanosheets (2D Co-Cu2S@C) through a one-pot sulfurization of 2D nanosheet-like Co-doped copperbased metal-organic frameworks (MOFs) precursors. When applied as LIBs anodes, the as-prepared 2D Co-Cu2S@C composites could deliver a specific capacity of 780 mAh g
1 at 0.5 A g
1 after 300 cycles and a high-rate capability with 209 mAh g
1 at 5 A g
1, superior to most reported copper sulfidebased anodes. The exceptional performance could be attributed to the synergism of ultrathin structure (~4 nm), appropriate cobalt doping and strong carbon coupling, resulting in the shortened paths for Li+ transportation, enlarged exposing surface for Li+ adsorption, enhanced electric conductivity for charge transfer as well as robust mechanical property against volume expansion. Ó 2020 Elsevier Inc. All rights reserved.

1. Introduction Nickel-metal hydride batteries and lithium-ion batteries (LIBs) have been extensively applied in the field of portable electronics [1–10]. However, with the sustainable development of society and economy, the energy density and rate capability of these E-mail address: [email protected] (R. Wu) https://doi.org/10.1016/j.jcis.2020.01.068 0021-9797/Ó 2020 Elsevier Inc. All rights reserved.

energy storage devices should be further improved to fulfill the requirements of the large-scale vehicle power systems for electric vehicles (EVs) and hybrid electric vehicles (HEVs) [11–16]. Since the performance of LIBs is mainly determined by their anode and cathode materials, numerous endeavors have been undertaken to explore the novel electrodes to substitute the conventional ones [17,18]. In the search of the advanced anode materials, transition metal sulfides (TMSs) have demonstrated great potential as

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H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

alternatives to the conventional graphite anode owing to their abundant earth resource, intriguing crystal structure and high lithium storage capability [19–24]. Nevertheless, TMS bulks suffer from the inferior electrical conductivity, unsatisfied charge adsorption ability, and large volume expansion during charge/discharge process, leading to the poor capacity, stability and rate performance [25–27]. In this regard, the improvement of the TMS electrodes in terms of the high-rate and long-term cyclability for EVs and HEVs is significant. In order to improve the cycle stability, hybridizing TMS with conductive carbon matrix (TMS@C) has been demonstrated as an effective approach because the carbon matrix not only enhances the electrical conductivity of composites but also can serve as the robust matrix to buffer the volume variation and maintain the structural stability [27–29]. For instance, Xu et al. reported that the composites constructed by rooting NixCo1-xSe2 nanosheets on surface of carbon dodecahedra could deliver a specific capacity of 1667 mAh g
1 at 2 A g
1 after 600 cycles [30]. Lin et al. synthesized a composite of NiS2 and CoS2 hetero-nanocrystals encapsulated in N-doped carbon nanocubes, showing a capacity of 600 mAh g
1 after 100 cycles at 1 A g
1 [31]. Despite great advances that have been achieved, the rate capabilities of most reported TMS@C composites are still unsatisfactory and need to be further ameliorated. Recent studies have revealed that the hetero-cation doping can powerfully enhance the lithium storage kinetics of TMS electrodes by introducing edge/surface defects and local charge polarization, which would remarkably contribute to the charge and electron transfer [32–34]. For instance, Gao et al. encapsulated iron-doped Co9S8 on N, S-doped carbon nanosheet as anode material for LIBs, which exhibits a high rate capability of 331 mAh g
1 at 1 A g
1 [35]. Meng et al. prepared Cu0.33Co0.67S2 hexagonal sheets and found they could achieve a dramatically improved rate capability of 558 mAh g
1 at 10 A g
1 by copper cation doping [36]. In addition to the cation-doping, construction of two-dimensional (2D) architecture can also greatly increase the rate capability as its high surface-to-volume ratio can readily trigger the charge-transfer reactions including surface or near surface redox reactions as compared with their three-dimensional (3D) counterparts [37–41]. Currently, there are two ways to prepare the 2D TMS-based electrodes. One is to directly synthesize 2D architecture of TMS, the other is to anchor the TMS nanoparticles on the 2D graphene [42]. However, the former one usually makes electrode suffer from collapses during the repeated charge/discharge process [43], while the latter one usually makes electrode possess a weak coupling between TMS and graphene, which may cause particle aggregation and increase the paths for ion transportation [44,45]. As such, it is highly desired to construct a robust and well-coupled 2D TMS@Cbased nanohybrid. Based on the above discussions, herein, we have proposed a 2D MOF-driven strategy to in-situ fabricate well-coupled TMS@C hybrid composite [46–48]. Specifically, Copper ions and 1,3,5benzenetricarboxylic acid (H3BTC) can form 2D Cu(HBTC)(H2O)3 (Cu(HBTC)-1) precursor and then 2D architecture consisting of cobalt-doped copper sulfides embedding within ultrathin carbon nanosheets (2D Co-Cu2S@C) can be prepared by one-pot sulfurization of 2D cobalt-doped Cu(HBTC)-1 precursor (Co-Cu(HBTC)-1). The in-situ fabrication not only allows hybridization of TMS with carbon but also ensures the strong coupling effect between the components, which is beneficial to preventing the aggregation during the iterative charge/discharge process and enhancing the charge transfer [23,35]. Such a nanohybrid through a MOF-driven strategy possesses the following merits for lithium storage: (i) the coupled carbon nanosheet functions as a buffering matrix to alleviate the volume variation of electrode during charge/discharge process and facilitates the structural stability, (ii) the Co doping in Cu2S lattice could effectively improve the conductivity and accelerate

the electron transfer, (iii) the porous ultrathin feature of architecture ensures a rapid mass diffusion and ion transport. Therefore, when employed as anode material for LIB, 2D Co-Cu2S@C composite exhibits a superior electrochemical performance in terms of an outstanding cycle durability (780 mAh g
1 at 0.5 A g
1 after 300 cycles) and excellent rate capability (209 mAh g
1 at 5 A g
1) in comparison with 2D Cu2S@C, 3D Cu2S@C and pure Cu2S anodes. 2. Experimental section 2.1. Chemicals The chemicals were used as received without further purification. Copper (II) nitrate trihydrate (Cu(NO3)23H2O, Sinopharm Chemical Reagent Co., 99.5%), copper (II) chloride dehydrate (CuCl22H2O, Sinopharm Chemical Reagent Co., 99%), cobalt (II) nitrate hexahydrate (Co(NO3)26H2O, Sinopharm Chemical Reagent Co., 98+%), 1,3,5-benzenetricarboxylic acid (H3BTC, Aldrich, 95%), sodium sulfide (Na2S, Sinopharm Chemical Reagent Co., 99%), polyvinylpyrrolidone (PVP, Aldrich, K-30), sodium L(+)-ascorbate (Ourchem, 99%), sodium hydroxide (Sinopharm Chemical Reagent Co., 96%), sublimed sulfur (Aldrich, 99.5%), N-methyl pyrrolidone (NMP), carbon nanotube (CNT), polyvinylidene fluoride (PVDF), methanol (Aladdin, 99%) and ethanol (Aladdin, 99.99%). Deionized water was used for all experiments. 2.2. Preparation of HKUST-1 precursors HKUST-1 is another MOF constructed by copper ion and H3BTC with 3D octahedron-like structure, which was reported and named by Williams et al. at Hong Kong University of Science and Technology (HKUST) in 1999 [49]. 875 mg of H3BTC was dissolved in methanol (50 mL). Then a 50 mL of methanol solution dissolved with 1.82 g of Cu(NO3)23H2O was added to above solution followed by stirring for 2 h at room temperature. The blue precipitates were collected by centrifugation and washed by methanol for several times. HKUST-1 precursors were finally obtained after drying blue precipitates in oven at 60 °C overnight. 2.3. Preparation of 2D Co-doped Cu(HBTC)-1-1 and 2D Cu(HBTC)-1 precursors A facile bottom-up route at room temperature and ambient pressure was used to synthesize 2D Cu(HBTC)-1 and 2D Co-doped Cu(HBTC)-1 (denoted as 2D Co-Cu(HBTC)-1) [50]. 136.8 mg of CuCl22H2O and 800 mg of PVP were dissolved into 160 mL of deionized water. Then, 10 mL of 0.8 mol L–1 NaOH solution and 10 mL of 0.4 mol L–1 sodium L(+)-ascorbate solution were mixed with the above solution under stirring to form an orange mixture. After continuously stirring for 5 min, the mixture was centrifuged and washed with ethanol for twice to obtain orange solids. The orange solids were resuspended into 40 mL ethanol to form Cu2O-ethanol suspension. 3.2 g of PVP and 425 mg of Co (NO3)26H2O were dissolved in 240 mL of deionized water and then 32 mL of ethanol solution dissolved with 1.6 g of H3BTC was added to the above solution under stirring. After that, the as-prepared 40 mL of Cu2O-ethanol suspension was poured to the above solution followed by stirring for 36 h. The light blue solids were collected by centrifugation and washed with deionized water for twice. After dried in oven overnight at 60 °C, 2D Co-Cu(HBTC)-1– 1 products were obtained. The formation of Cu(HBTC)-1was similar to that of 2D Co-Cu(HBTC)-1–1 but using additional of Co (NO3)26H2O. Note that the ratio of ethanol to water is the key for the formation of 2D morphology. To further investigate the influence of cobalt, Co-Cu(HBTC)-1 with different amounts of

H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

cobalt were synthesized. In the synthesis process, different amounts of Co(NO3)26H2O were added to form Co-doped Cu (HBTC)-1, denoted as 2D Co-Cu(HBTC)-1-n, respectively, where the amount of used Co(NO3)26H2O increases with number n. 2.4. Preparations of 2D Co-Cu2S@C, 2D Cu2S@C and 3D Cu2S@C hybrid composites A clean porcelain boat loaded with 2D Co-Cu(HBTC)-1-1 precursors and sulfur powders with a mass ratio of 1:1 was put into a tube furnace and then heated to 600 °C with a ramp rate of 3 °C min
1 under flowing argon. After being maintained at this temperature for 2 h, the 2D Co-Cu2S@C hybrid composites were successfully prepared. The preparations of 2D Cu2S@C and 3D Cu2S@C were similar to that making for 2D Co-Cu2S@C hybrid composites but with 2D Cu(HBTC)-1 and HKUST-1 as precursors, respectively. For comparison, 2D Co-Cu2S@C-2 and 2D Co-Cu2S@C-3 with different cobalt contents were formed via sulfurization of 2D Co-Cu (HBTC)-1–2 and 2D Co-Cu2S@C-1–3, respectively. 2.5. Preparation of pure Cu2S 32 mg of sodium sulfide powders were dispersed in 20 mL of deionized water. The sodium sulfide aqueous solution was then added into the aforementioned Cu2O-ethanol suspension. The mixture was aged for 25 min under stirring. The black products were collected and dried. A clean porcelain boat loaded with dry black products and sulfur powders with a mass ration of 1:1 was put into a tube furnace and then heated to 600 °C with a ramp rate of 3 °C min
1 under flowing argon. After being maintained at this temperature for 2 h, the pure 3D Cu2S particles were successfully prepared. 2.6. Characterizations X-ray diffractometer (XRD) (D8 ADVANCE) equipped with Cu Ka radiation was utilized to analyze the crystal structures of dried sample powders. The morphology and microstructure characterizations were performed on a field-emission scanning electron microscope (FESEM, ZEISS, Ultra-55) at 5.0 kV and transmission electron microscope (TEM, JEOL-2100) at 120 kV. The element distribution was investigated using energy dispersive X-ray spectrometry (EDS) elemental analysis attached on JEOL-2100. In order to identify the element states of products, X-ray photoelectron spectroscopy (XPS) measurements were conducted on a Kratos XSAM-800 with a radiation source of Al Ka X-ray. Witec Alpha 300 confocal microscope Raman spectrometer was applied to examine the carbon structure. The porosity was analyzed through N2 sorption measurements at 77 K using an Autosorb IQ Gas Sorption instrument. BET surface area, the mesopores with width over 2 nm were analyzed by Barrett-Joyner-Halenda (BJH) method and the micropores were studied by Horvath-Kawazoe (H-K) method. The mass contents of metal sulfides and carbon in the products were determined through Thermogravimetric analysis (TGA) using a METTLER TOLEDO TGA2. The thickness of product was determined by using a Bruker Dimension Icon atomic force microscope (AFM). 2.7. Electrochemical measurements Active material, carbon nanotubes (CNTs) and polyvinylidene fluoride (PVDF) were mixed in N-methyl pyrrolidone (NMP) solution with a weight ratio of 7:2:1 to produce a homogeneous slurry. The as-prepared slurry was then spread over the copper foil evenly, followed by a drying process in a vacuum oven at 70 °C for 12 h. Coin half cells were assembled in a glove-box (<0.1 ppm, H2O

3

and O2) filled with Ar with the cut electrode slice as the working electrode, Li metal as the counter electrode and Celgard 2400 membrane as the separator. The electrolyte was 1 M LiPF6 in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1:1 in volume). To evaluate the electrochemical performance of the above-assembled cells, galvanostatic cycling was carried out over the 0.01–3.0 V voltage range on a battery test system (LandCT2001A) at room temperature. The cyclic voltammetry (CV) tests ranging from 0.01 to 3.00 V and electrochemical impedance spectroscopy (EIS) in a fixed frequency range of 0.1 MHz to 0.01 Hz were conducted on an electrochemical measurement system (AutoLab-PGSTAT302N).

3. Results and discussion The synthesis process of ultrathin 2D Co-Cu2S@C nanosheets is schematically shown in Fig. 1. Firstly, Cu2O nanocubes are prepared via a co-precipitation reaction of CuCl2 and NaOH followed by a reduction of sodium ascorbate. After that, with a further bottom-up method by using Cu2O nanocubes as templates at room temperature, 2D Co-doped Cu(HBTC)-1 nanosheets are formed. Note that both small molecular (i.e., water and oxygen) and proper cobalt anions are needed to react with Cu2O in this process. Finally, the ultrathin 2D Co-Cu2S@C hybrid composites can be obtained by a sulfurization of 2D Co-Cu(HBTC)-1 at 600 °C under flowing Ar gas, during which the sulfur powders are sublimated and react with metal species to form Co-Cu2S while the organic linker (H3BTC) is transformed to carbon, resulting in the in-situ formation of strongly-coupled heterostructure composed of embedding CoCu2S within 2D ultrathin carbon nanosheet. The morphologies and microstructures of the as-prepared products were checked by FESEM and TEM, respectively. As shown in Fig. S1, the low-magnification FESEM images clearly present that the obtained 2D Co-Cu(HBTC)-1-1 precursors are rectangle nanosheets with size of 2–3 mm. As expected, the 2D structure could be kept well after sulfurization treatment (Fig. 2a and b). Interestingly, the Co-Cu2S@C composite shows an ultrathin nanosheet-like structure after sulfurization. Such a decrease of the thickness for the 2D nanosheet could be ascribed to the architecture shrunk caused by the thermal stress and decomposition. Moreover, a higher-magnification FESEM image (Fig. 2c) also reveals that the smooth surfaces for precursors have been changed to the rough and porous surfaces for 2D Co-Cu2S@C composites. In consistent with the FESEM observations, the ultrathin 2D sheetlike structure can be clearly seen in TEM image (Fig. 2d). As shown in an enlarged TEM image (Fig. 2e), plentiful black dots with an average diameter of 12 ± 4.5 nm are uniformly distributed on the nanosheets, suggesting the well-coupling between particles and carbon sheet. Two districts are perceived from the highresolution TEM image (Fig. 2f), where clear lattice fringes with a spacing of 0.324 nm indexed to the (1 0 2) crystal planes of Co-Cu2S and the surrounding disordered area corresponded to amorphous carbon. Note that lattice defects can also be clearly identified (Fig. 2f and Fig. S2). EDS element mapping results in Fig. 2g–k confirms the coexistence of S, Cu, C, and Co elements within the nanosheet. EDS spectra recorded from the middle sheet in Fig. S4 confirms that the atomic percentage of Co is ~0.43%. Furthermore, the ultrathin feature is corroborated by an AFM image, which shows an average thickness of ~4.3 nm for 2D Co-Cu2S@C composite (Fig. 2l). For comparison, the morphologies of 3D Cu2S@C and 2D Cu2S@C are also investigated by FESEM and the results are presented in Fig. S3, from which 3D Cu2S@C has the octahedral structures with rough surfaces while 2D Cu2S@C shows a similar morphology to that of 2D Co-Cu2S@C composite. Nevertheless, the AFM result (Fig. S5) suggests that 2D Cu2S@C

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H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

Fig. 1. Schematic synthesis route for ultrathin 2D Co-Cu2S@C composites.

Fig. 2. FESEM images of (a) 2D Co-Cu(HBTC)-1 precursor and (b, c) 2D Co-Cu2S@C composite; (d) low-magnification TEM image, (e) high-magnification TEM image and (f) high-resolution TEM (HRTEM) image of 2D Co-Cu2S@C; (g-k) the corresponding EDS elemental mapping images; (l) AFM image for the 2D Co-Cu2S@C. Fig. 2e inset: the distribution of particle size.

H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

composite possesses a slightly thinner thickness of ~3.4 nm than that of 2D Co-Cu2S@C, implying that cobalt doping might cause the increase of thickness. The phase structures of the as-synthesized hybrid composites were characterized by XRD. Fig. S6a shows the XRD patterns of precursors, indicating the successful synthesis of HKUST-1 and Cu (HBTC)-1. Notably, although the 2D Co-doped precursor has a semblable diffraction pattern with Cu(HBTC)-1, the further comparison between them (Fig. S6b) demonstrated that the positions of peaks in Co-Cu(HBTC)-1 are slightly shifted to a lower angle direction. As presented in Fig. 3a, the sulfurized products i.e. 2D Cu2S@C and 3D Cu2S@C show similar XRD patterns, where the whole diffraction peaks can be assigned to chalcocite Cu2S-type phase (JCPDS card No. 72-1071). A slight shift of peak positions to a lower angle direction caused by Co doping is observed in the XRD pattern of 2D CoCu2S@C when compared with Cu2S@C. Moreover, due to the fact that no diffraction peaks of Co or its sulfides are found in the XRD pattern of 2D Co-Cu2S@C nanosheet, it could be concluded the successful doping of Co into Cu2S lattice. On the other hand, the broad peak between 20 and 35 degree is attributed to carbon. The carbon content was further investigated by thermogravimetric analysis (TGA) at a heating rate of 10 °C min
1 from room temperature to 800 °C in air. As shown in Fig. 3b, there is a gradual increase of weight between 200 and 350 °C, which is accumulatively ascribed to the formation of CuOCuSO4. At approximately 300 °C, a sudden mass drop at a rapid rate in the curve is identified, which is due to the decomposition of residual oxygen groups [51]. Finally, the 2D Co-Cu2S@C, 2D Cu2S@C and 3D Co-Cu2S@C composites have been completely converted to CuO with a 73.3%, 58.3% and 64.5% weight remaining, respectively. Thus, the mass contents of carbon in 2D Co-Cu2S@C, 2D Cu2S@C and 3D Co-Cu2S@C were

5

calculated to be ~26.7%, 41.7% and 35.5%, respectively. The carbon structure of the composites was further investigated by Raman spectra (Fig. 3c). Four deconvoluted peaks at ~1190 cm
1, ~1350 cm
1, ~1440 cm
1 and ~1560 cm
1 could be ascribed to the sp2 carbon outside graphene (I-band), disordered carbon (Dband), distortion (D’’-band) and graphite carbon (G-band), respectively [52–55]. As an indicator for the degree of graphitization, the peak area ratios of D-band to G-band (AreaD/AreaG) are 1.08, 1.13 and 1.14 for 2D Co-Cu2S@C, 2D Cu2S@C and 3D Cu2S@C, respectively, indicating the partial graphitization. In comparison with the 2D Cu2S@C and 3D Cu2S@C, the carbon in 2D Co-Cu2S@C possesses the highest graphitization degree, suggesting that the doped cobalt can promote the carbon graphitization during the pyrolysis of organic linkers [56]. The porosity characteristics of as-prepared products are qualified by N2 sorption measurements. As shown in Fig. 3d and Fig. S7, the typical type-IV isotherms with H3 hysteresis loops from the adsorption-desorption isotherm are identified, indicating the mesoporous structures of composites. Fig. 3d inset shows the pore size distribution of 2D Co-Cu2S@C composite based on Barrett-Joyner-Halenda (BJH) method, suggesting the existence of mesopores. Moreover, 2D Co-Cu2S@C and 2D Cu2S@C have BET surface areas of 164.5 m2 g
1 and 231.8 m2 g
1, respectively, higher than that of 3D Cu2S@C (138.6 m2 g
1). The large surface area and mesoporous structure for 2D Co-Cu2S@C are believed to be beneficial to providing enriched active sites for lithium ions and vast channels for the penetration of ions and electrolyte [28,57,58]. To examine the chemical states of the elements on the surface of the sulfurized products, X-ray photoelectron spectroscopy (XPS) investigations were performed. The C 1s spectrum in Fig. 4a describes two main peaks centered at 284.8 and 285.5 eV

Fig. 3. (a) XRD patterns, (b) TGA curve and (c) Raman spectra of 3D Cu2S@C, 2D Cu2S@C and 2D Co-Cu2S@C composites; (d) N2 isothermal adsorption-desorption curve at 77 K of 2D Co-Cu2S@C composite. Fig. 3d inset: pore size distribution based on BJH method.

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Fig. 4. High-resolution (a) C 1s, (b) Co 2p, (c) Cu 2p, and (d) S 2p XPS spectra for the 2D Co-Cu2S @C composite.

at the representative of CAC/C@C bond and CAS bond, respectively [21]. The existence of CAS bond reveals the coupling between carbon and sulfur atoms. The oxidation state of Co 2p is shown in Fig. 4b, where two pairs of spin-orbit doublets are displayed, i.e., Co 2p3/2 located at 781.0 eV and Co 2p1/2 located at 798.8 eV, demonstrating the presence of Co species [21]. The highresolution Cu 2p spectrum of 2D Co-Cu2S@C is deconvoluted into two spin orbits at 952.5 eV for Cu 2p1/2 and 932.6 eV for Cu 2p3/2, with an energy difference of ~20 eV [59]. Additionally, the four peaks centered at 162.8, 161.7, 165.2 and 164.1 eV are assigned to Cu-S 2p1/2, Cu-S 2p3/2, CAS 2p1/2 and CAS 2p3/2, respectively [59,60]. Notably, the four peaks for Cu 2p1/2, Cu 2p3/2, Cu-S 2p1/2 and Cu-S 2p3/2 shift to lower binding energy upon introducing cobalt into the Cu2S lattice when compared with those of bare 2D Cu2S@C and 3D Cu2S@C (Figs. S8a and S8b). These deviations could be due to a changed electronic structure caused by cobalt doping [61,62]. Copper sulfide is a kind of conversion-typed anode materials, which has reversible oxidation-reduction reaction with Li+ in the discharge and charge process [63]:

Cu2 S + 2Liþ + 2e— $Li2 S + Cu

ð1Þ

The lithium storage mechanism and properties of the 2D CoCu2S@C nanosheets are firstly studied by cyclic voltammetry (CV). Fig. 5a manifests the first CV curves at a sweep rate of 0.1 mV s
1 in the voltage range of 0.01–3.0 V vs Li/Li+. In the first cathodic scan, a broad peak ranged 0.3–0.9 V and two sharp peaks located at around 2.10, and 1.56 V are observed. The broad peak at 0.3–0.9 V is related to the irreversible formation of solid electrolyte

interphase (SEI) film. It is likely to associate the small peak at 2.10 eV with the adsorption of Li+ on the surface of composite [59]. The peak at 1.56 V is attributed to the reduction of Cu+ and the formation of Li2S, which slightly moves to a lower voltage in the subsequent cycles. In the first anodic scan, two peaks centered at 1.91 V and 2.33 V are identified, which could be due to the oxidation of Cu and the formation of Li2S [59,64]. In the subsequent three cathodic-anodic scans, the CV curves are almost overlapped, indicating the great reversibility of the electrochemical reaction in the 2D Co-Cu2S@C composite electrode. Fig. 5b displays the first, fifth and 100th discharge-charge voltage profiles of the 2D CoCu2S@C composite electrode at a current density of 0.5 A g
1. There are several plateaus consistent with the previous analysis from CV result in the profiles (Fig. 5a). Note that a high initial discharge capacitance of 1356 mAh g
1 and a reversible charge capacitance of 858 mAh g
1 are observed with an initial columbic efficiency of 63.3%, mainly resulting from the formation of SEI film. The result of cycle performance at 0.5 A g
1 is presented in Fig. 5c. After 300 cycles, 2D Co-Cu2S@C nanosheet electrode can deliver a reversible discharge capacity of 780 mAh g
1, which is not only much higher than those of 2D Cu2S@C (670 mAh g
1), 3D Cu2S@C (234 mAh g
1) and pure Cu2S (206 mAh g
1), but also outperforms the recently reported Cu-based anode materials (Table S1). The rate capability of 2D Co-Cu2S@C nanosheet electrode is studied at different current densities ranging from 0.2 to 5 A g
1 (shown in Fig. 5d). At current densities of 0.2, 0.5, 1.0, 2.0 and 5.0 A g
1, the electrode can deliver an average reversible capacity of 950, 783, 584, 386, and 209, respectively, which shows much higher capacities in comparison with those of 2D Cu2S@C, 3D Cu2S@C and Cu2S, especially at the high current density. When the current density returns to 0.2

H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

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Fig. 5. (a) CV curves of the first four cycles at the scan rate of 0.1 mV s
1 and (b) charge-discharge profiles of 2D Co-Cu2S @C composite at the current density of 0.5 A g
1; (c) cyclic performance at 0.5 A g
1 and (d) rate performance at various current densities for pure Cu2S, 3D Cu2S@C, 2D Cu2S@C and 2D Co-Cu2S@C; (e) cycling performance of the 2D Co-Cu2S@C composite electrode at 1 A g
1 after activation.

A g
1, the electrode can still recover to a reversible capacity as high as 988 mAh g
1. To further investigate the stability of 2D CoCu2S@C nanosheet electrode, a long-term cycling performance is carried out at a current density of 1 A g
1 after activation (Fig. 5e). It is obviously seen that after 1000 cycles, the electrode still exhibits a reversible capacity of 570 mAh g
1, showing an excellent cycle stability. Moreover, after 300 cycles at 0.5 A g
1, the batteries were disassembled and the morphologies of the 2D Co-Cu2S@C and Cu2S electrodes were characterized by FESEM. As seen in Fig. S9, the pristine structure of Cu2S is dramatically damaged, while 2D structure in Co-Cu2S@C is well-maintained as expected, demonstrating its good structural integrity. By hybridizing with carbon nanosheet, the volume variation of electrode is also greatly alleviated, and the structural stability is efficiently facilitated. To investigate the electrochemical lithium storage kinetics of 2D Co-Cu2S@C nanosheet electrode, CV tests at different scan rates ranging from 0.2 to 1.0 mV s
1 were carried out after chargedischarge cycle testing. As shown in Fig. 6a, the current increases with the scan rate. The electrochemical lithium storage kinetics can be analyzed by the following equations [65,66]:

i ¼ a  vb

ð2Þ

log i ¼ b  log v þ log a

ð3Þ

where a and b mean variable constants, i is the current, and v is the scan rate. Four peaks are picked to modulate the relationship between log i and log v. By linear fitting, the values of b are 0.76 and 0.87 for peak 1 and peak 2 in the oxidation process, respectively (Fig. 6b). For peak 3 and 4 in the reduction process, the values of b are 0.85 and 0.84, respectively. The values of b represent that 2D Co-Cu2S@C nanosheet electrode includes diffusion-controlled and capacitive-controlled storage. A quantitative assessment to evaluate the contribution ratios of diffusion-controlled and capacitivecontrolled storage can be conducted based on the following equation:

iðvÞ ¼ k1  v þ k2  v 1=2

ð4Þ

where i(v) means the current at a certain voltage and k1 is a variable constant as well as k2. By calculating k1 and k2 constants, the current contributed by diffusion and capacitive effect can be extracted. Fig. 6c demonstrates that capacitive effect totally contributes to

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Fig. 6. 2D Co-Cu2S@C composite electrode: (a) CV curves at various scan rates from 0.2 to 1.0 mV s
1, (b) linear relationship between log i and log v at the oxidized and reduced state, (c) the capacitive and diffusion-controlled contribution to the overall charge storage composite electrode at a sweep rate of 0.4 mV s
1, (d) the normalized contribution percentage from capacitive-controlled and diffusion-controlled behavior to the overall charge storage at different sweep rates.

83.6% current at the scan rate of 0.4 mV s
1. As shown in Fig. 6d, the contribution ratio of capacitive effect increases with the scan rate and well maintains over 80%. At a sweep rate of 1 mV s
1, it could even reach as high as 92.9%. The lithium storage kinetics of 3D Cu2S@C and 2D Cu2S@C are further investigated. As displayed in Fig. S10 and Fig. S11, only 74.1% and 88.3% of total capacity are resulted from the capacitive contribution for 3D Cu2S@C and 2D Cu2S at 1 mV s
1, respectively, much smaller than that of 2D Co-Cu2S@C. The dominant of capacity contribution from pseudocapacitive reaction for 2D Co-Cu2S@C could be ascribed to 2D structure and cobalt doping. Firstly, the ultrathin 2D structure of 2D Co-Cu2S@C and 2D Cu2S@C plays an essential role in narrowing Li+ diffusion and increasing the contact surface with electrolyte, which could greatly boost the kinetics when compared with octahedral structure of 3D Cu2S@C. Furthermore, cobalt doping can create a wealth of lattice defects as demonstrated in the above-presented highresolution TEM images (Fig. S2). These defects have been proven to have lower adsorption energy, resulting in energetically favorable structures to alleviate the ion diffusion during lithiation and delithiation process [34,56,67]. The electrochemical impedance spectroscopy (EIS) measurements are also carried out to further verify the fast lithium storage kinetics of 2D Co-Cu2S@C (Fig. S12). As presented in Fig. S13a, all of the electrochemical impedance spectra are fitted with separate semicircles in high and middle frequency as well as a straight line in low frequency, which are related to SEI film resistance (Rs), change transfer resistance (Rct), and Warburg impedance (W0), respectively. It is noticeable that 2D Co-Cu2S has the smallest semicircle among all measured composites, revealing the lowest charge transfer

resistance (Fig. S13b). Additionally, in the low frequency, the lines of 2D Co-Cu2S@C and 2D Cu2S@C composites are much straighter than that of 3D Cu2S@C, indicating the facilitated lithium ion diffusion. Moreover, the amount of doped Co may have an important influence on the thickness of composite. FESEM images shown in Fig. S14 suggest the thickness of composite increase with the increasing of added cobalt. The cycle performance (Fig. S15) and EIS results (Fig. S16) of composites with different cobalt contents demonstrate that the 2D Co-Cu2S@C has the highest specific capacity and the smallest Rct in comparison with 2D Co-Cu2S@C-2 and 2D Co-Cu2S@C-3. By cobalt doping, the electrochemical property of electrode can be improved. Nevertheless, with the continuous addition of cobalt, the thickness of composites increases, which may lead to a decline of conductivity and unsatisfied electrochemical performance. In short, 2D Co-Cu2S@C electrode renders an extraordinary electric conductivity and prominent lithium ion diffusion as a result of synergetic effects of deliberately-designed 2D structure and appropriate cobalt doping, which leads to a superior electrochemical property for LIB.

4. Conclusions In summary, ultrathin 2D Co-Cu2S@C nanosheets have been synthesized through a novel 2D MOF-driven strategy, which involves the preparation of well-defined 2D Co-Cu(HBTC)-1 nanosheet precursors and subsequent sulfurization of the precursor at moderate temperature. The as-synthesized 2D Co-Cu2S@C nanosheets simultaneously integrates several appealing design rationales for high-performance electrodes including the porous

H. Qing et al. / Journal of Colloid and Interface Science 566 (2020) 1–10

ultrathin feature of architecture, remarkable conductivity and excellent mechanical stability of strongly coupled carbon nanosheets. As compared with those of 2D Cu2S@C and 3D Cu2S@C, the electrochemical property of 2D Co-Cu2S @C composite was significantly improved by constructing 2D architecture and cobalt doping, representing a new record for copper sulfides-based anodes [59,60,64]. The current work may provide guidance for the structural design of advanced transition metal-based materials, which shows great potential to develop advanced materials for next-generation anodes of LIBs. CRediT authorship contribution statement Huilin Qing: Conceptualization, Methodology, Investigation, Writing-Original Draft, Writing-Review & Editing. Ruirui Wang: Investigation, Formal Analysis. Ziliang Chen: Formal Analysis, Writing - Orignal Draft. Mingming Li: Investigation. Lilei Zhang: Investigation. Yong-Ning Zhou: Formal analysis. Renbing Wu: Conceptualization, Supervision, Formal Analysis, Writing - Original Draft, Writing - Review & Editing. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements This work was financially supported by the fund of The State Key Laboratory of Refractories and Metallurgy in Wuhan University of Science and Technology (G201709), the China Postdoctoral Science Foundation (Grant No. 2018M640337 and 2019T120304), and the Fudan’s Undergraduate Research Opportunities Program (FDUROP, No. 18060). Appendix A. Supplementary material Supplementary data to this article can be found online at https://doi.org/10.1016/j.jcis.2020.01.068. References [1] H. Pan, Y. Liu, M. Gao, Y. Lei, Q. Wang, A study of the structural and electrochemical properties of La0.7Mg0.3(Ni0.85Co0.15)x (x=2.5–5.0) hydrogen storage alloys, J. Electrochem. Soc. 150 (2003) A565–A570. [2] H. Pan, Y. Liu, M. Gao, Y. Zhu, Y. Lei, Q. Wang, An investigation on the structural and electrochemical properties of La0.7Mg0.3(Ni0.85Co0.15)x (x=3.15–3.80) hydrogen storage electrode alloys, J. Alloys Compd. 351 (2003) 228–234. [3] S. Goriparti, E. Miele, F.D. Angelis, E.D. Fabrizio, R.P. Zaccaria, C. Capiglia, Review on recent progress of nanostructured anode materials for Li-ion batteries, J. Power Sources 257 (2014) 421–443. [4] B. Liao, Y. Lei, L. Chen, G. Lu, H. Pan, Q. Wang, A study on the structure and electrochemical properties of La2Mg(Ni0.95M0.05)9 (M=Co, Mn, Fe, Al, Cu, Sn) hydrogen storage electrode alloys, J. Alloys Compd. 376 (2004) 186–195. [5] B. Liao, Y. Lei, L. Chen, G. Lu, H. Pan, Q. Wang, Effect of the La/Mg ratio on the structure and electrochemical properties of LaxMg3
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