Integrated Charge Transfer in Colloidal Cu–MnO Heterostructures for

Jul 14, 2014 - TUM CREATE, 1 CREATE Way, #10-02 CREATE Tower, Singapore 138602 ... 1 CleanTech Loop, #06-04 CleanTech One, Singapore 637141, ...
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Integrated Charge Transfer in Colloidal Cu−MnO Heterostructures for High-Performance Lithium Ion Batteries Hui Teng Tan,†,‡,§ Xianhong Rui,†,∥ Ziyang Lu,† Chen Xu,†,‡ Weiling Liu,† Huey Hoon Hng,† and Qingyu Yan*,†,‡,§ †

School of Materials Science and Engineering, Nanyang Technological University, Singapore 639798, Singapore TUM CREATE, 1 CREATE Way, #10-02 CREATE Tower, Singapore 138602, Singapore § Energy Research Institute@NTU, 1 CleanTech Loop, #06-04 CleanTech One, Singapore 637141, Singapore ∥ School of Energy and Environment, Anhui University of Technology, Maanshan, Anhui 243002, China ‡

S Supporting Information *

ABSTRACT: Nanodimensional monodispersed Cu−MnO heterostructures were synthesized via a facile solution-based method. Cu−MnO heterostructures with narrow size distribution were successfully achieved in the presence of organic surfactants to prevent agglomeration during the growth process. Furthermore, the unique architecture of carboncoated Cu−MnO (Cu−MnO@C) core−shell heterostructures obtained after a thermal annealing process preserved the electrical integrity of the electrode via the conductive copper “nanobridge” to provide an efficient electron transfer pathway between the active materials and the current collector. On the other hand, the amorphous carbonaceous shell evidently gives a protective layer to attain structural integrity throughout the electrochemical measurements. Enhancement in the electrochemical performance can be reflected by the excellent cycling stability and high rate capability of the Cu−MnO@C heterostructures as compared to the MnO@C nanoparticles.

1. INTRODUCTION Lithium ion batteries (LIBs) have dominated the commercial market as power sources in the past few decades. The important breakthrough in LIBs can be traced back to the implementation of nanosized transition metal oxides (TMOs) as the electrode materials, endowing them with unexpected electrochemical kinetics by exhibiting both bulk and surface properties.1 Followed by this pioneering research discovery, successful attempts were also made to demonstrate the contribution of nanostructural features of TMOs in tailoring the energy storage and conversion system.2 In an effort to explore appealing TMOs as electrode materials, MnO holds great promise to provide distinct lithium storage properties for energy-related applications. Besides having high theoretical capacity (755.6 mA h g−1), MnO also exhibits low voltage hysteresis as compared to other transition-metal oxides.3,4 More importantly, from the thermodynamics point of view, MnO has a relatively low electromotive force (emf) value (1.032 V vs Li+/Li) that is beneficial for the lithium uptake process, making the conversion reaction more favorable.5 However, the low electronic conductivity characteristic of MnO becomes an obstacles to realizing the full potential of its electrochemical performance. In addition, a problematic issue that impedes its long-term cycling performance falls on the huge volumetric variation of MnO during the phase transformation (MnO ↔ Mn), with repeated swelling and shrinking of the crystal structure rendering structural collapse. Furthermore, the constraints imposed by nanostructure engineering remained a © 2014 American Chemical Society

research concern owing to its higher surface energy. Upon restricting the dimension to the nanoregime, the surface energy of nanoparticles (NPs) becomes inevitably high, creating a strong tendency to agglomerate. In addition, they are also highly reactive toward side reactions and subsequently induced some uncontrollable phenomena, such as depletion of electrolytes, self-discharge, and shorter usage life.1 In order to compensate these disadvantages, hybridization approaches were always adopted to reduce the detrimental effects encountered. Apart from direct growth of TMOs on carbonaceous materials such as graphenes6,7 and carbon nanotubes,8,9 to increase the electronic conductivity, in addition to addressing the severe agglomeration issue, anchoring active materials on 3-D metallic current collectors such as nickel foam,10,11 copper nanocables,12,13 and stainless steel foil14,15 is a new trend in LIBs. The integrated conductive pathway provides very efficient charge transfer via the intimate contact between the active materials and the current collector, giving rise to better electron transfer and rate capability. On the other hand, some research groups also exploited the conformal coating of amorphous carbon on the surface of NPs as the conductive as well as the protective matrix.16,17 In the literature, the high-temperature solution-based method was extensively applied in chemical synthesis to Received: May 22, 2014 Revised: July 9, 2014 Published: July 14, 2014 17452

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decompose thermally. The heating mantle was then removed and the mixture was cooled to room temperature naturally. The MnO NPs were precipitated and washed with hexane/acetone (volume ratio 3:1) several times, followed by drying at 50 °C in a vacuum oven overnight. 2.1.3. Synthesis of Cu−MnO Heterostructures. In the synthesis of Cu−MnO heterostructures, 10 mL of Mn−oleate (0.5 M) was added into a round-bottom flask containing 20 mL of 1-octadecene, and the mixture was degassed at 120 °C for 1 h under protection of argon. The mixture was held at 300 °C for 10 min before the temperature was stabilized at 200 °C. Subsequently, 0.1 or 0.5 mL of Cu−oleate solution (0.1 M) was injected rapidly into the mixturedenoted as 0.1Cu−MnO and 0.5Cu−MnO, respectivelyand the mixture was aged for 5 min. The Cu−MnO heterostructures were precipitated and washed with hexane/acetone (volume ratio 3:1) several times, followed by drying at 50 °C in a vacuum oven overnight. 2.1.4. Synthesis of Carbon-Coated Cu−MnO Heterostructures. In order to transform the organic surfactant into a thin layer of carbon shell surrounding the MnO and Cu−MnO NPs, a thermal annealing step was performed to carbonize the samples. The as-prepared samples were annealed under the protection of argon at 400 °C for 1 h with 2 °C min−1 ramping rate. Carbon-coated MnO, 0.1Cu−MnO, and 0.5Cu−MnO are denoted as MnO@C, 0.1Cu−MnO@C, and 0.5Cu−MnO@C, respectively. 2.2. Materials Characterizations. X-ray diffraction (XRD) was performed on a Shimadzu thin film diffractometer with Cu Kα radiation (λ = 0.15406 nm) for compositional analysis. The diffraction pattern was collected within 10°−80° (2θ). A JEOL JEM 2100 transmission electron microscope (TEM), operating at 200 kV, was used for capturing high-resolution TEM (HRTEM) images, bright-field scanning transmission electron microscopic (BF-STEM) images, and energy dispersive X-ray (EDX) mapping. Inductively coupled plasma (ICP, Dual-View Optima 5300 DV ICP-OES system) was employed to measure the elemental contents of the as-prepared samples. Concentrated hydrofluoric acid (40%) was used to digest the samples prior to the measurements. Raman spectroscopy was conducted using the WITec CRM200 confocal Raman microscopy system with a laser wavelength of 488 nm and a spot size of 0.5 mm. Calibration was done using the Si peak at 520 cm−1 as a reference. 2.3. Electrochemical Characterizations. The electrodes were prepared by dispersing the active materials, carbon nanotubes (conductive matrix), and 8% wt of polyvinylidene fluoride (PVDF) binder in a weight ratio of 7:2:1 in Nmethylpyrrolidinone (NMP). After homogeneous mixing, it was casted on a copper foil and dried in avacuum oven overnight to remove the solvent. The electrochemical measurements were carried out on a coin-type cell using metallic lithium as reference and counter electrode, a Celgard 2400 membrane as the separator, and the electrolyte solution obtained by dissolving 1 M LiPF6 into a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (EC/DMC, 50:50 w/w). The galvanostatic charge and discharge cycles in the potential range of 0.01−3.0 V were conducted on a NEWARE multichannel battery test system, while the cyclic voltammetry (CV) was performed on Solartron analytical equipment (model 1470E), sweeping on a scan rate of 0.2 mV s−1. Electrochemical impedance spectroscopy (EIS) measurements were conducted at open circuit potential on an impedance spectrum analyzer (Solatron, SI 1255B impe-

precisely control the dimension and geometry of the NPs without sacrificing their size uniformity. This strategy has been advocated by several research groups in yielding high-quality colloidal NPs upon thermal decomposition of the precursor in surfactant-assisted organic solvents. For instance, self-assembled two-dimensional ferrite/carbon composites and monodisperse MnxNi1−xO and MnyCo1−yO metal alloy mono-oxide NPs were successfully synthesized for LIB application.18,19 Besides, Sn and Sn/SnO2 nanocrystals with tunable sizes were also successfully achieved.20 Although monodispersed nanocrystals have been successfully synthesized by many research groups, it is still of great challenge to functionalize them into a heterostructure in the hopes of enhancing their effectiveness. On the other hand, multiple fabrication steps have made the hybridization approach complicated. Here, we reported a facile, nonhydrolytic heating-up method to synthesize colloidal Cu−MnO heterostructures that afford high crystallinity and narrow size distribution. The versatility of this synthetic approach provides a platform for in situ growth of metallic Cu that allows strong coupling between the conductive species with the electrochemical active materials. Therefore, promising electrochemical performances were achieved by breaking down the obstacles associated with the electronic conductivity issue of MnO. In this regard, the metallic copper contributed in maintaining intact contact between the active materials and the current collector, enhancing the electron transfer process to guarantee the electronic integrity of the electrode. On the other hand, the protective layer of amorphous carbon that surrounded the MnO NPs that formed after annealing was able to sustain the tensile stresses upon the phase transformation. The synergistic effects could be reflected by the stability performance and rate capability, allowing the Cu−MnO@carbon heterostructures to experience less capacity fading upon cycling and to preserve greater capacity at high C rate. Since this solution-based method is versatile for synthesis of TMOs, it might open up a new milestone in energy storage systems, paving the way for lighter and longer calendar life batteries for commercial use.

2. EXPERIMENTAL SECTION 2.1. Synthesis. 2.1.1. Synthesis of Metal−Oleate Complex. Metal−oleates complexes were synthesized by modifying a reported procedure.21,22 Using Mn−oleate as an example, the preparation steps are described as follows: MnCl2 (2 mmol) and Na−oleate (4 mmol) were dispersed in 10 and 20 mL of methanol, respectively, in separate bottles to obtain clear solutions. After that, the manganese precursor solution was added dropwise into the Na−oleate solution, and highly viscous brownish precipitate can be observed instantly upon sonication. The Mn−oleate complex was washed with methanol several times to remove the residual sodium salts and stored in a vacuum oven overnight to evaporate the solvent completely. The preparation step for Cu−oleate resembled that of Mn− oleate, except the MnCl2 was substituted with CuCl2·2H2O. 2.1.2. Synthesis of MnO NPs. The MnO NPs were synthesized by thermal decomposition of manganese−oleate precursor under protection of argon. In the synthetic steps, the Mn−oleate complex was redispersed in 20 mL of 1-octadecene (ODE) in a three neck round-bottom flask. Then, 10 mL of the solution (0.5 M) was transferred into a round-bottom flask and degassed at 120 °C for 1 h to get rid of the moisture. Subsequently, the temperature was raised to 300 °C within 15 min and refluxed for 10 min to allow the precursors to 17453

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diffraction peaks for Cu upon increasing the concentration of Cu−oleate precursor. The TEM and BF-STEM were employed for morphological and compositional analyses, respectively. Figure 2 provided the TEM and HRTEM images of the as-prepared samples with a different concentration of Cu− to Mn−oleates as well as the elemental X-ray mapping for the heterostructures. The narrow particle size distribution of the MnO@C and Cu−MnO@C heterostructures shown in Figure 2a,d,g suggested that the surfactant-assisted thermal decomposition approach is promising in producing monodisperse nanocrystals. More precisely, the bare MnO@C NPs were found to be achievable in the size range of 5−8 nm. Increasing the molar concentration of Cu precursor inevitably gave rise to particles with larger sizes (10− 18 nm), but their size homogeneity was retained. The lattice dspacing of MnO@C shown in Figure 2b was estimated to be 0.222 nm, which was consistent with the diffraction peak corresponding to the MnO(200) plane. In the subsequent step, the epitaxial growth of Cu species on MnO can be evidenced from the observation of Cu(111) plane with 0.208 nm lattice dspacing in Figure 2e,h, which was indicative of epitaxial growth of Cu on MnO NPs Additionally, the compositional information on Cu species extracted from the TEM-EDX shown in Figure S1 (see the Supporting Information) was also inconsistent with such an observation. Coupling with the compositional EDX mapping in BF-STEM, the elemental distribution of Mn and Cu could be evidently perceived in Figure 2c,f,i. In this case, no preferential growth of Cu on MnO NPs was observed, giving rise to random elemental distribution of Cu on the surface of MnO (Figure 2f,i). In Figure S2 (see the Supporting Information), the SAED patterns of the asprepared samples that presented as sharp diffraction spots encompassed in a ring structure were indicative of their highly crystalline nature and random crystallite orientation. Upon increasing the concentration of Cu to 0.6 M, severe particle agglomeration was found, as illustrated in Figure S3 (see the Supporting Information). In addition, the ICP results shown in Table 1 further confirmed the coexistence of the Mn and Cu species and also resolved the metal contents of Cu−MnO@ carbon heterostructures, in which the molar ratio of Cu:Mn increased (0.06:1 and 0.2:1 for 0.1Cu−MnO and 0.5Cu−MnO, respectively) with increasing the concentration of Cu precursors. 3.2. Formation Mechanism of the as-Prepared Samples. The concept of complete separation between the nucleation and growth steps is always applied in an attempt to synthesize monodisperse NPs, which are synonymous with high surface area and shorter ions/charges transport pathways for surface-mediated electrochemical reactions.25 Hence, metal− oleates were used as metal−surfactant precursors to induce homogeneous nucleation and growth in the synthetic step, which can eliminate the undesirable nucleation of intermediate phases during the growth process.26 The schematic representation of the plausible formation mechanism of colloidal Cu− MnO@carbon heterostructures is proposed in Scheme 1. When the Mn−oleate complex was heated in the noncoordinating solvent ODE, the thermal energy transformed the oleate precursors into MnO NPs at 300 °C. In preparing the Cu− MnO heterostructures, hot injection above the decomposition temperature of Cu−oleate was executed to drive a high supersaturation level in the reaction solution instantaneously, inducing the Cu NPs to grow on the surface of MnO to reduce the free energy of heterogeneous nucleation. In contrast to the

dance/gain-phase analyzer; computer software ZView) in the frequency range of 10 kHz to 0.01 Hz with an ac perturbation of 5 mV. It should be noted that the mass of the active materials was counted as the total mass of the carbon-coated Cu−MnO nanocomposites.

3. RESULTS AND DISCUSSION 3.1. Morphology and Phase Composition. The colloidal Cu−MnO@carbon heterostructures were synthesized via a thermal decomposition of oleate-chelated precursors in noncoordinating hydrocarbon solvent. In brief, pyrolysis of the carboxylate salts (Mn−oleates and Cu−oleates) in noncoordinating hydrocarbon solvent at elevated temperature produced highly crystalline Cu−MnO NPs with narrow size distribution by taking advantage of the homogeneous nucleation and growth of their corresponding oleate precursors in ODE solvent. Owing to the higher reactivity of Cu−oleate precurors, the Cu nanocrystals were epitaxially grown on MnO in the latter stage by injecting the precursors into the solution at lower temperature. The one-pot synthesis method not only functionalized the MnO with metallic Cu but also induced the strong anchoring of the capping agents (oleates) on the surface of MnO simultaneously. Hence, the sterically bulky organic molecules could stabilize the nanocrystals against Brownianmotion-driven aggregation. With the aim to carbonize the capping layer into amorphous carbon, an annealing process was executed at 400 °C under argon protection. The XRD patterns shown in Figure 1a evidenced the formation of the rock-salt structure of highly crystalline MnO@

Figure 1. XRD pattern of the as-prepared samples: (a) MnO@C, (b) 0.1Cu−MnO@C, and (c) 0.5Cu−MnO@C.

C. Sharp diffraction peaks located at 34.9°, 40.6°, 58.8°, 70.3°, and 73.9° can be indexed to the (111), (200), (220), (311), and (222) planes, respectively, of cubic phase MnO (JPDS 780424).23 No characteristic peaks were observed for crystalline impurities besides a broad hump that appeared at 23° ascribed to the amorphous glass holder. In the presence of Cu foreign species, the peak intensity of the MnO was attenuated gradually while the diffraction peaks of elemental Cu emerged for 0.1Cu−MnO@C and 0.5Cu−MnO@C, illustrated in parts a and b or Figure 1, respectively. The peak positions were consistent with metallic copper (JCPDS 04-0836) with a facecentered cubic (fcc) arrangement.24 On the other hand, the crystallographic information also showed gradual enhanced 17454

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Figure 2. Structural characterizations of the as-prepared samples: (a) TEM images, (b) HRTEM images, and (c) EDX mapping of MnO@C; (d) TEM images; (e) HRTEM images, and (f) EDX mapping of 0.1Cu−MnO@C; and (g) TEM images, (h) HRTEM images, and (i) EDX mapping of 0.5Cu−MnO@C. The regions marked in parts d and g were magnified in parts e and h, respectively. The scale bars in parts c, f, and i are 10 nm.

sequential addition of precursors, the great discrepancy in chemical reactivity between the Mn− and Cu−oleates caused the formation of discrete Cu and MnO NPs instead of forming Cu−MnO heterostructures when both of them were heated to 300 °C simultaneously (see the Supporting Information, Figure S4a,b). As a result, the high reaction temperature caused Cu to grow larger at the decomposition temperature of MnO,

Table 1. Elemental Composition of the Heterostructures a

metal content (wt %)

a

entry

Cu

Mn

molar ratio (Cu:Mn)

0.1Cu−MnO@C 0.5Cu−MnO@C

3.08 11.29

42.92 49.71

0.06:1 0.20:1

Metal content was tested by ICP.

Scheme 1. Schematic Representation of the Proposed Formation Mechanism of MnO/Cu@C Heterostructures

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the first cycle, demonstrating the improved kinetics of the active materials upon reacting with lithium. This phenomenon was explained as the amorphization effect of the active material originating from the structural alteration after the first lithiation process.31 It is also worth pointing out that the possibility of trace amounts of MnOx was irreversibly reduced to Mn(II) at 1.2 V and overoxidation of MnO to higher oxidation states (Mn3+, Mn4+) remained by observing an insignificant hump located at 2.13 V.30 On the other hand, the charge/discharge voltage profiles of the as-prepared electrodes at a current density of 0.2 A g−1 (0.26C; 1C = 756 mA g−1) between 0.01 and 3.0 V were illustrated in Figure 4d−f. The initial discharge capacity of MnO@C, 0.1Cu−MnO@C, and 0.5Cu−MnO@C are 1627, 1560, and 1367 mA h g−1, respectively. It has been proposed that the extra capacity as compared to the theoretical value stemmed from the formation of a SEI layer in the low-voltage region and the surface reaction of the electrode materials with lithium.2 Although the first-cycle Columbic efficiency of the asprepared samples is relatively low (46%, 50%, and 52% for MnO@C, 0.1Cu−MnO@C, and 0.5Cu−MnO@C, respectively), the electrodes stabilized in the second cycle and gave rise to 80%, 89%, and 91% of Columbic efficiency for MnO@C, 0.1Cu−MnO@C, and 0.5Cu−MnO@C, respectively. In addition, the voltage hysteresis values of the as-prepared samples derived from the voltage profiles (Figure 4d−f) at 40% state of charge (SOC) in the second cycle were estimated to be 0.82, 0.78, and 0.72 V for MnO@C, 0.1Cu−MnO@C, and 0.5Cu−MnO@C, respectively, which is relatively low compared to other transition-metal oxides.32 The origin of voltage hysteresis is associated with the charge transfer kinetics, and hence, the incorporation of Cu could improve the sluggish phase transformation process of MnO species, alleviating this kind of hysteresis. To evaluate its composition−activity relationship, a series of comparative electrochemical measurements was also conducted. It is known that the rate capability of the active materials is highly dependent on its electronic conductivity, which is also reflected by our preliminary results. Herein, the eminent role of Cu was manifested at higher C rate charge/ discharge cycling, as shown in Figure 5a. It revealed that the 0.5Cu−MnO@C showed an outstanding rate capability at high C rate, although it performed inferiorly at lower C rate. At current density as high as 3 A g−1 (3.97 C), the 0.5Cu−MnO@ C could still deliver 200 mA h g−1, surpassing the other two counterparts that could only achieve 119 and 51 mA h g−1 for 0.1Cu−MnO@C and MnO@C, respectively. As shown in Figure 5b, the long-term cycling stability of the heterostructures outperformed the single phase active material. Although the initial discharge capacity of bare MnO@C exceed the Cu− MnO@carbon heterostructures, it suffered from rapid capacity fading and ultimately can only reach an average of 410 mA h g−1 after 100 cycles. On the contrary, the 0.1Cu−MnO@C and 0.5Cu−MnO@C heterostructures stabilized within the first 10 cycles and preserved about 604 and 629 mA h g−1, respectively, after 100 cycles. Despite the lower initial discharge capacity, the presence of Cu in the heterostructures resulted in excellent capacity retention with capacitive lost less than 15% after the 10th cycle, whereas a drastic capacity drop of 46% was experienced for bare MnO@C. To examine the cycling performance of the Cu-containing samples at high current density, the 0.5Cu−MnO@C heterostructures were cycled at 1 and 3 A g−1. As illustrated in Figure S6 (see Supporting

significantly affecting the uniformity in particle size and morphology. Current successful synthesis of monodisperse nanocrystals relied on the steric stabilization of the anionic oleate molecules. The long chain oleates of the precursor can contribute to different aspects: (1) function as the coordinating agents to stabilize the metal salts in hydrocarbon solvent and (2) prevent rapid coalescence of nuclei during the growth process via the steric repulsion of their bulky chains. To achieve full utilization of the surface-chelated organic molecules, the in situ carbonization of the organic surfactants can be done through the postsynthesis thermal annealing process under inert gas to provide a protective layer (∼2 nm) with amorphous nature surrounding the nanocrystals (see the Supporting Information, Figure S5).17 Furthermore, the existence of the poorly crystalline conformal carbon coating of the as-prepared samples was further examined using Raman spectroscopy, where the D and G bands of graphitic materials were manifested at ∼1580 and ∼1340 cm−1, respectively, as shown in Figure 3.

Figure 3. Raman spectra of the as-prepared samples.

3.3. Comparison of Electrochemical Performances of All Electrodes. The redox mechanistic study of the asprepared samples was investigated using CV to probe the electrochemical behaviors of the active materials. It was found that the CV curves of the bare MnO@C and the Cu−MnO@ carbon heterostructures were almost identical, indicating the electrochemical inertness of metallic Cu to lithium and eliminating the possibility of the presence of its oxide derivatives. As depicted in Figure 4a−c, irreversible reduction peaks that evolved at 0.58 and disappeared in the subsequent cycles can be attributed to formation of a solid electrolyte interface (SEI) as a consequence of organic electrolyte decomposition.27,28 Apart from those alloy anode systems, MnO adapted to a conversion reaction mechanism upon lithium uptake and extraction.29 In this regard, the cathodic peak at 0.18 V represents the prominent voltage trend of MnO vs Li (MnO + 2Li+ + 2e− ↔ Mn + Li2O), where complete reduction of MnO to metallic Mn and formation of amorphous Li2O occurred concurrently. Its redox couple along the anodic scan could be observed at 1.31 V, attributed to the oxidation of Mn back to its higher valence state.3,30 Interestingly, the reduction potential of MnO shifted positively to 0.26 V after 17456

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Figure 4. Cyclic voltammograms of (a) MnO@C, (b) 0.1Cu−MnO@C, and (c) 0.5Cu−MnO@C electrodes in the potential range of 0.01−3 V at scan rate of 0.2 mV s−1 for the first three cycles. Galvonostatic charged/discharged curves of (d) MnO@C, (e) 0.1Cu−MnO@C, and (f) 0.5Cu− MnO@C electrodes at current density of 0.2 A g−1 for the first three cycles.

Figure 5. (a) Cycling performance of the as-prepared samples at different current densities. (b) Cycling performance of the as-prepared samples at 0.2 A g−1 (0.265 C).

region corresponds to the Warburg impedance, W. As shown in Figure 6b, all the as-prepared samples showed similar blocking behavior by exhibiting small variation in Rsf+ct and Rb values, as summarized in Table S1 (Supporting Information) after the first five cycles. Subsequently, the values for Rsf+ct and Rb decreased with increasing Cu content, and a larger discrepancy in ohmic resistances that stemmed from the charge transfer process and intrinsic conductivity between the electrodes was perceived after 100 cycles (Table S1, Supporting Information), indicating that the internal resistance of the as-prepared samples can be arranged in the following order: MnO@C > 0.1Cu−MnO@C > 0.5Cu−MnO@C. This observation provided concrete evidence that the electronic conductivity increases upon increasing the Cu content, as the thickness of the amorphous carbon layer coated on the surface of MnO that may affect the charge and electron diffusion was almost the same. However, the existence Cu has degraded the specific capacity of the electrode, as this electrochemically inert species

Information), the 0.5Cu−MnO@C heterostructures retained 85% and 35% of their second-cycle discharge capacity at 1 and 3 A g−1, respectively, after 100 cycles. Such a promising electrochemical performance is a solution to the calendar life issue of lithium ion batteries. Figure 6a shows the equivalent circuit for data fitting, which consisted of Re (the ohmic resistance of electrolyte and cell components),33 Rsf+ct (surface film resistance arising from the diffusion of lithium ion across the solid electrolyte interface and charge transfer resistance at the electrode/electrolyte interface),34 CPEsf+ct (capacitance arising from surface film and charge transfer),35 Rb (bulk resistance associated with the electronic conductivity of the electrode),36 CPEb (bulk capacitance),35 and W (Warburg impedance arising from the solid-state diffusion of lithium in the bulk phase).37 The semicircle observed at the high-frequency region indicated the presence of Rsf+ct, while the second semicircle in the medium frequency region corresponded to Rb. The linear low-frequency 17457

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Figure 6. Nyquist impedance plots of the as-prepared samples (a) after 5 cycles and (b) after 100 cycles.

theless, the presence of weak diffraction spots in a circular pattern of the cycled samples indicated their lower degree of crystallinity as compared to fresh samples.

shows no capacitive contribution, in spite of providing an improvement in electronic conductivity. Furthermore, the trend was preserved after 100 cycles, as illustrated in Figure 6c.This indicated that the excellent electrochemical reversibility of 0.5Cu−MnO@C upon charge/discharge was endowed by the existence Cu conductive species, giving rise to the superior cycling stability. The enhanced electrochemical performance of the Cu−MnO@carbon heterostructures can be attributed to the following aspect: Incorporation of metallic Cu is beneficial in maintaining the electrical integrity between the electrode and the current collector. The conductor Cu bridged the electron transfer between the active materials MnO and the current collector via their intimate contact, as implied by the apparent improvement in stability test and high rate cycling. To investigate the implication of carbon coating on the nanocrystals, post-mortem analyses of the as-prepared samples were conducted, as shown in Figure 7a−c. It was found that the

4. CONCLUSIONS In the present study, the colloidal chemical approach provided an effective protocol to synthesize Cu−MnO heterostructures with uniform dimension and geometry, which could be summarized as the thermal decomposition and hot injection strategies. The nanodimensional monodispersed Cu−MnO heterostructures were readily transformed into the Cu−MnO@ C core−shell hybrid after executing the thermal carbonization step under argon protection. Electrochemical evaluation of the Cu−MnO@C evidenced the practicability of rationally designed heterostructures as an anode for LIBs by achieving excellent cycling stability and C rate capability. It was revealed that the conformal carbon coating prevented the structural deformation upon cycling, while the embedded metallic Cu preserved the electrical integrity between the electrochemical active species and the current collector. Rational design on the architecture of the Cu−MnO@C heterostructures endowed them with an integrated charge transfer capability, giving rise to significant electrochemical enhancement. Since the synthesis of heterostructures with different compositions falls within the realm of possibility, versatile hybrid system can also be developed to empower LIBs with capacitive enhancement responses.

Figure 7. TEM images of the electrodes after 100 cycles: (a) MnO@ C, (b) 0.1Cu−MnO@C, and (b) 0.5Cu−MnO@C.

■

ASSOCIATED CONTENT

S Supporting Information *

pulverization process broke them down into smaller particles after the cycling test, but their original morphology was retained, showing a particle size of approximately 5 nm for all of the as-prepared samples. Therefore, we believed that the spatial restriction of the graphitic layer has cushioned the drastic volume changes upon lithiation/delithiation processes and restricted them from undergoing severe structural deformation. In Figure S7 (see Supporting Information), the lattice d-spacing of the cycled samples could be clearly visualized, which ruled out the possibility of being amorphous after cycling. Never-

TEM−EDX elemental analysis of the Cu-containing MnO samples, TEM and SAED pattern of the as-prepared samples, TEM images of Cu−MnO with the addition of 0.6 M of Cu− oleate precursors, TEM images of the Cu-containing MnO samples where both the Mn and Cu precursors were mixed and heated simultaneously, TEM images of the as-prepared samples that indicated the thickness of amorphous carbon layer, cycling performances of 0.5Cu−MnO@C at different current densities, and TEM and SAED pattern of the cycled samples. This 17458

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: +65 6790 4583. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We gratefully acknowledge Singapore MOE AcRF Tier 1 grants RG2/13, A*STAR SERC grant 1021700144, Singapore MPA 23/04.15.03 grant, and Singapore National Research Foundation under CREATE program: EMobility in Megacities.

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