Achieving Fully Reversible Conversion in MoO3 for Lithium Ion

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Achieving Fully Reversible Conversion in MoO3 for Lithium Ion Batteries by Rational Introduction of CoMoO4 Wei Wang,† Jinwen Qin,† Zhigang Yin, and Minhua Cao*

Key Laboratory of Cluster Science, Ministry of Education of China, Beijing Key Laboratory of Photoelectronic/Electrophotonic Conversion Materials, School of Chemistry and Chemical Engineering, Beijing Institute of Technology, Beijing 100081, P. R. China S Supporting Information *

ABSTRACT: Electrode materials based on conversion reactions with lithium ions generally show much higher energy density. One of the main challenges in the design of these electrode materials is to improve initial Coulombic efficiency and alleviate the volume changes during the lithiation−delithiation processes. Here, we achieve fully reversible conversion in MoO3 as an anode for lithium ion batteries by the hybridization of CoMoO4. The porous MoO3−CoMoO4 microspheres are constructed by homogeneously dispersed MoO3 and CoMoO4 subunits and their lithiation/delithiation processes were studied by ex situ TEM to reveal the mechanism of the reversible conversion reaction. Co nanoparticles are in situ formed from CoMoO4 during the lithiation process, which then act as the catalyst to guarantee the reversible decomposition of Li2O, thus effectively improving the reversible specific capacity and initial Coulombic efficiency. Moreover, the pores in MoO3−CoMoO4 microspheres also greatly enhance their mechanical strength and provide enough cavity to alleviate volume changes during repeated cycling. Such a design concept makes MoO3 to be a potential promising anode in practical applications. The full cell (LiFePO4 cathode/MoO3−CoMoO4 anode) displays a high capacity up to 155.7 mAh g−1 at 0.1 C and an initial Coulombic efficiency as high as 97.35%. This work provides impetus for further development in electrochemical charge storage devices. KEYWORDS: molybdenum oxide, hybrid materials, pores, catalysis, lithium ion batteries studies on MoO3 for LIBs mainly focus on solving the first problem by reducing particle size to nanoscale or hybrizing MoO3 with carbon-based materials, and although great advances have been made,9 the resultant capacities are still not satisfactory.10 In order to break through the capacity barrier, making the conversion process of Li2O reversible is only way to achieve high capacity of MoO3. In view of the limitation of the electrochemical reaction, the reversible transformation of Li2O remains a great challenge. A fully reversible conversion reaction for Li2O was first reported by Tarascon et al. in nanosized CoO as an anode material for LIBs.6 They demonstrate the electrochemical process of CoO by ex situ TEM, and the specific mechanism of the reversible conversion reaction is as follows. During the lithiation process, CoO is reduced into metal Co nanoparticles (NPs) along with the formation of Li2O crystals. In the subsequent delithiation process, the formed Co NPs break

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ecently, MoO3, as a well-known lithium insertion material, has been intensively studied as an anode for lithium ion batteries (LIBs) because of its good chemical stability, high theoretical specific capacity (1111 mAh g−1), and very stable one-dimensional layered structure.1−5 The layered structure of MoO3 is capable of acting as a temporary support for intercalated lithium ions.2 However, as a high-specific-capacity anode material for LIBs, it does not follow the mechanism of insertion/deinsertion of lithium ions in layered graphite electrode, but the so-called conversion reaction (typically operating in the potential range of 0−3.0 V vs Li+/Li).1 Therefore, like most of transition-metal oxides with high theoretical specific capacities, MoO3 also faces two serious problems as an anode for LIBs, which greatly limit its practical applications.4 MoO3 shows significant capacity fading upon extended discharge/charge cycling due to its huge volume changes during the lithiation and delithiation process, thereby resulting in the pulverization of electrode materials accompanied by poor cycling performance.5 The other big disadvantage typically is its low Coulombic efficiencies in initial cycles, which is mainly caused by the formation of a large number of irreversible Li2O.4,6−8 Until now, the reported © 2016 American Chemical Society

Received: August 1, 2016 Accepted: November 3, 2016 Published: November 4, 2016 10106

DOI: 10.1021/acsnano.6b05150 ACS Nano 2016, 10, 10106−10116

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Figure 1. (a) Schematic diagram illustrating the procedure to fabricate P−Mo−Co−HMSs; (b) The schematic representation of the Li2O decomposition, facile strain relaxation as well as continuous electron/ion pathways during the lithiation-delithiation processes in P−Mo−Co− HMSs.

relieving the capacity fading and improving the cycling stability of the MoO3 electrode. On the other hand, constructing porous micro-/nanostructures assembled from nanosized subunits is an effective way to improve the electrochemical performance of electrode materials in view of their enhanced mechanical/structural integrity against large changes in volume and crystal structure over extended cycling.16,17 As well established previously, this kind of robust porous micro-/nanostructures not only possess the outstanding advantages of nanosized materials, such as the high electrode/electrolyte contact area and the shortened transport length for Li+ and e−, but also guarantee good structure stability during the electrochemical process due to their desirable mechanical strength properties.18,19 Moreover, the porous structure can provide additional free volume to alleviate the structural strain for active materials during cycling.20−23 Herein, we report the rational design and synthesis of porous MoO3−CoMoO4 hybrid microspheres (P−Mo−Co−HMSs) assembled from uniform NPs subunits obtained via a facile solvothermal method followed by thermal treatment. Benefiting from its appealing component and structural features, the P− Mo−Co−HMSs anode exhibits extremely excellent lithium storage performance for LIBs, such as high reversible capacity (1598.7 mAh g−1 at 0.2 A g−1 after 100 cycles), superior rate capability (640 mA h g−1 at 3.0 A g−1), and excellent cycling stability (875.4 mAh g−1 at 1.0 A for 100 cycles). We further construct a LiFePO4/P−Mo−Co−HMSs full cell, which displays high reversible capacities of 155.7 mAh g−1 at 0.1 C and 90.75 mAh g−1 at 1.0 C after 70 cycles. The detailed studies on the electrochemical lithiation and delithiation processes reveal that the in situ formed Co nanoclusters from CoMoO4 can act as a catalyst to promote the fully reversible conversion

down into even smaller clusters, which then react with the oxygen anion, leading to the decomposition of Li 2O accompanied by the formation of CoO (CoO + 2Li+ ⇆ Li2O + Co).7 As a consequence, these two steps will lead to a fully reversible conversion of Li2O, release extra Li+, and eventually offer an additional capacity.8 Subsequently, similar phenomena have also been observed in nanosized Co@carbon, CoO quantum dots/graphene, CoO/Li2O, and Co3O4/graphene composites.11−14 In light of these findings, we seek to hybridize MoO3 with highly active nanosized Co-based oxides in an atomically homogeneous manner, which is expected to be able to achieve the fully reversible conversion of Li2O, thus overcoming the second problem mentioned above and greatly enhancing the final reversible capacity. Among various Cobased materials, ternary CoMoO4 has been regarded as an ideal candidate due to its synergistically improved electrochemical properties including electrical/ionic conductivity, mechanical stability, and high specific capacity (980 mAh g−1).15 The introduction of CoMoO4 in MoO3 is believed to hold promising prospects for MoO3 as an anode material because the high content of Mo in the constructed MoO3−CoMoO4 hybrid can provide super capacity via the conversion reaction mechanism. In addition, the in situ formed Co nanoclusters from CoMoO4 during the electrochemical process can promote the formation/decomposition of Li2O, thus providing extra capacity. Furthermore, the in situ formed Co nanoclusters can also be a good buffer matrix for the MoO3 electrode. Specifically, the Co nanocluster matrix can not only prevent the aggregation of Mo nanocrystals into large grains but also function as a cushion to buffer volume change and structural stress during the discharge/charge process, thus effectively 10107

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shakeup-type peaks of Co at the high binding energy side of the Co 2p3/2 and Co 2p1/2 edge.25,26 Finally, the O 1s signal shows four types of oxygen with different chemical states, which appear at 529.8 (O2−), 530.4 (OH−), 531.4 (C−O, O−CO), and 532.7 eV (H2O), respectively (Figure 2d).25 Parts a−f of Figure 3 show field-emission scanning electron microscopy (FESEM) images of the Mo−Co precursor and its calcined sample (P−Mo−Co−HMSs). It can be seen from the low-magnification FESEM image that the Mo−Co precursor is composed of well-defined MSs with nonuniform size distribution, and most of the MSs have a size of 2.0 μm in diameter (Figure 3a and its inset). The high-magnification FESEM images reveal that each MS has a smooth surface (Figure 3b,c). Besides, on the basis of the time-dependent experimental observations, a stepwise self-assembly mechanism is proposed for understanding the formation process of the well-defined Mo−Co precursor MSs (Figure S4). After thermal treatment in air, the Mo−Co precursor subjects to decomposing and oxidizing to form P−Mo−Co−HMSs, which completely maintain the spherical morphology of the precursor. P−Mo−Co−HMSs have a rough surface instead of the smooth surface, and meanwhile, the sizes of the P−Mo−Co−HMSs are a little larger than those of the Mo−Co precursor, with an average size of 2.36 μm in diameter (Figure 3d and its inset). Close inspection for the MS surface reveals that the surface possesses an obvious pore structure (Figure 3e,f), which may result from the release of some gases such as CO, CO2, or H2O during the thermal treatment process. Like the precursor, the P−Mo−Co−HMSs also have a nonuniform size distribution, and this feature is particularly beneficial for improving their tap density when used as active materials for LIBs. The low- and high-magnification transmission electron microscope (TEM) images (Figure 3g,h) further disclose that the MSs are assembled from NPs with sizes in the range of 10−30 nm. Meanwhile, the combustion of organic components in the Mo− Co precursor results in the formation of the pores, which extend from the outer of the MSs to the inner (Figure 3h). The corresponding selected area electron diffraction (SAED) pattern of an individual P−Mo−Co−HMS (the inset in Figure 3h) gives its polycrystalline nature. Although the content of CoMoO4 in this hybrid is low, we still can distinguish CoMoO4 from MoO3 in high-resolution transmission electron microscope (HRTEM) image. As shown in Figure 3i, the distinct lattice spacings of 0.69 and 0.349 nm are clearly observed. The former can be indexed to the (020) plane of MoO3 (JCPDS card no. 05-0508) and the latter to (201) plane of CoMoO4 (JCPDS card no. 21-0868). In addition, the energy-dispersive X-ray spectrum (EDS) and elemental mapping images further demonstrate the existence and homogeneous distribution of Mo, Co, and O elements in the P−Mo−Co−HMSs (Figure 3j,k). The pore nature of the P−Mo−Co−HMSs is further determined by N2 adsorption/desorption measurements. As shown in Figure 3l, the isotherms have a steep increase of sorption at relatively high pressures (P/P0 > 0.8), and a hysteresis loop, although less obviously, belongs to the type-IV hysteresis, implying that the P−Mo−Co−HMSs contain macropores and mesopores resulting from the combustion of organic components.17,18 The sizes of the macro- and mesopores based on desorption data are mainly in the range of 25− 75 nm, in which the pore amount at 44.4 nm is maximum (the inset in Figure 3l). The calculated Brunauer−Emmett−Teller (BET) surface area of the P−Mo−Co−HMSs is evaluated to be 12.7 m2 g−1. This porous structure is likely to create more active

of Li2O, which is mainly responsible for the significantly improved electrochemical performance. The concept presented in current work can be readily extended to other high-capacity electrode materials and provides a strategy to design and synthesize future electrode materials for high-performance LIBs.

RESULTS AND DISCUSSION The synthesis procedure of P−Mo−Co−HMS is schematically illustrated in Figure 1a, which involves two steps. First, an amorphous Mo−Co precursor was obtained via the solvothermal reaction of MoO2(acac)2 and CoCl2·6H2O in ethanol solvent at 200 °C for 48 h (Figure S1). Then, the as-resultant Mo−Co precursor was calcined at 500 °C for 3 h in air, thus leading to the formation of P−Mo−Co−HMSs accompanied by the release of a large number of gases. The determination of the calcination temperature was based on the thermogravimetric and differential scanning calorimetry (TG-DSC) analysis of the amorphous Mo−Co precursor (Figure S2). The phase composition of the final sample was first characterized by powder X-ray diffraction (XRD) and X-ray photoelectron spectra (XPS) measurements. As shown in Figure 2a, those

Figure 2. (a) XRD pattern of the as-obtained P−Mo−Co−HMSs. High resolution XPS spectra of the P−Mo−Co−HMSs: (b) Mo 3d, (c) Co 2p, and (d) O 1s.

diffraction peaks marked with ∗ in the XRD pattern can be indexed to MoO3 (JCPDS card no. 05-0508) and the ones marked with # can be assigned to CoMoO4 (JCPDS card no. 21-0868). No other crystalline phases are detected. Moreover, the mass ratio of MoO3 to CoMoO4 is determined to be 88.7:11.3 by inductively coupled plasma optical emission spectroscopy (ICP-OES) analysis. To further determine the composition of the sample, XPS measurements are performed. The survey XPS spectrum clearly indicates the presence of Mo, Co, and O elements (Figure S3), consistent with the above XRD result. The high-resolution Mo 3d XPS spectrum (Figure 2b) shows two strong peaks centered at 232.3 and 235.5 eV, which could be attributed to Mo 3d5/2 and Mo 3d3/2 of Mo6+, respectively,24 indicating the presence of only the Mo6+ oxidation state. For the Co 2p XPS spectrum (Figure 2c), two major peaks with binding energies at 784.2 eV (Co 2p3/2) and 800.7 eV (Co 2p1/2) are the characteristic of Co2+, while the satellite peaks at around 789.3 and 807.2 eV are two 10108

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Figure 3. (a−c) FESEM images of the Mo−Co precursor; (d−f) FESEM images of the P−Mo−Co−HMSs; (g−i) TEM images, SAED pattern, and HRTEM image of the P−Mo−Co−HMSs; (j,k) EDS and elemental mapping images recorded on an individual P−Mo−Co−HMSs. (l) N2 adsorption−desorption isotherms and Barrett−Joyner−Halenda (BJH) pore size distribution of P−Mo−Co−HMSs. The insets in (a) and (d) are particle size distributions of the Mo−Co precursor and the P−Mo−Co−HMSs, respectively.

MoO3 MSs (Figure 4d,e).30,31 In the anodic scans, besides the oxidation peaks at ca. 0.5, 1.32, and 1.8 V that correspond to the multistep phase transition process of MoO3,27,32 an obvious peak at ca. 1.50 V is detected, which is related to the decomposition of Li2O (Figure 4c).31 Clearly, this peak did not appear in the CV profiles of bare MoO3 MSs (Figure 4d−f) and CoMoO4 NPs (Figure 4g−i). As we know, Li2O is primarily irreversible and is almost impossible to decompose under unprompted conditions, which will eventually cause the irreversible capacity losses. However, in P−Mo−Co−HMS electrode, the Co NPs in situ formed during the Li + intercalation process act as a catalyst to prompt the reduction of Li2O.11−14 Thus, the introduction of CoMoO4 in MoO3 plays a strategic point to realize the reversible reduction of Li2O, which is expected to significantly improve the reversible capacity of MoO3 electrode. Furthermore, different oxidation/ reduction pairs of 0.21/0.5, 0.8/1.32, and 1.49/1.8 V are highly reversible after the first cycle, which can be attributed to the Li+ insertion/extraction with different site energies and the phase transition process of MoO3. In addition, the cathodic/anodic peaks at 1.49/1.8 V are also related to the oxidation and reduction between metallic Co and CoO. On the basis of the above analysis and the literatures,27,29 the reversible electrochemical reaction process of Li with the P−Mo−Co−HMSs can be summarized as follows:

sites, facilitate the transport of lithium ions, and enlarge the contact of electrolyte with active material during the electrochemical reactions. The lithium storage performance of the resultant product is evaluated as an anode material for LIBs. For comparison, bare MoO3 MSs and CoMoO4 NPs were also prepared (Figures S5−S7). Parts a−i of Figure 4 show cyclic voltammetry (CV) profiles of P−Mo−Co−HMSs, bare MoO3 MSs, and CoMoO4 NPs for the first five cycles at a scan rate of 0.2 mV s−1. In the first cathodic process, three reduction peaks can be observed at 2.22, 1.64, and 0.55 V for P−Mo−Co−HMSs (Figure 4a,b). The broad peak centered at 2.22 V corresponds to the intercalation process of Li+ into the crystalline MoO3 to form LixMoO3, and the minor peak located at 1.64 V can be attributed to the decomposition of CoMoO4 and the reduction of Mo6+ to Mo4+ (Figure 4b),27 whereas the intense peak below 0.55 V can be assigned to the reaction between lithium and the as-formed LixMoO3, the further reduction of Co2+ and Mo4+ to metallic Co and Mo as well as the formation of solid electrolyte interface (SEI) (Figure 4a).28,29 After the first cycle, four new reduction peaks at ca. 1.49, 1.2, 0.8, and 0.21 V are observed. The peak at 1.49 V corresponds to lithium insertion into LixMoO3, and the peaks at 0.8 and 0.21 V are commonly recognized as the conversion of LixMoO3 into metallic Mo and Li2O, while the other peak at 1.2 V represents the formation of Li2O (Figure 4b), which has also been detected for the bare 10109

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Figure 4. CV curves and corresponding enlarged curves of P−Mo−Co−HMSs (a−c), MoO3 MSs (d−f), and CoMoO4 NPs (g−i).

Figure 5. (a, b) Discharge−charge voltage profiles and cycling performance of P−Mo−Co−HMSs, bare MoO3 MSs, and CoMoO4 NPs at a current density of 0.2 A g−1 as well as the corresponding CE of P−Mo−Co−HMSs. (c) Rate performance of the three samples at current densities from 0.2 to 3.0 A g−1. (d) EIS of the three samples tested after 100 cycles at a current density of 0.2 A g−1. (e) Cycling performance of P−Mo−Co−HMS electrode at different current densities. (f) FESEM image of P−Mo−Co−HMS electrode after 100 cycles at a current density of 0.2 A g−1.

MoO3 + x Li+ + x e− → LixMoO3

(1)

CoMoO4 + 8Li+ + 8e− → Co + Mo + 4Li 2O

(2)

voltage plateau at ca. 2.2 V and a sloping voltage plateau located below 0.5 V for both P−Mo−Co−HMSs and MoO3 MSs are quite visible, consistent with their CV results. To be specific, the initial discharge curve for P−Mo−Co−HMSs (Figure S8a) can be divided into two regions: (I) above 1.5 V and (II) below 1.5 V (vs Li+/Li). In region I, the potential plateau at around 2.2 V corresponds to the Li+ insertion into the crystalline MoO3 to form LixMoO3 (eq 1), providing a capacity of 188.8 mAh g−1.36,37 As the voltage drops to below 1.5 V (region II), the conversion reaction of the LixMoO3 with Li+ to form metallic Mo and Li2O takes place, giving rise to a long plateau at around 0.4−0.5 V (eq 3) and contributing a capacity as high as 1503.4 mAh g−1.38 For the same region, the capacities for P−

LixMoO3 + (6 − x)Li+ + (6 − x)e− ↔ Mo + 3Li 2O (3)

Co + Li 2O ↔ CoO + 2Li+ + 2e−

(4)

Figure 5a displays the discharge/charge voltage profiles of the P−Mo−Co−HMSs, bare MoO3 MSs, and CoMoO4 NPs at a current density of 0.2 A g−1 over a voltage range from 0.01 to 3.0 V. In the initial discharge profile, an obvious extending 10110

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ACS Nano Table 1. Comparison of the Initial CE of Present Work with Reported MoO3-Based Materials materials

current density (A g−1)

discharge capacity (mAh g−1)

charge capacity (mAh g−1)

initial CE (%)

year

ref

MoO3/CNT MoO3 HfO2-coated MoO3 α-Fe2O3@α-MoO3 composites yolk-shell MoO3 microspheres MoO3 nanowires one-dimensional α-MoO3 ultralong α-MoO3nanobelts MoO2@C MoO3-carbon P−Mo−Co−HMSs MoO3 MSs CoMoO4 NPs

0.1 0.1 0.1 0.1 1.0 0.2 0.1 0.2 0.1 0.15 0.2 0.2 0.2

1604.5 1685.4 1728 1899 1481 1988 800 2349 1278.8 945 1692.2 934.5 1015.5

1138.3 1028.3 1120 1416 998 1078 200 1126 769.3 813 1328.9 1140.0 1373.0

70.9 61 64.8 74.5 67.4 54.2 25 47.9 60 86 78.5 81.9 73.9

2015 2015 2015 2014 2014 2014 2013 2012 2012 2011

33 33 34 35 36 37 38 39 40 1 this work

is 117.7% of the original value. The additional capacity can be recognized as being related to the partial amorphization of MoO3 during cycling,11 which improves Li diffusion kinetics, creates more accessible active sites for Li+ insertion, and thus expresses the gradual increase of the capacity.12 A similar cycling-induced capacity increase has also been observed for the bare MoO3 MS electrode, which maintains the specific capacity of 1036.6 mAh g−1 after 100 cycles (Figure 5b). The mass percentage of CoMoO 4 in our P−Mo−Co−HMSs is determined to be 11.3 wt % by ICP analysis, which has been optimized. For comparison, we also examined the cycling performances of P−Mo−Co−HMSs with other mass percentages (9.6 and 13.0 wt %) of CoMoO4 under the same tested conditions (Figure S9). Obviously, they both show fast capacity fading, and their capacity retentions are as low as 51.8% and 41.7%, respectively, far lower than that of P−Mo−Co−HMSs. The rate performance was further investigated at different cycling rates (Figure 5c). Clearly, P−Mo−Co−HMSs exhibit super rate performance, and average discharge capacities of 1342.6, 1153.3, 1011.8, 883.7, and 732.8 mAh g−1 can be obtained when the current densities increase stepwise to 0.2, 0.4, 1.0, 1.5, and 2.0 A g−1, respectively. Even at a high current density of 3.0 A g−1, P−Mo−Co−HMSs still deliver a capacity of more than 640 mAh g−1, which is almost two times higher than the theoretical capacity of the graphite. After undergoing the high rate testing for 60 cycles, P−Mo−Co−HMSs display a capacity as high as 1244 mAh g−1 when the current density rolls back to 0.2 A g−1. For bare MoO3 MSs, the capacities are around 838.2, 741.8, 593.1, 561.5, 491.6, and 380.2 mAh g−1 at 0.2, 0.4, 1.0, 1.5, 2.0, and 3.0 A g−1, respectively, which are much lower than those of P−Mo−Co−HMSs, while for CoMoO4 NPs, the capacities are only 195 and 71 mAh g−1 at 1.0 and 3.0 A g−1, respectively. The significantly enhanced rate performance of P−Mo−Co−HMSs may be closely related to their excellent conductivity, which has been further confirmed by electrochemical impedance spectroscopy (EIS) analysis. Figure 5d shows Nyquist plots of the three samples over the frequency range from 100 kHz to 0.01 Hz after 100 cycles at a current density of 0.2 A g−1. As can be seen, the Nyquist plots are composed of one depressed semicircle followed by a slope line. The semicircle represents charge-transfer impedance, while the slope line represents the solid-state diffusion impedance.25 Obviously, P−Mo−Co−HMSs have smaller charge-transfer impedance compared to bare MoO3 MSs and CoMoO4 NPs, indicating the improved conductivity and enhanced reaction kinetics of for P−Mo−Co−HMSs. Moreover, P−Mo−Co−

Mo−Co−HMSs are much higher than those of MoO3 MSs. The enhanced capacities derive mainly from the redox reactions occurring between Li+ and CoMoO4,31 which has been confirmed by the smaller slope at ca. 1.6 V and the much longer plateau between 0.01 and 0.5 V for P−Mo−Co−HMSs (eq 2, Figure S8a). At the end of the conversion reaction, the structure of MoO3 transforms into an amorphous state, resulting in smooth voltage profiles for the first charge process and subsequent cycles. The initial charge capacity of P−Mo− Co−HMSs is about 1328.9 mAh g−1, showing a capacity loss of about 363.3 mAh g−1. The capacity loss during the first cycle can be attributed to the formation of Li2O and the irreversible intercalation of Li+ into the crystal lattice of MoO3 and CoMoO4, as well as other irreversible electrochemical processes such as the electrolyte decomposition and the formation of SEI layer, which is common for most transition metal oxide anodes. P−Mo−Co−HMSs have an initial Coulombic efficiency (CE) of 78.5%, which is comparable to that (81.9%) of bare MoO3 MSs and higher than that (73.9%) of CoMoO4 NPs. According to eqs 1 and 2, we know that Co nanoclusters are in situ formed from CoMoO4 in the first discharge cycle. In the subsequent first charge cycle, the Co nanoclusters are oxidized to CoO instead of CoMoO 4. As the cycle continues, the Co nanoclusters begin to promote the decomposition of Li2O as a catalyst. As a consequence, the CE of P−Mo−Co−HMSs rapidly increases to ∼96.67% in the second cycle and remains above 98% after the 100th cycle. Furthermore, the initial CE of P−Mo−Co−HMSs is also higher than those of previously reported MoO3 anodes (see details in Table 1). Moreover, from the second cycle on, the discharge/charge curves overlap well, and no evident capacity loss is observed, revealing superior cyclability of P−Mo−Co−HMSs, which is closely related to the contribution from the Co catalyst promoting the reversible decomposition of Li2O.31 In contrast, for individual MoO3 MSs and CoMoO4 NPs, they both have shorter discharge plateaus. The bare CoMoO4 NP electrode, in particular, suffers serious capacity fading and only delivers a reversible capacity of 300 mAh g−1 after 100 cycles. The cycling performance of the three samples at a current density of 0.2 A g−1 is shown in Figure 5b. It can be seen that compared with the bare MoO3 MSs and CoMoO4 NPs, P− Mo−Co−HMSs display the highest cycling capacity. The capacity of the P−Mo−Co−HMS anode remains at ∼1350 mAh g−1 during the first 20 cycles and then increases continuously. An exceptionally high discharge specific capacity of ∼1598.7 mAh g−1 has been retained after 100 cycles, which 10111

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addition, the interplanar spacing of 0.51 nm, which can be assigned to LixMoO3 phase (Figure 6h), has also been detected in this lithiation process. This phenomenon is probably due to the incomplete transformation of LixMoO3 into Mo during the first several cycles.8 When the electrode was recharged to 3.0 V, the P−Mo− Co−HMS electrode kept its porous spherical morphology (Figure 7a). The HRTEM image in Figure 7b shows crystal

HMSs also display outstanding cycling performance at high current densities. As shown in Figure 5e, P−Mo−Co−HMSs deliver specific capacities of 1182.8, 875.4, 750.7, and 587 mA h g−1 at 0.4, 1.0, 1.5, and 2.0 A g−1 after 100 cycles, respectively. Impressively, under all test conditions, no obvious capacity fading is observed, and high CE values of nearly 100% are throughout the whole cycles. Even after 200 cycles, P−Mo− Co−HMSs still display relatively high capacities (Figure S10). To further investigate the structure integrity of P−Mo−Co− HMSs, FESEM measurements were performed. As shown in Figure 5f, most of the P−Mo−Co−HMSs maintain the integrated spherical structure after 100 cycles at 0.2 A g−1. Those irregular particles belong to the binder and the carbon black used in assembling the cell as well as individual broken spheres. This result indicates that P−Mo−Co−HMSs have relatively stable microstructure during the cycling tests. To further confirm the electrochemical reaction mechanism (eqs 1−4) of the P−Mo−Co−HMS electrode proposed above, we used a TEM technique to study the microstructure changes of P−Mo−Co−HMS electrode after 10 discharge−charge cycles. Figure 6a presents high-magnification TEM image of

Figure 7. TEM and HRTEM images of P−Mo−Co−HMS electrode recharged to 3.0 V after 10 cycles: (a, c) High-magnification TEM images; (b, d−f) HRTEM images recorded on different areas.

lattices with interlayer distances of 0.213 and 0.51 nm that can be ascribed to the (100) plane of CoO and the (003) plane of LixMoO3.28 We also performed HRTEM images on other different areas, as shown in Figure 7c−f. Clearly, the interplanar spacing of 0.213 nm (Figure 7d,e) can be assigned to the (200) plane of CoO, while the crystal lattice with the interplanar spacing of 0.51 nm (Figure 7f) agrees well with the LixMoO3 phase. These results reveal that the metallic Co and Mo are oxidized to CoO and LixMoO3 during this delithiation process, which is associated with the irreversible reactions in eq 3 and 4. The detection of the CoO phase in the delithiation process further demonstrates that the reduced Co NPs act as the catalyst to prompt the reversible conversion of Li2O, releasing extra Li+ simultaneously.11,14 Next, in an effort to evaluate potential practical application of the as-prepared P−Mo−Co− HMSs, a full cell was constructed with prelithiated P−Mo− Co−HMSs as the anode and commercial LiFePO4 as the cathode, as illustrated in Figure 8a. Figure 8b displays the voltage profiles of the LiFePO4/P−Mo−Co−HMSs full cell with a voltage range of 1.0−3.9 V. The initial charge profile has a flat plateau at around 3.4 V (vs Li+/Li), and the polarization between the charge and discharge profiles is high (ΔV ≈ 1 V), indicating the slow Li+ diffusion rate for the LiFePO 4 cathode.41,42 This is probably due to the hindering of Li+ penetration on the boundary of the LiFePO4 electrode as well as the “activation” process of LiFePO4 during the first cycle.42 Besides, the considerable polarization may be also related to the

Figure 6. TEM and HRTEM images of P−Mo−Co−HMS electrode discharged to 0.01 V after 10 cycles: (a, b, e, f) high-magnification TEM images; (c, d, g, h) HRTEM images recorded on different parts.

the electrode discharged to 0.01 V after 10 cycles, in which we can see that the porous structure of the MSs has been well preserved. The HRTEM images (Figure 6c,d) recorded on two different parts of an individual P−Mo−Co−HMS (Figure 6b) show two distinct interplanar spacings of 0.205 and 0.222 nm, which match well with the (111) plane of Co and the (110) plane of Mo, indicating that MoO3−CoMoO4 was transformed into crystalline Co and Mo.28 Besides the Co and Mo phases, Li2O phase is also detected. As shown in Figure 6e,f (part 1), the interplanar spacing of 0.266 nm can be readily indexed to the (111) plane of Li2O (Figure 6g).7 These results agree well with the reduction reactions of eqs 1−3 described above. In 10112

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to 1.0 C, the capacity of 109.9 mAh g−1 can be delivered. Even suffering from a rapid variation of the current density, the full cell exhibits a relatively stable capacity at each current density. Moreover, when the rate was turned back to 0.1 C, the initial capacity was mostly recovered. These results reveal the superior rate capability of the full cell device. The remarkable electrochemical properties of the asprepared P−Mo−Co−HMSs can be attributed to their CoMoO4-hybrid component and porous structure. On the one hand, the introduction of CoMoO4 into MoO3 can not only promote the reversible decomposition of Li2O due to the catalytic reaction by the in situ formed Co NPs during the electrochemical process (Figure 1b) but also significantly improve the conductivity of the whole electrode (Figure 5c). On the other hand, the porous structure in this hybrid not only ensures the presence of the additional free volume to strengthen the structural integrity associated with repeated Li+ insertion/extraction processes but also provides more electroactive sites, which can increase the contact area, shorten the Li+/e− diffusion distance, and improve the specific capacity of the electrode.16 With the advantages of all these features, P− Mo−Co−HMSs present exceptionally high specific capacity, good cycling stability, and excellent rate capability both in lithium ion half and full cells.

Figure 8. Electrochemical test of the full cell constructed by LiFePO 4 cathode and P−Mo−Co−HMS anode. Schematic illustration (a) and voltage profiles (b) of the full cell. Cycling performance (c) and corresponding CE of the full cell at 0.1 and 1.0 C; (d) rate capability of the full cell at various current densities.

irreversible processes such as the electrolyte decomposition and the formation of SEI layer.43 The initial charge and discharge capacities are 173.8 and 169.2 mAh g−1, respectively, yielding a CE as high as 97.35%. After the first cycle, the full-cell voltage profiles display a sloping potential plateaus at ca. 2.5 V (from 2.0 to 3.5 V) and overlap very well, demonstrating a stable electrochemical behavior. Figure 8c shows the cycling performance and corresponding CE of the full cell at 0.1 C. Clearly, the full cell delivers a specific capacity of 155.7 mAh g −1 after 70 cycles, corresponding to 0.019% capacity fade per cycle. Moreover, the full cell presents an average capacity of 90.75 mAh g−1 at 1.0 C, further demonstrating its excellent cycling stability. In addition, the CE is found to be approaching 100% at both 0.1 and 1.0 C except for the initial cycle. To the best of our knowledge, the above lithium storage performance of the full cell is much better than most of the reported Mo-based and other metal or metallic oxide materials (Table 2). The rate capability of the full cell was also investigated by varying charge−discharge rates from 0.1 to 4.0 C for each 10 cycles. As shown in Figure 8d, the specific capacity is as high as 167.4 mAh g−1 at a rate of 0.1 C. As the rate increases 10 times

CONCLUSION In summary, P−Mo−Co−HMSs have been successfully designed and fabricated through a simple hydrothermal method followed by thermal treatment. The as-obtained P−Mo−Co− HMSs exhibit excellent lithium storage performance in terms of specific capacity, cycling stability, and rate capability when a half cell is assembled using P−Mo−Co−HMSs as an anode material. Ex situ TEM studies on the electrochemical lithiation and delithiation processes reveal that the introduction of CoMoO4 in MoO3 is considered to be mainly responsible for the excellent electrochemical behavior. CoMoO4 can be reduced to Co during the discharge process, and the in situ formed Co NPs can promote the reversible decomposition of Li2O, thus effectively improving the reversible specific capacity and the initial Coulombic efficiencies. Furthermore, the pore structure in P−Mo−Co−HMSs greatly enhances their mechanical strength and provides enough inner void cavity to alleviate volume changes during repeated cycling, thus ensuring

Table 2. Comparison of the Capacity of the Present Work with Previously Reported Full-Cell LIB Devicesa full cells (cathode/anode)

current density (C)

cycle no.

capacity (mAh g−1)

voltage window (V)

year

ref

LiCoO2/TiO2−MoO3 LiFePO4/CoMoO4-ppy LiFePO4/MoO2 LiNi0.5Mn1.5O4/Fe3O4@F-doped carbon nanoparticles LiMn2O4/MnOx/C nanocomposites Li0.85Ni0.46Cu0.1Mn1.49O4/CuO/mesocarbon microbeads LiMn2O4/TiNb2O7 nanofibers LiNi0.5Mn1.5O4/Sn−C composites LiNi1/3Mn1/3Co1/3O2/Si-graphene LiCoO2/Ge-graphene-carbon LiFePO4/Ge nanowires LiFePO4/P−Mo−Co−HMSs

0.18 0.2 0.2 0.65 0.2 0.3 1.0 1.0 0.25 1.0 0.1 0.1 (1.0)

100 60 100 300 100 50 200 100 15 100 30 70

90 148 90 110.7 76.6 83 91.6 120 77 77.1 112.5 155.7 (90.75)

1.0−4.0 2.0−3.8 1.0−3.8 2.0−4.6 2.0−4.1 1.5−5.0 1.8−2.8 3.0−5.0 3.0−4.3 2.5−4.2 2.0−3.8 1.0−3.9

2015 2015 2014 2014 2013 2014 2014 2007 2012 2014 2012

4 44 45 46 47 48 49 50 51 52 53 this work

a The specific capacity is calculated according to the mass of the cathode at various rates (cycling rate 1C = 274 mAh g−1 vs LiCoO2; 170 mAh g−1 vs LiFePO4; 146.7 mAh g−1 vs LiNi0.5Mn1.5O4; 148 mAh g−1 vs LiMn2O4; 145.8 mAh g−1 vs Li0.85Ni0.46Cu0.1Mn1.49O4; 180 mAh g−1 vs LiNi1/3Mn1/3Co1/3O2).

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carried out on a CHI-760E electrochemical workstation with a scan rate of 0.2 mV s−1. The electrochemical impedance spectroscopy (EIS) was obtained by applying a sine wave with amplitude of 5 mV over the frequency range from 100 kHz to 0.01 Hz. In the present work, a full cell was also constructed by combining the prelithiated P−Mo−Co−HMS anode with a LiFePO4 cathode. The prelithiation of P−Mo−Co−HMS anode was performed according to a previously reported method.22,23 For the preparation of the LiFePO4 cathode film, the active material (80 wt %), conductive carbon black (10 wt %), and PVDF (10 wt %) with NMP as the solvent were mixed in a mortar to make a homogeneous slurry, which was uniformly pasted on aluminum foil and finally dried overnight under vacuum at 120 °C. LiPF6 (1 M) dissolved in a mixture of EC/ DMC/DEC (1:1:1, in vol %) was used as the electrolyte, and Celgard 2400 polypropylene film was used as the separator. For the full cell assembly, the mass loadings of the cathode and the anode were 5.2 and 1.0 mg cm−2, respectively. The battery was cathode-limited and was cycled at 0.1 C (1C = 170 mAh g−1 based on the cathode weight) in the voltage range of 1.0−3.9 V.

excellent cycling stability. Additionally, a full-cell constructed by P−Mo−Co−HMSs as an anode and commercial LiFePO4 as a cathode manifests a remarkable Li storage performance, capacity retention, and rate capability. This design concept can be readily extended to other electrode materials for LIBs, especially for the development of those materials with superior capacity and rate performance.

METHODS Materials. All chemicals were used as received without further purification. Cobalt chloride hexahydrate (CoCl2·6H2O, 97%), cobalt nitrate hexahydrate (Co(NO3)2·6H2O, 97%), and ethanol (C2H5OH, 99.9%) were purchased from Beijing Chemical Works. Molybdenyl acetylacetonate (MoO2(acac)2, 99%) was purchased from SigmaAldrich. Synthesis of P−Mo−Co−HMSs. CoCl2·6H2O (2 mmol) and MoO2(acac)2 (1 mmol) were added to 50 mL of C2H5OH under vigorous stirring to form a transparent solution, which was then transferred into an 80 mL Teflon-lined stainless steel autoclave and maintained at 200 °C for 48 h. The resultant black powder was harvested by centrifugation, washed with deionized water and ethanol several times, and then dried at 50 °C for 8 h. Finally, the black powder was calcined at 500 °C for 3 h in air, leading to the formation of P−Mo−Co−HMSs. For comparison, we also prepared another two samples by changing the molar ratios of CoCl2·6H2O to MoO2(acac)2 to 1:1 and 3:1 while keeping other conditions constant. Synthesis of Bare MoO3 MSs. MoO2(acac)2 (1 mmol) was added into 50 mL of C2H5OH under vigorous stirring to form a transparent solution. Then the clear solution was transferred into an 80 mL Teflon-lined stainless steel autoclave and kept at 200 °C for 48 h. The resultant black precipitate was collected and calcined at 450 °C for 3 h in air to obtain the bare MoO3MSs. Synthesis of CoMoO4 NPs. The preparation of CoMoO4 NPs was basically identical to that of the P−Mo−Co−HMSs except the CoCl2· 6H2O was replaced with Co(NO3)2·6H2O. In addition, the molar ratio of Co(NO3)2·6H2O to MoO2(acac)2 was fixed at 1:1. Characterizations. Powder X-ray diffraction (XRD) patterns were obtained with a Bruker D8 X-ray diffractometer at 40 kV, 40 mA, with Cu Kα radiation (λ = 1.54056 Å). The field emission scanning electron microscopy (FESEM), energy dispersive X-ray spectra (EDS), and element mapping of the samples were taken on a Hitachi S-4800 SEM unit. DTG-60AH was used to achieve the thermogravimetric and differential scanning calorimetry analysis (TG-DSC). The element contents of the samples were acquired by using inductively coupled plasma optical emission spectroscopy (ICP-OES) (PerkinElmer). Transmission electron microscope (TEM) images, high-resolution transmission electron microscope (HRTEM) images, and selected area electron diffraction (SAED) patterns were performed on a JEM-2100F with an acceleration voltage of 200 kV. X-ray photoelectron spectra (XPS) were recorded on an ESCALAB 250 spectrometer (PerkinElmer). Nitrogen sorption isotherm was carried out at 77 K on a Belsorp-max surface area detecting instrument (ANKERSMID, Holland). Electrochemical Measurements. For the anode preparation, the active material, conductive carbon black, and polyvinylidene fluoride (PVDF) with a mass ratio of 80:10:10 were mixed and ground in a mortar. N-Methyl-2-pyrrolidone (NMP) was used as the solvent to make homogeneous slurry. The as-resultant slurry was uniformly pasted on a Cu foil and dried at 120 °C for 24 h in vacuum oven. The used electrolyte was 1 M LiPF6 dissolved in a mixture of ethylene carbonate (EC)/dimethyl carbonate (DMC)/diethyl carbonate (DEC) (1:1:1, in vol %). The lithium metal was employed as both the counter and reference electrode. Celgard 2400 polypropylene film was used as the separator. The loading of the active material was about 1.0 mg cm−2. The cell assembly was performed in an Ar-filled glovebox. The electrochemical performance of the cell was evaluated by galvanostatic charge/discharge on a LAND CT2001A battery tester at room temperature under various current densities in the voltage range of 0.01−3.0 V for the half-cell. Cyclic voltammetry (CV) was

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b05150. XRD pattern of the Mo−Co precursor. TG-DSC curve of Mo−Co precursor recorded in air. Typical XPS survey spectrum for the P−Mo−Co−HMSs. FESEM images of the Mo−Co precursors obtained with different reaction times. Typical XRD patterns and FESEM images for the MoO3 precursor and the MoO3 MSs. Typical XRD patterns and FESEM images for the CoMoO4 precursor and the CoMoO4 NPs. N2 adsorption−desorption isotherms of MoO3 MSs and CoMoO4 NPs. The voltage profiles of the first two cycles for P−Mo−Co−HMSs and MoO3 MSs. FESEM images, XRD patterns, and cycle performances for P−Mo−Co−HMSs with 9.6 wt % CoMoO4 and 13.0 wt % CoMoO4, respectively. Cycle performance and corresponding CE of P−Mo−Co− HMSs at 0.4 and 1.5 A g−1 for 200 cycles (PDF)

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Author Contributions †

W.W. and J.Q. contributed equally to this work.

Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (21601014, 21471016, and 21271023) and the 111 Project (B07012). REFERENCES (1) Tao, T.; Glushenkov, A. M.; Zhang, C. F.; Zhang, H. Z.; Zhou, D.; Guo, Z. P.; Liu, H. K.; Chen, Q. Y.; Hu, H. P.; Chen, Y. MoO3 Nanoparticles Dispersed Uniformly in Carbon Matrix: a High Capacity Composite Anode for Li-Ion Batteries. J. Mater. Chem. 2011, 21, 9350−9355. (2) Hassan, M. F.; Guo, Z.; Chen, Z. X.; Liu, H. Carbon-Coated MoO3 Nanobelts as Anode Materials for Lithium-Ion Batteries. J. Power Sources 2010, 195, 2372−2376. (3) Meduri, P.; Clark, E.; Kim, J. H.; Dayalan, E.; Sumanasekera, G. U.; Sunkara, M. K. MoO3−x Nanowire Arrays as Stable and High10114

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DOI: 10.1021/acsnano.6b05150 ACS Nano 2016, 10, 10106−10116

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DOI: 10.1021/acsnano.6b05150 ACS Nano 2016, 10, 10106−10116