Iron-Doping-Induced Phase Transformation in Dual-Carbon-Confined

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Iron-Doping-Induced Phase Transformation in Dual Carbon Confined Cobalt Diselenide Enabling Superior Lithium Storage Yang Liu, Ziliang Chen, Huaxian Jia, Hongbin Xu, Miao Liu, and Renbing Wu ACS Nano, Just Accepted Manuscript • Publication Date (Web): 09 May 2019 Downloaded from http://pubs.acs.org on May 9, 2019

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Iron-Doping-Induced Phase Transformation in Dual Carbon Confined Cobalt Diselenide Enabling Superior Lithium Storage Yang Liu,†, ‡,& Ziliang Chen,†,& Huaxian Jia,§, ǁ Hongbin Xu,† Miao Liu,§, ǁ, ⁋ Renbing Wu†,* †Department

of Materials Science, Fudan University, Shanghai 200433, P. R. China

‡Department

of Chemistry, Fudan University, Shanghai 200433, P. R. China

§Institute ǁCenter

of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China

of Materials Science and Optoelectronics Engineering, University of Chinese Academy of

Sciences, Beijing 100049, P. R. China ⁋Songshan

Lake Materials Laboratory, Dongguan, Guangdong 523808, P. R. China

*Corresponding author: [email protected]

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ABSTRACT: Transition metal chalcogenides (TMCs) have been investigated as promising anodes for high-performance lithium-ion batteries, but they usually suffer from poor conductivity and large volume variation, thus leading to unsatisfied performance. Although nanostructure engineering and hybridization with conductive materials have been proposed to address this concerning, a better performance towards practical device applications is still highly required. Herein, we reported an iron-doping-induced structural phase transition from pyrite-type (cubic) to marcasite-type (orthorhombic) phases in porous carbon/rGO-coupled CoSe2. The dual carbon confined orthorhombic CoSe2 (oFexCo1–xSe2@NC@rGO) composites exhibit dramatically enhanced lithium storage performance (920 mAh g−1 after 1000 cycles at 1.0 A g−1) over cubic CoSe2-based composites (c-CoSe2@NC@rGO). The combined experimental studies and densityfunctional theory calculations reveal that this doping-induced structural phase transformation strategy could create a favorable electronic structure and assure a rapid charge transfer. These results demonstrate that the phase-transformation engineering may provide another opportunity in the design of high-performance TMCs-based electrodes. KEYWORDS: cobalt diselenide; carbon matrix; phase transition; anode; lithium-ion batteries

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As one promising energy storage and supply device, nickel-metal hydride batteries and lithium-ion batteries (LIBs) have been widely utilized in the portable electronics territory.1–10 Nowadays, with rapidly increasing market demands for the large-scale electric vehicles and renewable power stations, it is an urgent and critical issue to further develop LIBs with higher energy density, better cycling stability, and more superior rate capability.11,12 Given the fact that the performances of LIBs are mainly dependent on the properties of electrode materials, intensive endeavors have been devoted to exploring for advanced cathode and anode materials to replace conventional ones.13–15 Recent advances in the exploration of advanced anode materials have shown that most of transition metal chalcogenides (TMCs) are attractive alternatives because of their intrinsic safety and much higher theoretical capacity (500~1000 mAh g–1) as compared to the commercial graphite anode (372 mAh g–1).16–20 Unfortunately, these bulk TMCs suffer from large volume change and sluggish Li-ion diffusion kinetic during the charge-discharge process, leading to an inferior cyclability and poor rate capability. To solve these problems, combining the nanosized TMC with carbon matrix (e.g., amorphous/graphitic carbon, carbon nanotubes and graphene) into rationally designed hybrid composites has been demonstrated to be a well-established strategy. In such composites, the nanosized TMC shortens the pathway for Li+ diffusion while the carbon matrix not only improves the electrical conductivity but also alleviates the large volume change during the repeated charge-discharge cycles, thus improving the lithium storage performance.21–26 For instance, Lou et al. embedded the CoSe nanoparticles within the amorphous carbon nanoboxes, delivering a capacity of 711 mAh g−1 after 100 cycles at 500 mA g−1.21 Wang et al. reported that CoS2 growing on amorphous carbon‐coated multiwalled carbon nanotubes could exhibit a reversible capacity of up to 598 mAh g−1 after 500 cycles at 2 A g−1.22 Joo et al. fabricated the three-dimensional (3D) free-standing 3


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anode consisting of FeS nanoparticles anchored on carbon nanofiber, showing a capacity as high as 535 mAh g−1 after 320 cycles at 0.5 A g−1.23 Despite continuous efforts that have been made in the research direction of nanostructure and hybridization engineering, there is still an urgent requirement to further boost the lithium storage performance of TMCs and make them satisfy the practical applications. On the other hand, TMCs can be formed in different polymorphic phases by modulating electron population in the d orbitals. Polymorph engineering of TMCs as well as the structure-property correlation originating from its polymorphs have received particular interest.27−31 For example, structural phase transition from the α-MnS to metastable β-MnS could be induced by a lithiation/delithiation process, which leads to an improved kinetic property due to the fact that metastable β-MnS has smaller charge-transfer resistance and more favorable ion mobility in the host lattice than the stable counterpart.27 2H (trigonal prismatic) to 1T (octahedra) phase transformation in MoS2 achieved by an electrochemical lithiation was also reported to improve the ion/charge transfer ability in the host lattice, which enabled the 1T-MoS2 to exhibit good lithium storage property.28 Despite these pioneering achievements, the application of polymorph engineering in lithium storage is still in its infancy. Moreover, the developed structural phase transition is mainly induced by the electrochemical lithiation, which not only is susceptible to the charge uniformity of the sample but also generates lattice stress during the process, resulting in inhomogeneous phase change. Herein, we have proposed an iron-doping-induced structural phase transition engineering in designing CoSe2-based free-standing anodes, which successfully achieves a significantly improved capacity, cyclability and rate capability in lithium-ion batteries. The preparation of anodes involves a precipitation and solvothermal process to form a 3D reduced graphene oxide (rGO)-wrapped bimetallic (Fe, Co)-Prussian blue analogue (PBA) 4


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precursor, followed by a simultaneous carbonization and selenization process. Specifically, during the thermal‐induced reaction process, the introduction of Fe-doping enables the phase transition from pyrite-type (cubic) to marcasite-type (orthorhombic) CoSe2, while organic ligands liberated from PBA converted into N-doped graphitic carbon (NC) accompanied by the formation of rGO, generating a hybrid composite consisting of dual carbon-coupled o-FexCo1–xSe2 nanoparticles (denoted as o-FexCo1–xSe2@NC@rGO). Benefiting from the free-standing 3D configuration, the highly dispersed bimetal selenides active nanoparticles and higher electrical conductivity, the orthorhombic FexCo1– xSe2@NC@rGO

anode exhibits an exceptional electrochemical performance in terms of

ultrahigh reversible capacity (920 mAh g−1 at 1.0 A g−1 after 1000 cycles) and superior rate capability (618 and 466 mAh g−1 at 6.4 and 12.8 A g−1, respectively). RESULTS AND DISCUSSION The synthetic route for 3D o-FexCo1–xSe2@NC@rGO aerogel was schematically illustrated in Figure 1. First, Co2+ and Fe2+ ions with positive charge were deposited on the surface of graphene oxide (GO) with carboxylic group through the electrostatic adherence. By further adding excessive [Co(CN)6]3‒ into above solution, a coordination interaction occurs between metal ions and [Co(CN)6]3‒, leading to the in-situ formation of Fe-Co-PBA on the surface of GO (Fe-Co-PBA@GO). Following that, water-dispersible Fe-CoPBA@GO was self-assembled as 3D Fe-Co-PBA@rGO aerogel via chemical reduction followed by freeze-drying. Finally, the 3D hierarchical o-FexCo1–xSe2@NC@rGO aerogels could be fabricated by a simultaneous carbonization and selenization of 3D FeCo-PBA@rGO aerogel at 600 °C for 2 h under flowing nitrogen gas. The as-fabricated precursors and their selenization products were firstly analyzed by X-ray diffraction (XRD). Figure S1 displayed the XRD patterns of the precursors and the 5


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GO treated by a VC-Na reduction. As expected, the precursors were consisted of Fe-CoPBA and rGO phases, suggesting the successful preparation of Fe-Co-PBA@rGO. For comparison, the XRD pattern of the selenization products was shown in Figure S2, from which orthorhombic CoSe2-type (o-CoSe2) phase and graphitic carbon could be identified, indicating a complete conversion from precursors to o-CoSe2 phase and carbon. Note that as compared to the theoretical o-CoSe2 phase (PDF#53-0449), the diffraction peaks of our o-CoSe2 type phase shifted to the higher 2θ angle, which should be due to the doping of Fe2+ with smaller atomic radius (0.61 Å) than Co2+ (0.65 Å), inducing the lattice shrinkage of CoSe2. To further confirm this point, on the one hand, inductively coupled plasma mass spectrometry (ICP-MS) was employed to analysis the atomic ratio of metal specie in selenization products. The atomic ratio of Fe to Co is 2.3:3.5, close to the designed value (3:7). On the other hand, the crystal structure model based on o-CoSe2 is employed to perform the Rietveld refinement for the XRD pattern of selenization products. As seen in Figure 2a, the refined results are well consistent with the observed data and the refined lattice parameters were indeed smaller than those of theoretical o-CoSe2 (Table S1). Above results strongly demonstrate the formation of orthorhombic FexCo1–xSe2 phase (o-FexCo1– xSe2).

To better elucidate the role of Fe-doping in the crystal structure of CoSe2, the XRD

pattern of the products obtained by the selenization of bare Co-PBA@rGO precursor under the same thermal-annealing condition was analyzed. As presented in Figure 2b and Table S2, they were basically composed of cubic CoSe2 phase (c-CoSe2) and graphitic carbon. Notably, the effect of Fe-doping amount on the phase structure of products was also investigated. The designed atomic ratio of Fe to Co was varied from 0:1, 1:9, 2:8, 3:7, 5:5, to 1:0 (the Fe content x was 0, 0.1, 0.2, 0.3, 0.5 and 1, respectively). As shown in Figure S3 and Figure S4, the relative abundance of o-FexCo1–xSe2 phase became more as the increase of the amount of Fe-doping, and a complete phase conversion occurred when the 6


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designed Fe/Co atomic ratio reached to 3:7. Such results strongly corroborated the Fedoping-induced structural phase-transition from c-CoSe2 to metastable o-FexCo1–xSe2 (schematic illustration was displayed in Figure 2c). Since c-CoSe2 phase has higher thermodynamic stability than o-CoSe2 phase, the orthorhombic-to-cubic phase transition in CoSe2 could be easily triggered by chemical or thermal ways;31,32 while the reverse cubic-to-orthorhombic phase transition has rarely been reported. In this regard, the Fedoping-induced cubic-to-orthorhombic phase transition in CoSe2 is possibly due to the fact that the doping of Fe with weaker electronegativity than Co in c-CoSe2 could drive the rotation of half Se–Se pairs (Figure S5).30 This metastable bimetal chalcogenide induced by anion doping is expected to show an excellent lithium storage kinetic. To analyze the carbon structure, Raman spectra of the products obtained using both Fe-Co-PBA@rGO and Co-PBA@rGO precursors were displayed in Figure 2d, in which two broad peaks were observed at 1344 cm‒1 (D-band) and 1590 cm‒1 (G-band), respectively. The relative intensity ratios of D-band and G-band (ID/IG) for both two products were close to 1.1, indicating their high graphitization degree of carbon, which is beneficial to the electron transfer during electrochemical process.33 To further confirm the presence of porous carbon in the selenized Fe-Co-PBA@rGO product, the Raman spectra of pristine rGO and the product obtained by direct selenization of pristine Fe-Co-PBA precursor were given in Figure S6, from which it could be apparently observed two peaks corresponding to D- and G-band of carbon in both samples, indicating the formation of porous carbon during the selenization of Fe-Co-PBA. Moreover, the relative intensity ratio of ID/IG for rGO and the selenized Fe-Co-PBA product was 0.98 and 1.32, respectively. Since the relative intensity ratio of ID/IG for the selenized Fe-Co-PBA@rGO product was 1.1, which located between 0.98 and 1.32, further demonstrating that the porous carbon and rGO were coexisted in the selenized Fe-Co-PBA@rGO product. Further 7


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thermogravimetric (TG) tests revealed that the weight ratios of porous carbon and rGO in the selenized Fe-Co-PBA@rGO product were about 17.8% and 26.9%, respectively, close to those in the selenized Co-PBA@rGO product, implying that the Fe-doping has a negligible effect on the amount of porous carbon and rGO in the resulted selenized products (Figure S7). Furthermore, N2 sorption measurement was carried out to reveal the pore features of the products using Fe-Co-PBA@rGO as precursor. As presented in Figure 2e, both micropores and mesopores are existed in the products, being well consistent with the pore size distribution (inset in Figure 2e). The specific surface area determined by the Brunauer–Emmett–Teller (BET) method is approximately 188 m2 g−1. Such a large BET specific surface area with hierarchical pores is believed to contribute to accelerate the mass transfer and the penetration of electrolyte into the active materials.34 The morphology evolution from Fe-Co-PBA@rGO precursor to the resulting product was further revealed by field emission scan electron microscope (FESEM). As shown in Figure 3a, the monolithic Fe-Co-PBA@rGO aerogel exhibited an interconnected 3D porous network. The magnified FESEM image further showed that the PBA nanoparticles with sizes mainly ranging from 10 to 100 nm were tightly wrapped within the interior of the wrinkled rGO framework without aggregation (Figure 3b and 3c). After a selenization at 600 °C, the interconnected 3D porous network maintained well (Figure 3d) and the particles still inherited the polyhedron-like morphology (Figure 3e and 3f) without collapsing. The magnified FESEM images suggested that the wrinkled rGO layer conformably coated on the surface of polyhedron-like particles (Figure 3g and 3h), implying a strong confinement of rGO towards nanoparticles. In addition, the elemental mapping recorded from the representative particle further demonstrated the uniform distribution of Co, Fe and Se elements within the polyhedron-like particle, which was confined by rGO as well as the N-doped carbon matrix (Figure 3i‒n). The morphology of 8


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the c-CoSe2@NC@rGO composite has also been examined by FESEM (Figure S8). Obviously, similar to the o-FexCo1–xSe2@NC@rGO, the monolithic c-CoSe2@NC@rGO aerogel had also an interconnected 3D porous network with the CoSe2@NC subunits uniformly wrapped by the wrinkled rGO framework. Transmission electron microscope (TEM) characterizations were further conducted to investigate the microstructure of products. Being well agreement with the FESEM observations, the low-magnification TEM images (Figure 4a‒c) clearly displayed the uniform wrapping of polyhedron-like particles by rGO layers. The high-resolution TEM (HRTEM) images (Figure 4d and 4e) further revealed that the polyhedron-like particles were mainly composed of the o-FexCo1–xSe2 core and ultrathin graphitic carbon shell. Note that the lattice fringes with interlayer spacings of 0.259 nm and 0.289 nm correspond to the (111) and (101) planes of o-FexCo1–xSe2 phase, respectively. These lattice distances are slightly smaller than those of o-CoSe2, in accordance with the XRD results, which may be ascribed to the doping of Fe in o-CoSe2. Additionally, both Figure 4d and 4e presented the coupled rGO matrix, showing a strong dual carbon confinement to polyhedron-like particles. Thus, based on above characterization results, the product obtained through the selenization of Fe-Co-PBA@rGO was defined as the o-FexCo1–xSe2@NC@rGO, as schematically illustrated in Figure 4f. To further examine the chemical states of the elements on the surface of o-FexCo1– xSe2@NC@rGO

composites, the X-ray photoelectron spectroscopy (XPS) measurements

were employed. The full survey spectrum displayed in Figure 5a proclaimed the presence of Co, Fe, Se, C and N. The Co 2p spectrum (Figure 5b) could be deconvoluted into two pairs of spin-orbit doublets, i.e., Co 2p3/2 located at 780.6 eV and Co 2p1/2 located at 796.8 eV.35 For the Fe 2p spectrum, the peaks only belonged to the Fe2+ was deconvoluted (Figure 5c), indicating the formation of (Co, Fe)Se2 phase. The peaks located at 54.6 eV 9


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and 55.7 eV could be assigned to Se 3d5/2 and Se 3d3/2, respectively (Figure 5d). Additionally, in the high-resolution C 1s spectrum, two main peaks with binding energies of 284.5 and 285.6 eV were ascribed to C–C and C–N bonds, respectively (Figure 5e).36 The presence of C–N bond strongly confirms the successful doping of N element in the carbon matrix derived from organic linker in PBA. Correspondingly, the high-resolution N 1s spectra (Figure 5f) revealed the existence of pyridinic N, pyrrolic N, and graphitic N, which is believed to promote the lithium storage capability.37,38 To investigate whether the o-FexCo1–xSe2@NC@rGO has enhanced lithium storage performance as compared with c-CoSe2@NC@rGO, the electrochemical behavior of assynthesized products was checked as free-standing anode materials for LIBs. Firstly, the possible lithium storage mechanism of o-FexCo1–xSe2@NC@rGO was explored by evaluating the CV curves against initial three cycles at a scan rate of 0.1 mV s−1 between 0.01 and 3.0 V. As seen in Figure 6a, in the first discharge process, a small peak positioned at ~1.41 V might be due to the insertion of Li into the host lattice of o-FexCo1–xSe2.39 With further discharging, a strong peak appeared at around 1.24 V could be assigned to the conversion reaction from o-FexCo1–xSe2 to Co, Fe and Li2Se. Following that, peak positioned at around 0.7 V might be ascribed to the formation of SEI layer.40 In the charge process, no peaks assigned to the Fe-based selenides could be found and two anodic peaks positioned at 2.14 and 2.28 V might be attributed to the stepwise formation of o-FexCo1– xSe2.

Upon subsequent cycling, two cathodic peaks during the lithiation process were

observed at 1.38 and 1.68 V, corresponding to the stepwise formation of Co/Fe and Li2Se, respectively. The curves in the second cycle is different from that in the first cycle, probably due to the microstructure alteration after the first lithiation/delithiation cycle, which also commonly appears in other conversion reaction‐type anode materials.41−44 The CV curves of the o-FexCo1–xSe2@NC@rGO composite are almost overlapped after the first 10


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cycle, implying a high degree of reversibility of the composite for the lithiation and delithiation reactions. Such results can be also confirmed by the galvanostatic discharge/charge (GDC) curves at a current density of 0.2 A g−1 (Figure 6b), in which several voltages extracted from the discharge and charge plateaus are corresponding to the peaks observed in the CVs, and the GDC curves in the second cycle are overlapped well with those in the third cycle, corroborating the high reversibility. The discharge and charge capacities of o-FexCo1–xSe2@NC@rGO electrode in the first cycle are ~1617 and ~1124 mAh g−1, respectively, corresponding to an initial coulombic efficiency (CE) of 67.4%. The first CE for our o-FexCo1–xSe2@NC@rGO is relatively low45,46, which might be due to the following two reasons: i) The porous hierarchical o-FexCo1–xSe2@NC@rGO electrode with a large specific surface area consumed much more electrolyte to form the solid electrolyte interphase (SEI) layer during the first discharge process.15,20,41,42,47,48 Because the formation of SEI layer is largely irreversible, the capacity during the initial charge process can be apparently decreased. ii) The inserted lithium ions during the initial discharge process could not be completely detached from the o-FexCo1–xSe2@NC@rGO electrode, which may also result in the decrease of charge capacity.41,49Above results also suggest the possible reversible transformation of o-FexCo1–xSe2 phase during charge/discharge process. To check this point, ex-situ XRD tests for products at different charge/discharge states were carried out. Figure S9 shows FESEM images of the o-FexCo1–xSe2@NC powders pressed onto nickel foam. It can be seen that the o-FexCo1–xSe2@NC powders are tightly anchored on the nickel foam and their hierarchical microstructure can be still maintained very well. The ex-situ XRD test results were shown in Figure S10, from which it could be known the reversible phase transformation of o-FexCo1–xSe2 phase during charge/discharge process (detailed analysis is given below Figure S10). Such a reversible transformation for bimetal selenides 11


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might be due to the facile reaction thermodynamic and dynamic condition.50 It should be remarked that the redox peaks in the CV curves was asymmetry, which could be due to the complex de/lithiation reaction of o-FexCo1–xSe2 and carbon as well as the formation/decomposition of SEI layer. To further evaluate the electrochemical performance of o-FexCo1–xSe2@NC@rGO electrode, the cycle performance of o-FexCo1–xSe2@NC@rGO was displayed in Figure 6c, in which a discharge capacity as high as 920 mAh g−1 could be maintained at a current density of 1.0 A g−1 after 1000 cycles, demonstrating an outstanding cycle stability with high reversible capacity. To the best of our knowledge, such a high capacity after longterm cycling at high current density may be the best among the Co-based selenides (Table S3). The morphology of the cycled o-FexCo1–xSe2@NC@rGO was checked and the result was shown in Figure S11. Obviously, it could maintain well as compared to the fresh one, showing a high structural stability against cycles. On the other hand, the Fe-free cCoSe2@NC@rGO anode only exhibited a capacity of 453 mAh g−1 after 1000 cycles under the same test condition, which was much lower than that of o-FexCo1–xSe2@NC@rGO, implying that the Fe-doping-induced structural phase transition could effectively boost the electrochemical performance of Co-based selenides. Notably, the phenomenon that capacity firstly decreased and then increased was also observed in the metal-based compounds, which might be related to the optimization of stable SEI and the enhanced kinetics of Li-diffusion caused by a fast lithiation-induced reactivation.46,51–55 In particular, under initial several lithiation/delithiation cycles, metal chalcogenides with porous structure may suffer from mechanical degradation caused by the large volume changes accompanying by the conversion reaction. Meanwhile, the outside SEI layer would fracture and re-arrange, leading to an instable and thick SEI layer and thus possibly resulting in the capacity fading. Nevertheless, with further cycling, the porous structure 12


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could be self-restructured well and exhibited an optimized thin SEI layer and enlarged space buffering volume variations, which may bring a greatly reactivated lithium storage capacity and improved cyclability. To alleviate the issue that capacity obviously change against the cycles, the prelithiation might be an effective approach, by which the porous structures can be optimized firstly.56,57 In order to investigate the effect of multicomponent composition on the electrochemical properties of o-FexCo1–xSe2@NC@rGO, o-FexCo1– xSe2@NC,

c-CoSe2@NC, rGO and NC samples have been evaluated as anode materials in

LiBs. As shown in Figure S12, after 1000 cycles, o-FexCo1–xSe2@NC, c-CoSe2@NC, rGO and NC electrodes could deliver capacities of only 275, 187, 259, and 56 mAh g1 at a current density of 1.0 A g1, respectively, strongly highlighting the synergistic effect of multicomponent composition. Figure 6d further compared the rate performance of o-FexCo1–xSe2@NC@rGO anode and c-CoSe2@NC@rGO anode. When the current densities are 1.6, 3.2, 6.4 and 12.8 A g−1, capacities of o-FexCo1–xSe2@NC@rGO anode are 899, 770, 618 and 466 mAh g−1, respectively, much higher than those of c-CoSe2@NC@rGO anode (440, 226, 120 and 50 mA h g−1 at 1.6, 3.2, 6.4 and 12.8 A g−1). When the current density recovers to 2.0 A g−1, the capacity of o-FexCo1–xSe2@NC@rGO could be stabilized at 786 mAh g−1. It is noteworthy that such excellent high-rate cycling stability of o-FexCo1–xSe2@NC@rGO anode is also superior to previously reported state-of-the-art Co-based anodes (Figure 6e). To further highlight the exceptional cycle stability and rate performance, the long-term cycling of o-FexCo1–xSe2@NC@rGO anode was performed at a high current density of 10.0 A g−1. Notably, a capacity as high as 412 mAh g−1 was maintained after 3000 cycles at 10.0 A g−1 (Figure S13), showing outstanding rate performance. To explore the underlying reason responsible for the outstanding rate capability, the kinetic study of synthesized o-FexCo1–xSe2@NC@rGO anode is further analyzed. Figure 13


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7a shows the CV curves at various scan rate from 0.1 to 1.0 mV s−1. It can be seen that the peaks in CV curve show relatively small variation in potential as a function of sweep rate, indicating the fast charge transfer processes.20 Moreover, it is generally believed that the Li storage mechanism, including diffusion-controlled electrochemical reaction and pseudocapacitive process, can be revealed on the basis of following equations:36 i = avb

(1)

log(i) = b·log(v) + log(a)

(2)

where i and v are the current and the scan rate, respectively; and a and b are adjustable parameters. When the b-value approaches to 1.0, the lithium storage mechanism is pseudocapacitive. When the b-value is 0.5, a diffusion-controlled electrochemical reaction should be the lithium storage mechanism. As shown in Figure 7b, the calculated b values of the corresponding redox peaks are 0.81, 0.72, 0.74 and 0.71, respectively, suggesting that the electrochemical process of o-FexCo1–xSe2@NC@rGO anode are associated with both diffusion-controlled reaction and pseudocapacitive behavior. Moreover, the contribution of pseudocapacitive behavior and diffusion-controlled reaction is further quantitatively extracted according to the following equation:36 i = k1vb + k2v1/2

(3)

where the pseudocapacitive behavior and diffusion-controlled contribution is weighted by k1 and k2, respectively. Figure 7c and Figure 7d show the contribution ratios for the pseudocapacitive current (light orange region) in comparison with the total current. The pseudocapacitive behavior increases along the increase of the scan rates. Especially, when the scan rate is increased to 1.0 mV s−1, the pseudocapacitive contributes above 78% of the total capacity, which is much larger than that 68% of c-CoSe2@NC@rGO (Figure S14). Such results strongly implied that the Fe-doping-induced phase transition could effectively improve the lithiation kinetics. To better clarify the excellent electrochemical performance 14


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of o-FexCo1–xSe2@NC@rGO anode, the electrochemical impedance spectra (EIS) were further investigated (Figure S15). All of the Nyquist plots consist of semicircles in the high-to-middle frequency regions and slanted lines in the low-frequency regions, in which the semicircles are mainly attributed to the charge-transfer resistance (Rct).61 Obviously, the fresh o-FexCo1–xSe2@NC@rGO electrode exhibits the lower charge-transfer resistance as compared to c-CoSe2@NC@rGO, which strongly demonstrated its faster charge transfer ability (Figure S15a). Furthermore, the Nyquist plots of o-FexCo1–xSe2@NC@rGO and c-CoSe2@NC@rGO electrodes after the 1st and 100th cycles have also been provided and compared (Figure S15b), from which two main conclusions could be extracted: i) The charge-transfer resistance after the 100th cycle is smaller than those without cycle and after the 1st cycle for both o-FexCo1–xSe2@NC@rGO and c-CoSe2@NC@rGO electrodes, suggesting the enhanced reaction kinetics against cycling, which might be due to the optimization of SEI film and the activation process during cycling.20,58 ii) After 100 cycles, the o-FexCo1–xSe2@NC@rGO electrode exhibited much lower charge-transfer resistance (Rct=47 Ω) compared to that of c-CoSe2@NC@rGO (Rct=78 Ω), strongly indicating the superior charge transfer ability of o-FexCo1–xSe2 to c-CoSe2 during cycles. To provide more insight into the electrical conductivity and lithium storage kinetics enhanced by the Fe-doping-induced phase transition, density functional theory (DFT) calculations were performed to compare c-CoSe2, o-CoSe2, and Fe-doped o-CoSe2. The charge-density distributions of c-CoSe2, o-CoSe2, and Fe-doped o-CoSe2 are visualized in Figure 8a‒c, all of which show the strong interaction between Co/Fe and Se ions. Noticeably, upon incorporating the Fe atoms into o-CoSe2 host lattice, a charge polarization was generated around the local environment of Fe ion (through comparing the circle parts marked by the red dot in Figure 8b and 8c), which might be beneficial to enhancing the binding ability with the lithium ions.59 On the other hand, the densities of 15


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states (DOS) results show that the Fe-doped o-CoSe2 has the stronger density of states (integrated area: ~4.4) near the Fermi level (within the range from ‒0.2 to 0.2 eV) as compared to those of pristine c-CoSe2 (integrated area: ~3.9) and o-CoSe2 (integrated area: ~4.1), suggesting its more metallic behavior and higher electronic conductivity, which will effectively improve the electron transfer capability of the composites. To investigate the effect of Fe doping amount on lithium storage performances, the electrochemical performance of FexCo1–xSe2@NC@rGO composites with various Fe/Co atomic ratios has also been evaluated. As shown in Figure 6c and Figure S16a, with increasing the designed Fe/Co atomic ratio from 0:1 to 3:7, the cycle performance is gradually improved, and the FexCo1-xSe2@NC@rGO (x=0.3) electrode has the highest discharge capacity (947 mAh g-1) after 500 cycles at 1.0 A g–1. However, with further increasing the designed Fe/Co atomic ratio, the cycle performance of FexCo1xSe2@NC@rGO

becomes worse, e.g., only capacities of 595 mAh g1 and 500 mAh g1

for FexCo1-xSe2@NC@rGO (x=0.5) and FeSe2@NC@rGO could be retained after 500 cycles at 1.0 A g–1, respectively. Similar change tendency could be also observed by comparing the rate performance of FexCo1-xSe2@NC@rGO composites with various Fe/Co atomic ratios (Figure 6d and Figure S16b), from which the FexCo1-xSe2@NC@rGO (x=0.3) has the best rate capability. Such results indicate that the amount of Fe-doping has a vital effect on the electrochemical performance of CoSe2. The FexCo1-xSe2@NC@rGO (x=0.3) has the best electrochemical performance, which might be due to its optimum balance among chemical composition, phase structure and surface property. CONCLUSION In summary, cubic-to-orthorhombic phase transformation in dual carbon confined CoSe2 induced by Fe-doping have been reailized and employed in designing anodes towards lithium-ion batteries. Benefiting from structural phase-transition enginering that enables 16


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higher electrical conductivity, enhanced ion/electron transfer kinetics and more active sites for Li+ ion storage, the o-FexCo1–xSe2@NC@rGO anode exhibits dramatically improved lithium storage performance (920 mAh g−1 after 1000 cycles at 1.0 A g−1 and 412 mAh g−1 after 1000 cycles at 10.0 A g−1) in comparison to c-CoSe2@NC@rGO anode in lithiumion batteries. We hope that the explored strategy and findings can provide more insight for developing high-performance electrodes for energy storage and conversion. MATERIALS AND METHODS Synthesis and Purification of Graphene Oxide (GO). GO with high purity was synthesized by the oxidation reaction of graphite powder according to the modified Hummers’ method.60 Synthesis of Fe-Co-PBA@rGO, Fe-PB@rGO and Co-PBA@rGO Aerogel. In a typical synthesis, 0.1 mL of 0.5 mol L−1 Co(NO3)2, 0.1 mL of 0.5 mol L−1 Fe(NO3)3, and 0.15 mL of 0.5 mol L−1 sodium L-ascorbate (or VC-Na) were well dispersed in 2 mL of 2 mg mL−1 graphene oxide suspension, during which the Fe3+ ions were reduced to Fe2+. Then, the mixed solution was centrifuged, and precipitates were well dispersed in 1 mL deionized water, followed by adding 0.15 mL of 0.5 mol L−1 K3[Co(CN)6] in the solution. After standing, the sediments were collected by centrifugation. To induce the selfassembly of graphene oxide, 0.1 mL of 0.5 mol L−1 VC-Na was added into above sediment, which was well dispersed in 1 mL deionized water. The 3D Fe-Co-PBA@rGO hydrogel was firstly achieved by taking out the above mixture solution into glass bottle, kept in the oven for 95 °C for 2 h. Finally, the as-prepared Fe-Co-PBA@rGO hydrogel was purified with the 50 °C deionized water and performed a freeze-drying for 24 h to obtain the 3D Fe-Co-PBA@rGO aerogel. The synthesis of Co-PBA@rGO aerogel was similar to that makes Fe-Co-PBA@rGO except that only Co(NO3)2 was used. By change the amounts of Co(NO3)2 and Fe(NO3)3, Fe-Co-PBA@rGO with different Fe/Co ratios can be obtained. 17


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The Fe-PB@rGO was synthesized by similar procedure except using Fe(NO3)3 and K3[Fe(CN)6]. Synthesis of Fe-Co-PBA and Co-PBA Aerogel. Typically, 0.003 mol FeCl3∙6H2O and 0.003 mol CoCl2∙6H2O was dissolved in 15 mL deionized water before adding 15 mL deionized water with 0.004 mol K3[Co(CN)6]. After standing for 2 h, the mixed solution was centrifuged and parched in an oven of 75 °C for 12 h. Co-PBA was synthesized using the similar strategy except that only CoCl2∙6H2O was used. Synthesis xSe2@NC

of

o-FexCo1–xSe2@NC@rGO,

c-CoSe2@NC@rGO,

o-FexCo1–

and c-CoSe2@NC Composite. The o-FexCo1–xSe2@NC@rGO was synthesized

through a thermal-induced carbonization and selenization process. Typically, the Fe-CoPBA@rGO precursors and selenium powders were heated up to 600 °C and maintained at this temperature for 2 h under a flowing argon gas. c-CoSe2@NC@rGO, FexCo1– xSe2@NC@rGO

with various Fe/Co atomic ratios, o-FexCo1–xSe2@NC and c-CoSe2@NC

were synthesized using the similar strategy except that Co-PBA@rGO, Fe-Co-PBA@rGO with different Fe/Co ratios, Fe-Co-PBA and/or Co-PBA were used as precursor, respectively. Synthesis of FexCo1–xSe2 and CoSe2 Composite. The bare FexCo1–xSe2 composite can be synthesized by reacting Co3O4 and Fe3O4 with Se powders. Typically, 580 mg of Co(NO3)2∙6H2O and 800 mg Fe(NO3)3∙9H2O were mixed by recrystallization and then oxidized in the air at 350 °C for 4 h. After that the Fe3xCo3–3xO4 and selenium powders were heated up to 350 °C for 10 h under a flowing mixed Ar flow with 10% H2. The synthesis of CoSe2 was similar to FexCo1–xSe2 composite except using Co(NO3)2∙6H2O only. Synthesis of N-doped Carbon (NC) Powder and rGO Gel Composite. The NC powders were achieved by using an acid treatment towards o-FexCo1–xSe2@NC 18


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composites, during which the selenide could be etched in dilute nitric acid (HNO3). The rGO gel was obtained by one-pot hydrothermal reaction of a mixture aqueous solution consisted of GO and VC-Na. Characterization. Powder X-ray diffraction (XRD) measurements were conducted on a powder X-ray diffraction system (XRD, D8 ADVANCE X-ray diffractometer, which equipped with Cu Kα radiation, λ = 1.5406 Å). The phase structure and components were further analyzed by the RIETAN-2000 program based on the Rietveld method.61 The morphology and microstructure of the composite were investigated by employing the Field emission scanning electron microscope (FESEM, ZEISS, Ultra55) and the transmission electron microscope (TEM, JEOL JEM-2100F). The chemical state of the elements in the materials was identified by utilizing the X-ray photoelectron spectroscopy (XPS, Kratos XSAM-800, which equipped with a monochromatic Al Kα1 X-ray, hν=1486.6 eV). The chemical composition of samples was identified by the inductively coupled plasma optical emission spectrometry (ICP, Perkin-Elmer Optima 5300). To identify the carbon structure, Raman spectroscopy measurement was carried out on Renishaw via plus laser Raman spectrometer (RENISHAW). The pore size volume and specific surface area of the composites were carried out on Autosorb IQ Gas Sorption System at 77 K. Thermogravimetric analysis (TGA) was tested with a Mettler Toledo TGA under the 20 mL min-1 air flow with the heating rate of 20 °C min–1. Electrochemical measurements. Electrochemical performance was tested at room temperature by using CR2032 coin-type half-cells. The coin-type cells, in which working electrodes are free-standing and binder-free, were assembled in a glove box filled with argon, using lithium foil as the counter/reference electrode. The working electrode of the powder was prepared by mixing the synthesized active materials and polyvinylidene fluoride (PVDF) in a weight ratio of 4:1 using the N-methyl-2-pyrrolidinone (NMP) as 19


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solvent. Celgard-2400 was used as the separator, and commercial LiPF6 (1 mol L‒1) in ethylene carbonate (EC) and dimethyl carbonate (DMC) (the volume ratio of EC and DMC was 1:1) was used as the electrolyte. The galvanostatic charge-discharge performance was checked using a LAND CT2001 testing system in the voltage range from 0.01 to 3.00 V. Cyclic voltammetry (CV) was carried out on an electrochemical workstation (AutoLab PGSTAT302N) at a scan rate of 0.11.0 mV s‒1. Electrochemical impedance spectroscopy (EIS) was conducted on the electrochemical measurement system (AutoLab PGSTAT302N) in the frequency from 100 kHz to 0.01 Hz. Theoretical calculation. We used Vienna ab initio simulation package (VASP)62 to implement density functional theory (DFT)63 calculations with the projector augmented wave (PAW) method.64 The exchange and correlation functional used in our calculations was Perdew-Burke-Ernzerhoffunctional (PBE).65 The plane-wave energy cutoff was set at 520 eV and convergence criterion for geometric optimization was 10‒6 eV. MonkhorstPack66 generated 11 × 11 × 11 k-point grid was used to sample the Brillouin zone for cubic lattice relaxation, for orthorhombic pristine CoSe2 and Fe-doped CoSe2, 13 × 11 × 9 kpoint grid was used. To obtain accurate density of states (DOS), fine k-point grids with 6 × 6 × 6 were used. ASSOCIATED CONTENT Supporting Information. This material is available free of charge via the Internet at http://pubs.acs.org. Refined XRD data, comparison of electrochemical performance, additional XRD patterns, the sketch map of phase transition, Raman spectra, TGA curves, additional FESEM images, ex-situ XRD patterns, and electrochemical measurements (PDF). AUTHOR INFORMATION 20


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Corresponding Author *E-mail: [email protected] (R. Wu) Author Contributions &Y.

Liu and Z. Chen contributed equally to this work.

Notes The authors declare no competing financial interest. ACKNOWLEDGEMENTS This work was financially supported by the National Natural Science Foundation of China (Grant Nos. 51672049, 51871060 and 51831009), the China Postdoctoral Science Foundation (Grant No. 2018M640337), CURE (Hui-Chun Chin and Tsung-Dao Lee Chinese Undergraduate Research Endowment) (No. 18928), and National University Student Innovation Program (No. 201810246085).

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Framework Assembled by Cobblestone-Like CoSe2@C Nanospheres for Ultrastable Sodium-Ion Storage. ACS Appl. Mater. Interfaces 2018, 10, 14716– 14726. (45) Zhu, Y.; Huang, Z.; Hu, Z.; Xi, L.; Ji, X.; Liu, Y. 3D Interconnected Ultrathin Cobalt Selenide Nanosheets as Cathode Materials for Hybrid Supercapacitors. Electrochimica Acta 2018, 269, 30–37. (46) Wei, Y.; Huang, L.; He, J.; Guo, Y.; Qin, R.; Li, H.; Zhai, T. Healable Structure Triggered by Thermal/Electrochemical Force in Layered GeSe2 for High Performance Li-Ion Batteries. Adv. Energy Mater. 2018, 8, 1703635. (47) Wang, D.; Yu, Y.; He, H.; Wang, J.; Zhou, W.; Abruña, H. D. Template-Free Synthesis of Hollow-Structured Co3O4 Nanoparticles as High-Performance Anodes for Lithium-Ion Batteries. ACS Nano 2015, 9, 1775–1781. (48) Yu, J.; Chen, S.; Hao, W.; Zhang, S. Fibrous-Root-Inspired Design and Lithium Storage Applications of a Co–Zn Binary Synergistic Nanoarray System. ACS Nano 2016, 10, 2500–2508 (49) Chen, J.; Mao, Z.; Zhang, L.; Wang, D.; Xu, R.; Bie, L.; Fahlman, B. D. NitrogenDeficient Graphitic Carbon Nitride with Enhanced Performance for Lithium Ion Battery Anodes. ACS Nano 2017, 11, 12650–12657. (50) Ali, Z.; Asif, M.; Huang, X.; Tang, T.; Hou, Y. Hierarchically Porous Fe2CoSe4 Binary-Metal Selenide for Extraordinary Rate Performance and Durable Anode of Sodium-Ion Batteries. Adv. Mater. 2018, 30, 1802745. (51) Sun, H.; Xin, G.; Hu, T.; Yu, M.; Shao, D.; Sun, X.; Lian, J. High-Rate LithiationInduced Reactivation of Mesoporous Hollow Spheres for Long-Lived LithiumIon Batteries. Nat. Commun. 2014, 5, 4526. (52) Jin, J.; Zheng, Y.; Kong, L. B.; Srikanth, N.; Yan, Q.; Zhou, K. Tuning ZnSe/CoSe in MOF-Derived N-Doped Porous Carbon/CNTs for HighPerformance Lithium Storage. J. Mater. Chem. A 2018, 6, 15710–15717. (53) Wang, X.; Tang, Y.; Shi, P.; Fan, J.; Xu, Q.; Min, Y. Self-Evaporating from inside to Outside to Construct Cobalt Oxide Nanoparticles-Embedded Nitrogen-Doped Porous Carbon Nanofibers for High-Performance Lithium Ion Batteries. Chem. 27


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Figure 1. Schematic illustration for the synthesis of o-FexCo1–xSe2@NC@rGO aerogel.

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Figure 2. Rietveld refinement for the XRD patterns of the selenization products using (a) Fe-Co-PBA@rGO and (b) Co-PBA@rGO precursors; (c) schematic illustration of Fe-doping induced cubic-to-orthohombic phase transition in CoSe2; (d) Raman spectra of the selenization products using Fe-Co-PBA@rGO and Co-PBA@rGO; (e) specific surface area of the selenization products of Fe-Co-PBA @rGO and the inset showing the pore size distribution.

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Figure 3. FESEM images of (a‒c) Fe-Co-PBA@rGO precursor and (d‒h) its selenization products; (i) the representative FESEM image of the selenization products using Fe-Co-PBA@rGO precursor and its corresponding (j) Co, (k) Fe, (l) Se, (m) C, and (n) N element mapping. Inset in Figure 3d: the digital image of the free-standing Fe-Co-PBA@rGO aerogel.

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Figure 4. (a‒c) TEM images and (d, e) high-resolution TEM images of o-FexCo1– xSe2@NC@rGO;

(f) schematic illustration for the dual carbon confined o-FexCo1–xSe2

nanoparticles.

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Figure 5. (a) Survey spectrum of o-FexCo1–xSe2@NC@rGO aerogel and the corresponding high-resolution (b) Co 2p, (c) Fe 2p, (d) Se 3d, (e) C 1s and (f) N 1s XPS spectra.

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Figure 6. Electrochemical properties of the o-FexCo1–xSe2@NC@rGO electrode: (a) CV curves at a scan rate of 0.1 mV s‒1, (b) initial three discharge-charge profiles at a current density of 0.2 A g‒1, (c) cycling performance at 1.0 A g‒1, (d) rate capability at current densities between 0.2 to 12.8 A g‒1, and (e) comparison of the high-rate cycling performance of the o-FexCo1–xSe2@NC@rGO electrode with other reported Co-based anodes of LIBs.

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Figure 7. (a) CVs at different scan rates for o-FexCo1–xSe2@NC@rGO electrodes after 120 cycles; and their derived (b) log i vs log v plots; (c) capacitive (light orange) and diffusion-controlled (light bean green) contribution at 0.4 mV s–1; (d) normalized contribution ratios of capacitive (light orange) and diffusion-controlled (light green) capacities at different scan rates.

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Figure 8. The charge density distributions for (a) c-CoSe2, (b) o-CoSe2, and (c) Fedoped o-CoSe2; the total density of sates for (d) c-CoSe2, (e) o-CoSe2, and (f) Fe-doped o-CoSe2, the Fermi level is defined as zero. The atoms with purple, green and yellow colors represent the Co, Se, and Fe atoms, respectively.

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Iron-Doping-Induced Phase Transformation in Dual-Carbon-Confined

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