Porous CoO Nanostructure Arrays Converted from Rhombic Co(OH)F ...

Article pubs.acs.org/JPCC

Porous CoO Nanostructure Arrays Converted from Rhombic Co(OH)F and Needle-like Co(CO3)0.5(OH)·0.11H2O and Their Electrochemical Properties Liping Zhu,* Zhen Wen, Weimin Mei, Yaguang Li, and Zhizhen Ye State Key Laboratory of Silicon Materials, Department of Materials Science and Engineering, Cyrus Tang Center for Sensor Materials and Applications, Zhejiang University, Hangzhou 310027, People’s Republic of China S Supporting Information *

ABSTRACT: Novel CoO nanostructure arrays on nickel foam with needle-like and rhombic morphologies have been prepared by using urea and hexamethylenetetramine as hydrolysis agents through fluoride-assisted hydrothermal method, respectively. The possible formation mechanism and effect factors of the novel arrays were systematically investigated by X-ray diffraction, scanning and transmission electron microscopies, N2 sorption, X-ray photoelectron spectroscopy, and Fourier transform infrared spectroscopy. The precursor Co(CO3)0.5(OH)·0.11H2O or Co(OH)F plays a crucial role in the formation of needle-like or rhombic arrays. As-prepared precursor arrays can be further converted to corresponding CoO arrays by annealing at 450 °C for 2 h without signification alteration of one-dimensional (1D) morphology. When tested as anodes for lithium ion batteries (LIBs) without the addition of other ancillary materials (carbon black and binder), the synthesized CoO arrays with needle-like and rhombic morphologies deliver ultrahigh initial discharge capacities of 1973.3 and 1447.9 mAh g−1, respectively. In addition, they also maintain high reversible capacities of 710 and 719 mAh g−1 at 0.2 C after 50 cycles, respectively.

1. INTRODUCTION Over the past decade, research in the synthesis of onedimensional (1D) materials has become a focus area due to their unique properties and application.1−4 Particularly, 1D nanostructures are acknowledged as one of the promising solutions to the future generation of lithium ion batteries (LIB) due to their higher surface to volume ratio, faster electron transportation, and lower volume change during the charge− discharge process as compared to the bulk materials.5−8 Recently, 1D nanostructures (nanorods,9 nanowires,10 naotube,11 etc.) with controlled size and morphology have been developed by employing various synthetic approaches. These mainly include sol−gel, thermal decomposition, microemulsion, electrodeposition, solvothermal, and hydrothermal methods.6,12−15 Hydrothermal method has emerged as an efficient and easily scaled-up strategy for preparation of low dimensional nanomaterials. Transition metal oxides (cobalt oxide, nickel oxide, iron oxide, etc.) have been considered as such promising anodes for LIB, because they can exhibit reversible capacities around 2−3 times larger than that of graphite (372 mAh g−1).16 Among them, cobalt monoxide (CoO) has been conceived as a promising anode to display the highest capacity and best cycling performance.17−22 Moreover, 1D CoO nanostructured with porosity can facilitate the molecule/ion diffusing to the inside of an active particle and buffer the volume expansion.20,21,23 © 2013 American Chemical Society

Up to now, most of the CoO-based electrodes of LIB are normally mixed with carbon black and binder and compressed into pallets, in which a large portion of surface of CoO is blocked from the contact with electrolyte, resulting in low electron transfer velocities and electrolyte diffusion efficiency.24 Although reversible capacities of bulk materials can be retained at large values in the first few cycles, these cobalt oxide anode materials exhibit rapid capacity fading during the prolonged cycles. Therefore, free-standing and binder-free 1D CoO nanostructure arrays grown directly on conducting substrates have been highly desirable electrode materials to achieve ideal electrochemical performance. For example, Jiang et al. prepared well-aligned CoO porous nanowire arrays on Ti substrate via hydrothermal method and demonstrated them as advanced anode materials for LIB.25 However, it is still a great challenge to introduce a controllable method of low cost, process simplicity, and high scale-up yield production to synthesize a unique architecture of CoO nanostructure arrays, which can maintain a large reversible capacity. In this work, we report for the first time that porous CoO nanostructure arrays can be induced to grow directly on nickel foam into different morphologies by changing the alkaline reagent via fluorine-assisted hydrothermal synthesis. More Received: June 21, 2013 Revised: September 25, 2013 Published: September 25, 2013 20465

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Figure 1. Schematic illustration of the formation process of porous CoO rhombic and needle-like arrays, respectively.

35 mm × 60 mm was immersed into the reaction solution close to the inner wall of the autoclave. After being sealed and heated at 95 °C for 24 h, the autoclave was cooled naturally to room temperature. The nickel foam coated with pink cobalt hydroxide carbonate nanowire arrays was picked out, rinsed with distilled water several times, and dried at 60 °C under vacuum for 2 h. Afterward, the dried sample was annealed at 450 °C in high-purity Ar atmosphere for 2 h with a ramp of 2 °C/min until the color changed from pink to brown. 2.4. Materials Characterization. The morphologies and structures of both the precursors and the products were characterized by an X-ray diffraction (XRD, Bede D1) system with Cu Kα radiation (λ = 0.15406 nm), field emission scanning electron microscopy (FE-SEM Hitachi S-4800), and high-resolution transmission electron microscopy (HRTEM, FEI F20). Specific surface areas were computed from the results of N2 physisorption at 77 K (TriStar II 3020) by using BET (Brunauer−Emmet−Teller) and BJH (Barrett−Joyner−Halenda). X-ray photoelectron spectroscopy (XPS, Thermo ESCALAB 250) measurement was performed with a monochromatic Al Kα (hν = 1486.6 eV) X-ray source. Fourier transform infrared spectroscopy (FTIR, TENSOR 27) was carried out with a DTGS detector by making pellets with KBr powder, and the resolution was set at 4 cm−1 with a scan number of 32. 2.5. Electrochemical Measurement. Electrochemical measurements were carried out using the coin-type (CR2025) test cells, which were assembled in an argon-filled glovebox (MBRAUM) by using the rhombic or needle-like cobalt monoxide nanostructure arrays as working electrode (diameter, 15 mm). Lithium foil was used as counter electrode and reference electrode. Celgard 2300 was used as the separator membrane. The electrolyte was 1 M LiPF6 in ethylene carbonate (EC)−dimethyl carbonate (DMC) (1:1 by volume). The galvanostatic charge−discharge tests were conducted using a LAND battery program-control test system over a voltage of 0.01−3.0 V at 25 °C. Cyclic voltammogram (CV) tests were carried out on a CHI660D electrochemical workstation at a scanning rate of 0.5 mV s−1 between 0.01 and 3.0 V.

attractively, both of the CoO arrays exhibit combined properties of mesoporous and quasi-single-crystalline structure, and meanwhile exhibit robust mechanical adhesion with nickel foil. When directly tested as LIB anode electrode without any ancillary materials, both CoO arrays with needle-like and rhombic morphologies deliver a high reversible capacity, good cycling performance, and rate capability.

2. EXPERIMENTAL SECTION 2.1. Chemical Materials. All chemicals or materials were of analytical grade and used directly without any further purification prior to usage. The cobalt nitrate, ammonium fluoride, hexamethylenetetramine, and urea were received from Shanghai Chemical Reagent Co. Deionized water (18.3 MΩ) was produced by using a Millipore Direct-Q System and was used throughout the experiments. 2.2. Synthesis of Rhombus-Shaped CoO. In a typical synthesis, 1.45 g of cobalt nitrate (Co(NO3)2·6H2O), 0.37 g of ammonium fluoride (NH4F), and 1.5 g of hexamethylenetetramine (C6H12N4, HMT) were dissolved in 50 mL of deionized water under stirring for 10 min at room temperature. The homogeneous solution was transferred into a 100 mL Teflonlined stainless steel autoclave. Next, a piece of cleaned nickel foam substrate with the size of 35 mm × 60 mm was immersed into the reaction solution close to the inner wall of the autoclave. After being sealed and heated at 95 °C for 24 h, the autoclave was cooled naturally to room temperature. The nickel foam coated with pink cobalt hydroxide fluoride nanorod arrays was picked out, rinsed with distilled water several times, and dried at 60 °C under vacuum for 2 h. Afterward, the dried sample was annealed at 450 °C in high-purity Ar atmosphere for 2 h with a ramp of 2 °C/min until the color changed from pink to brown. 2.3. Synthesis of Needle-like CoO. In a typical synthesis, 1.45 g of cobalt nitrate (Co(NO3)2·6H2O), 0.37 g of ammonium fluoride (NH4F), and 1.5 g of urea (CO(NH2)2) were dissolved in 50 mL of deionized water under stirring for 10 min at room temperature. The homogeneous solution was transferred into a 100 mL Teflon-lined stainless steel autoclave. Next, a piece of cleaned nickel foam substrate with the size of 20466

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3. RESULTS AND DISCUSSION Here, we describe a reliable synthetic strategy for the selforganized nanostructured CoO arrays with tunable morphologies. An illustration of the formation process of porous CoO arrays is shown in Figure 1. The precursor Co(OH)F arrays were formed by hydrothermal method using hexamethylenetetramine (HMT) as an alkaline reagent in the presence of NH4F at 95 °C, and porous CoO rhombic nanorod (NR) arrays could be obtained by annealing the Co(OH)F in an inert atmosphere. When using urea as the alkaline reagent rather than HMT in the same reaction condition, Co(CO3)0.5(OH)· 0.11H2O nanowire (NW) arrays could be obtained on the substrate, and these nanoarrays will convert to porous CoO needle-like NW arrays after heat treatment. Figure 2a shows the X-ray diffraction (XRD) patterns of the Co(OH)F prepared through hydrothermal method using HMT Figure 3. Morphological and structural characterizations of the Co(OH)F NR arrays: (a) top-view SEM image (inset shows a magnified rhombus shape) and (b) cross-sectional SEM images of Co(OH)F NR arrays growing on nickel foam; (c) TEM image of a single Co(OH)F NR; and (d) HRTEM image of the area indicated by the red square in (c) (inset shows the SAED pattern).

Figure 2. XRD patterns of the products: (a) the precursor Co(OH)F NR and (b) final CoO rhombic NW arrays before and after heat treatment, respectively.

F, and O elements (Figure S3, Supporting Information). A typical TEM image (Figure 3c) shows the as-prepared NR has a diameter of approximately 400 nm. The HRTEM image of lattice fringes (Figure 3d) for the area marked with a red square in Figure 3c is 0.252 nm, in agreement with the distance of the (111) crystal face of Co(OH)F. The selected-area electron diffraction (SAED) pattern (inset in Figure 3d) indicates the single-crystalline nature of the product of rhombus-shaped Co(OH)F precursor. Figure 4a shows a typical SEM image of the CoO NR arrays, showing a 1D rhombus-shaped morphology similar to that of the Co(OH)F precursor after annealing. As compared to the

as a hydrolysis reagent in the presence of NH4F at 95 °C for 24 h. All of the diffraction peaks correspond perfectly with orthorhombic Co(OH)F (a = 10.305 Å, b = 4.677 Å, c = 3.126 Å; JCPDS card no. 50-0827). No other detectable peaks from impurities were observed, which indicates the high purity of the prepared product. After annealing at 450 °C for 2 h in Ar atmosphere, the Co(OH)F precursor was completely converted to pure-phase CoO [space group: Fm3m (225); JCPDS card no. 43-1004] (Figure 2b), which could also be confirmed by the visual color changes of the product from pink to brown (Figure S1, Supporting Information). The representative morphological and structural characterizations of the Co(OH)F precursor NR arrays on nickel foam were investigated systematically. A typical low-magnification SEM image shows that large-scale and high density arrays of Co(OH)F NR arrays are uniformly grown on the substrate (Figure S2, Supporting Information). As can be seen in highmagnification SEM images (Figure 3a), the Co(OH)F NR arrays are vertically deposited on the nickel foam and consist of uniform and smooth rhombus-shaped surfaces with a mean edge length of 400−500 nm. A cross-sectional SEM image (Figure 3b) displays that well-organized Co(OH)F NR arrays are aligned perpendicular to the nickel foam and the average thickness is ca. 8 μm. The elemental maps of the Co(OH)F NR arrays clearly demonstrate a homogeneous distribution of Co,

Figure 4. Morphological and structural characterizations of the CoO NR arrays with rhombus-shaped morphology: (a) SEM images of CoO NR arrays (inset shows a magnified rhombic NR); (b) TEM image of a single CoO NR; (c) HRTEM image of the area indicated by the red square in (c) (inset in panel (d) shows the corresponding FFT patterns); and (d) N2 adsorption/desorption isotherm curve of the rhombic CoO NRs and porous volume distribution of the pore size. 20467

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precursor, the surfaces of CoO NR are apparently rough. A high-magnification SEM image of the broken CoO NRs (Figure S4, Supporting Information) shows that numerous nanopores exist in the rhombic building-blocks, arising from the transition from Co(OH)F to CoO. An HRTEM image (Figure 4c) taken from the red square area of Figure 4b reveals a lattice plane with the interplanar distance of 0.25 nm, corresponding to the (111) plane of the fcc CoO cubic cystalline structure. The corresponding fast Fourier trasformation (FFT) patterns (inset in Figure 4c) imply that the CoO NRs are quasi-singlecrystalline in nature and the preferential growth is along the [110] direction, which is also the NR’s longitudinal axis. Interestingly, mesoporosity is introduced to the converted rhombic CoO with quasi-single crystallinity preserved. N2 sorption was used for characterization of the porous structure of rhombic CoO NR arrays and gathering information about the specific surface area and pore size. It can be observed that the isotherms (Figure 4d) are characteristic of a type IV with type H3 hysteresis loop, which confirms the mesoporous structure. BET measurement shows that the CoO NR arrays have a surface area of 16.8 m2/g. The pore size distribution shows a peak at 33.1 nm (inset in Figure 4d). This can be explained by the loss of water molecules and fluorine ions in the lattice during calcination. By using urea rather than HMT as an alkaline reagent in the presence of NH4F, Co(CO3)0.5(OH)·0.11H2O needle-like NWs were formed. After 24 h of reaction, the color of Ni foam changed from original white to pink. After heat treatment, the surface color turned to brown, indicating the chemical composition had been converted (Figure S5, Supporting Information). Figure 5a shows the corresponding XRD pattern

Figure 6. Morphological and structural characterizations of the Co(CO3)0.5(OH)·0.11H2O needle-like NW arrays: (a) top-view SEM image (inset shows a magnified needle shape) and (b) crosssectional SEM image of NW arrays growing on nickel foam; (c) TEM image of a single Co(CO3)0.5(OH)·0.11H2O and the corresponding SAED pattern (inset); and (d) HRTEM image of the area indicated by the red square in (c).

the products is needle-like with diameter in the range of 50− 100 nm. The cross-sectional SEM image indicates that the needle-shaped arrays with a length of about 5 μm are homogeneously well aligned. The length and loading weight of NW could be controlled by altering the hydrothermal time and the concentration of start reagents. The HRTEM image (Figure 6d) marked by the red frame in Figure 6c shows two mutually perpendicular interplanes with lattice fringes of 0.298 and 0.506 nm that correspond very well with the (300) and (020) crystal planes of orthorhombic Co(CO3)0.5(OH)· 0.11H2O SAED pattern (inset in Figure 6c), and HRTEM reveals that those precursors are single-crystalline in nature and the preferential growth is along the [100] direction, similar to those reported by Zeng and Wang.26,27 From the SEM image (Figure 7a), the needle-like morphology of precursor Co(CO3)0.5(OH)·0.11H2O is completely unaltered during the thermal conversion. As compared to the TEM image of the precursor (Figure 6c), porous structure could be clearly observed in the Figure 7b. The corresponding SAED patterns from the square area inset in Figure 7b reveal that these porous CoO are nearly quasi-singlecrystalline with longitudinal axis identified in the [110] direction, similar to the direction of rhombus-shaped CoO preferential growth. Furthermore, the lattice spacing is measured to be 0.21 nm, which agrees well with the separation between (200) interplane distance of the CoO phase. N2 sorption was also used for characterization of the porous structure of needle-like CoO NR arrays. BET measurement shows that the CoO NW arrays have a surface area of 22.9 m2/ g. The pore size distribution shows a peak at 15.4 nm (inset of Figure 7d). We attribute this to the pyrolysis and the release of CO2 and H2O during the thermal conversion. As observed above, it is evident to conclude that different anions greatly influence the morphology of the final products. It is noteworthy that fluorine ions play an essential role in obtaining the nanostructure arrays. The different fluoride source such as NaF or KF is found to have no significant effect on the product of precursor in our experiment (Figures

Figure 5. XRD patterns of the products: (a) the precursor Co(CO3)0.5(OH)·0.11H2O and (b) final CoO needle-like NW arrays obtained by using urea as an alkaline reagent.

of the as-prepared needle-like Co(CO3)0.5(OH)·0.11H2O (JCPDS card no. 48-0083) with orthorhombic crystal structure. The lattice parameters are as follow: a = 8.792 Å, b = 10.15 Å, c = 4.433 Å. The XRD characterizations verified the convertion process shown in Figure 5b. All of the patterns could be undoubtedly indexed to cubic CoO (JCPDS card no. 43-1004) except the peaks induced from Ni foam diffraction. Figure 6a and b shows the SEM images of as-prepared Co(CO3)0.5(OH)·0.11H2O. We can observe that large scale of 20468

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NH3 + H 2O → NH4 + + OH−

CoF+ + OH− → Co(OH)F

When using urea as a hydrolysis reagent, cobalt hydroxide carbonate forms: CO(NH 2)2 + H 2O → 2NH3 + CO2

CO2 + H 2O → CO32 − + 2H+ CoF+ + OH− + 0.5CO32 − + 0.11H 2O → Co(CO3)0.5 (OH)0.11H 2O + F−

After heat treatment, both of the precursors convert to cobalt monoxide: Co(OH)F → CoO + HF↑ Co(CO3)0.5 (OH)0.11H 2O

Figure 7. Morphological and structural characterizations of the CoO needle-like NW arrays: (a) SEM image of CoO NW arrays; (b) TEM image of CoO NW and the corresponding SAED pattern (inset); (c) HRTEM image of the area indicated by the red square in (c); and (d) N2 adsorption/desorption isotherm curve of the needle-like CoO NWs and porous volume distribution of the pore size.

→ CoO + 0.5CO2 ↑ + 0.61H 2O↑

The phase changes are also clearly reflected in Fourier transform infrared (FTIR) spectra (Figure 8). The strong peaks

S6 and S7, Supporting Information). However, if choosing NH4Cl as the structure agent instead of NH4F, we could only obtain the needle-like NW arrays rather than rhombus-shaped NR arrays (Figure S8, Supporting Information). It is worth mentioning that F− is a good complexing ligand for Co2+ to form CoF+ complex that both prevents Co(OH)2 generation and serves to reduce the concentration of free Co2+ ions, lowing the supersaturation to favor the NR or NS growth,9 similar to Zn2+ reported by Saito.29 HMT, a common ammonia-releasing agent, is extensively utilized as a pH buffer to release OH− slowly. The OH− and CoF+ can be combined into a single Co(OH)F phase, which is a prerequisite for preparing CoO with rhombus-shaped morphology. Urea, with a slower hydrolysis rate of releasing ammonia than that of HMT, however, provides carbonate ions simultaneously during the hydrolysis.30−32 Carbonate anions can be gradually intercalated into the interlayers by replacing F− and OH− anions due to the strong affinity to Co2+. The anion exchange process is kinetically driven and induced a dissolution−recrystallization in which carbonate anions act as a structure-directing agent, producing needle-like cobalt hydroxide carbonate.28,33 On the other hand, the disproportionation of formaldehyde may also produce carbonate ions during hydrolysis of HMT at 95 °C, but the amount is trivial and it is also verified by FTIR. The formation of the NR and NW arrays on Ni foam is based on heterogeneous nucleation and growth due to the functional groups (hydroxyl and carboxyl or carbonyl groups) on the substrate.26 With respect to the formation, reactions of the precursor of Co(OH)F or Co(CO3)0.5(OH)·0.11H2O and the final CoO phase can be expressed with the following steps. In the presence of NH4F, Co2+ ions form complexes with F−:

Figure 8. FTIR spectra of (a) Co(OH)F precursor and (b) Co(CO3)0.5(OH)·0.11H2O precursor.

at 3580 and 3503 cm−1 are attributed to the O−H stretching mode, which is characteristic of molecular water and hydrogenbond O−H groups. The shoulder vibration at 3417 cm−1 can be assigned to the O−H groups interacting with fluoride or carbonate anions. The peaks from middle to lower wavenumbers regions in Figure 8b can be indexed to stretching vibrations ν(OCO2), δ(CO3), and δ(OCO) centered at 1500, 829, and 743 cm−1, respectively. The presence of CO32− is due to hydrolysis of urea. The peaks at 964 and 522 cm−1 correspond to δ(Co−OH) and ρ w (Co−OH) bending modes.28 The peak at 1634 cm−1 in Figure 8a is well-known as the bending mode of water molecules. However, no peaks of CO32− are observed in Figure 8a. During the hydrolysis of HMT, it is noteworthy that the disproportionation of formaldehyde is trivial at 95 °C. After heat treatment, the transformation of precursors to cubic CoO is also supported by FTIR and XPS analyses (Figures S9 and S10, Supporting Information).

Co2 + + F− → CoF+

When using HMT as a hydrolysis reagent, cobalt hydroxide fluoride forms: (CH 2)6 N4 + 6H 2O → 4NH3 + 6HCHO 20469

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nanoparticles in the needle-like CoO as compared to that of rhombic morphology. However, as compared to the cycling shapes for needle-like CoO, rhombic CoO showed nicely overlapping redox peaks upon extended scans, indicating rhombic CoO anode materials exhibit better electrochemical stability. To evaluate the electrochemical properties of as-prepared electrodes with needle-like and rhombus-shaped morphologies for LIB, galvanostatic charge−discharge (Figure 10a and b) and cycling tests (Figure 10c) were carried out at a current rate of 0.2 C (1 C = 716 mA g−1) in the voltage range of 0.01−3.0 V. Both first discharge curves produced a long voltage plateau at 1.0−0.8 V, followed by sloping down to the cutoff voltage of 0.01 V. It is observed that initial discharge−charge capacities are 1973.3 and 1364.8 mAh g−1 with a Coulombic efficiency of 69.2% (Figure 10a) for needle-like CoO arrays and 1447.9 and 894.1 mAh g−1 with a Coulombic efficiency of 61.8% (Figure 10b) for rhombic CoO arrays, respectively. The initial capacity loss can be normally attributed to the possible irreversible processes such as electrolyte decomposition and inevitable formation of a SEI layer.22 From the XPS spectra (Figures S11 and S12, Supporting Information), there is evidence for the formation of LiF, Li2CO3, Li−O bonds (e.g., RCH2OCO2Li and LiOH), and a lithium-containing polymer, which reveal the composition of the SEI layer formed on the electrode surface. Despite the large irreversible loss in the first cycle, both electrodes can still attain very high charge capacities. The reversible capacities can reach 710 and 719 mAh g−1 for needlelike and rhombic CoO at the current rate of 0.2 C after 50 cycles, respectively. It should be mentioned that capacties of the needle-like CoO electrode can retain above 1000 mAh g−1 even after 30 cycles. Although the cycling performance of the electrode starts to show some fall off after 30 cycles, the capacity is still comparable to the theoretical capacity of bulk CoO (716 mAh g−1), which is predicted by the conversion reaction mechanism and calculated by the number of transferred mechanism and by the number of transferred electronics in the reaction.16 Actually, the phenomenon is normal in the metal oxide nanomaterials anodes and it has been reported several times,18,34,35 yet the measured capacity was higher than the theoretical capacity. Because LIB battery capacity not only depends on all of the materials available to react, materials with porous structure and large surface area would also display higher reversible capacities that exceeded theoretical capacity value. That is because Li+ ions stored in the interfaces and pores of the porous materials would take part in the reaction, and the electrolyte decomposition would lead to the formation and decomposition of a gel-like polymeric layer on the surface of the active nanoarrays particles.35−37 In addition, there are many other mesoporous surfaces due to the etching of the formation of HF in the solution, which can provide more storage space and thus store more energy. The further investigations on the influence mechanism of F− and other halogen elements to metal oxide anodes will be done to clarify the mechanism. Therefore, the mesoporous surfaces and electrolyte decomposition can contribute to the observed extra capacity. On the other hand, the third reason can be based on the two-phase capacitive behavior of the Co−Li2O interface that allows for the storage of Li+ at the lithium compound side, whereas electrons are localized at the metallic side, leading to a charge separation and extra capacity.38,39 For these reasons, thus, the synergetic effect of the porous rhombic and needlelike nanostructured arrays favors larger extra specific capacity

The presence of mesoporosity during calcination treatment allows easier electrolyte penetration and can buffer the massive volume expansion and contraction occurring during the electrochemical reactions. Additionally, the robust mechanical adhesion between CoO and Ni foam would make our products a potential candidate for anodes of LIBs directly. Figure 9

Figure 9. Cyclic voltammogram tests of (a) rhombic CoO arrays and (b) needle-like CoO arrays at a scan rate of 0.5 mV s−1.

presents cyclic voltammograms (CV) of the electrodes made from rhombic and needle-like CoO NW arrays at a scan rate of 0.5 mV s−1. In the first cycle, it can be seen that both samples have two cathodic peaks observed at 0.83 and 0.39−0.42 V, which can be assigned to the irreversible reactions with the electrolyte and the reduction reaction of CoO into Co and the formation of amorphous Li2O, respectively. A oxidation peak was located at around 2.23 V in the anodic process, corresponding to the oxidation of Co to CoO. Moreover, the anodic scan curves exhibit a broad peak located at about 1.5 V, which may be attributed to the decomposition of the solid electrolyte interphase (SEI) layer. In subsequent cycles, the cathodic peaks shift to 1.4 and 0.75 V and the peak shape is almost unchanged, suggesting high reversibility of lithium storage.25 Nevertheless, there was one obvious difference in the shapes of the two curves. The area and intensity of cathodic peaks at around 0.4 V of rhombic CoO (Figure 9a) are smaller than those of needle-like CoO (Figure 9b), which implies a greater amount of Li can react with needle-like CoO rather than rhombic CoO. This may be attributed to the smaller 20470

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Figure 10. Electrochemical characterization: galvanostatic charge−discharge curves of (a) needle-like and (b) rhombus-shaped CoO arrays electrode in the voltage range of 0.01−3.0 V (vs Li/Li+) at a current of 0.2 C rate. (c) Cycling performance of the needle-like and rhombus-shaped CoO arrays at a current rate of 0.2 C. (d) Rate capability of the needle-like and rhombus-shaped CoO arrays electrodes at various current rates of 0.2−4 C.

4. CONCLUSIONS In summary, we developed a novel strategy to synthesize CoO arrays on nickel foam substrate with tunable morphologies by changing the alkaline reagent during fluoride-assisted hydrothermal synthesis. The synergistic effects of the different alkaline reagents and ammonium fluoride on the morphology of the precursors have been investigated and discussed in detail. The obtained two oxide structures have a drastic impact on LIB performance, and both needle-like and rhombus-shaped CoO arrays display high discharge capacity and good cycling capability as promising anode materials for LIB.

than the theoretical capacity. Higher capacity in CoO arrays with needle-like morphology than rhombus shape arrays can be mainly attributed to the larger specific surface area. It is worth noting that the capacity of CoO with rhombic morphology is higher than that with needle-like morphology after 40 cycles. The tips of needle-like CoO nanorod are thinner than that of rhombus morphology, and may be broken after 40 cycles.7 An increased Coulombic efficiency of 96% from the second cycle of rhombic NR also indicates high charge−discharge reversibility of the anode material for LIBs. The excellent cycling performance of CoO arrays with rhombus morphology is evidently believed to benefit from the unique morphology and structure, which is in agreement with the above CV results. Figure 10d shows the rate capability of the two electrodes; CoO arrays with needle-like morphology exhibit 1465 mAh g−1, when first cycled at 0.2 C, and the following capacity values at other C-rates: 1050 mAh g−1 at 0.5 C, 400 mAh g−1 at 1 C, 150 mAh g−1 at 2 C, 90 mAh g−1 at 4 C, and back to 1453 mAh g−1 at 0.2 C. The CoO arrays with rhombus morphology present a discharge capacity of 990, 940, 611, 286, 100, and 950 mAh g−1 at the same rates. The capacities gradually fade at high discharge rates, which could be ascribed to the collapse of the as-prepared materials during the lithiation−delithiation at a high current density. Further investigations will be done to improve the capacity retention at high discharge rates.

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ASSOCIATED CONTENT

S Supporting Information *

Optical images of cobalt-precursors and CoO nanostructured arrays, SEM images of the pristine porous nickel substrate, element mapping of Co, F, and O, SEM image of broken rhombic CoO nanorods, XRD patterns and top-view SEM images of the cobalt-precursors nanoarrays using NaF or KF as fluoride source, SEM images of cobalt-precursors using NH4Cl as reaction agents, FTIR spectra and XPS spectrum of CoO nanostructured arrays, and XPS spectra of CoO nanostructured arrays cycled in electrolytes. This material is available free of charge via the Internet at http://pubs.acs.org. 20471

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AUTHOR INFORMATION

Corresponding Author

*Tel.: +86-571-87951958. Fax: +86-571-87952625. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

This work was supported by the National Natural Science Foundation of China (51072181 and 51372224), and the Science and Technology Department of Zhejiang Province (project no. 2010R50020).

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