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Article Cite This: ACS Omega 2017, 2, 7647-7657
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A Convenient Synthesis Route for Co3O4 Hollow Microspheres and Their Application as a High Performing Anode in Li-Ion Batteries Padalingam Gurunathan,† Pedda Masthanaiah Ette,†,§ Narayanan Lakshminarasimhan,‡,§ and Kannadka Ramesha*,†,§ †
CSIR−Network of Institutes for Solar Energy (CSIR−NISE), CSIR−Central Electrochemical Research Institute-Madras Unit, CSIR−Madras Complex, Taramani, Chennai 600113, India ‡ Functional Materials Division and §Academy of Scientific and Innovative Research (AcSIR), CSIR−Central Electrochemical Research Institute, Karaikudi 630003, India S Supporting Information *
ABSTRACT: A new, rapid, and cost-effective method for synthesizing hollow microspheres (HMSs) of cobalt oxide (Co3O4) using the phloroglucinol−formaldehyde gel route is reported here. Further, the synthesized hollow Co3 O 4 microspheres were investigated as an anode material for Liion batteries. The Co3O4 hollow spheres exhibited excellent electrochemical performance and cycling stability, for example, a capacity of 915 mA h g−1 was obtained at 1 C rate over 350 cycles. The material also exhibited good performance at high rates, viz., capacities of 500, 350, and 250 mA h g−1 at 10 C, 25 C, and 50 C, respectively, with good capacity retention over 500 cycles. The excellent electrochemical performance of Co3O4 can be ascribed to the porous nanoarchitecture that provides a short diffusion length for Li+ ions and high electrolyte percolation in the porous structure. Additionally, the thin porous wall of the nanocages provides an effective way to overcome the issues associated with the volume change occurring during Li charge/discharge. The conversion of Co3O4 into Co upon discharge was also probed by measuring the magnetic properties. sional nanowires, nanobelts,10 nanorods,11 nanoneedles,12 and nanotubes,13 two-dimensional nanosheets,14 three-dimensional architectures,15 nanoparticles,16 and nanocubes.17 Similarly, hollow porous structured nanocage materials that combine the advantages of both nanothickness and void space have attracted increased attention for use in Li-ion battery applications. Such a hollow architecture enables increased electrode−electrolyte contact area and also provides additional space to overcome the strain associated with the Li insertion−deinsertion process.18−20 Recently, we have used a resorcinol (R)−formaldehyde (F) gel method to prepare SnO2 hollow microspheres (HMSs) that exhibited superior electrochemical properties.21 Carbon aerogels or hollow structures can be prepared using a resorcinol (1,3-dihydroxybenzene) and formaldehyde mixture.21,22 The aqueous mediated polycondensation of resorcinol and formaldehyde gives a cross-linked polymeric gel (RF gel).22−25 However, because of the water content, the RF aquagel is not suitable for supercritical drying as it takes a long time to dry, so a tedious solvent exchange step is required while making the carbon aerogels or hollow porous structures.26,27 The production procedure for carbon aerogels or hollow
1. INTRODUCTION Rapid depletion of fossil fuels, global warming, and increased pollution levels have necessitated energy turnaround scenario to seek alternative renewable energy sources for the sustainable future. Various energy conversion/storage technologies have been rigorously developed, and among them, rechargeable batteries have gained much of the attention as they can be used in portable devices as well as for automobile and grid applications. Because of their high energy density, in addition to their long cycle life and light weight, Li-ion batteries are becoming the obvious choice for energy storage.1−5 Recently, significant progress has been made in Li-ion battery technology, both in terms of identifying new electrode materials and new mechanisms, pushing the realizable limits of energy density and power density further. Recently, it has been discovered that metal oxides, when reduced to nanosize, exhibit remarkable reactivity with lithium, usually referred to as “conversion reaction”, showing high capacity.6,7 For example, nanostructured cobalt oxide (Co3O4) is expected to exhibit a high capacity (about 890 mA h g−1), which is twice that of the graphite anode used in commercial Li-ion batteries. However, practical use of Co3O4 electrodes is still hindered by poor cycling stability and rate capability.8,9 To overcome such issues, several efforts have been devoted to prepare nanoarchitectured materials with different morphologies including one-dimen© 2017 American Chemical Society
Received: September 4, 2017 Accepted: October 26, 2017 Published: November 7, 2017 7647
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spheres. It can be seen that (Figure 2c) the particle surface is, in turn, made up of interconnected nanoparticles. The porous nature of the shell and its formation through aggregation of smaller nanoparticles were further supported by transmission electron microscopy (TEM) (Figure 2d). On the basis of the morphological studies, we propose the initial formation of a trough/pit on the surface of the Co/C hollow sphere, which on further heating in an oxygen atmosphere, opens up due to the stress developed during the annealing process as well as due to the escape of CO2 from inside21,31,38 as shown in Scheme S1. Further, the Co/C composite and Co3O4 HMSs were characterized by Fourier transform infrared (FT-IR) spectroscopy (Figure S1a,b). The absorption bands in the regions of 800 and 1050−1125 cm−1 could be assigned to the Co−OH, −CH2, and C−O stretching vibrations.39 A small amount of cobalt oxide is present in the Co/C sample, as understood from the peak at around 519 cm−1. The band at 1620 cm−1 represents the bending vibration mode of water molecules.39 The broad and sharp band between 3392 and 3555 cm−1 could be assigned to the O−H stretching modes of water and surface hydroxyl groups.39 After annealing the Co/C composite at 500 °C for 2 h in an oxygen atmosphere, strong peaks at 574 and 670 cm−1 are observed, which are characteristic absorption peaks of the cobalt−oxygen (Co−O) bond.39 The broad peak at 3417 cm−1 and the shoulder peak at 1632 cm−1 are typical of O−H vibrations due to adsorbed moisture.39 The Raman spectra of the Co/C composite (Figure S2a) show two broad peaks: the peak at 1585 cm−1 [G band] indicates sp2 graphitic carbon, whereas the band at 1345 cm−1 [D band] corresponds to sp3 disordered carbon.40 Figure S2b shows the Raman spectra of the Co3O4 HMSs, revealing four Raman active modes at 470 (A1g), 510 (Eg), 608 (2F2g), and 675 cm−1 (A1g).41,42 The Raman mode at 675 cm−1 (A1g) is due to the atoms at octahedral sites whereas the Eg and F2g modes can be assigned to atoms at the tetrahedral site and octahedral oxygen.41,42 Thermal stability and carbon content were evaluated using thermogravimetric analysis (TGA) by heating the dried precursor in air. TGA for the Co/C composite (Figure S3a) shows a weight loss of 18−20% below 100 °C, which can be assigned to the removal of adsorbed moisture.43,44 A weight loss of 65% observed between 350 and 450 °C is mainly due to the removal of carbon as CO2.43 No weight loss was observed in the case of the Co3O4 sample (Figure S3b). 2.2. Growth Mechanism of Co3O4 HMSs and Their Characterization. To understand the growth mechanism, the product formed during the synthesis was characterized using powder XRD (Figure S4). Initially, the Co/C composite was formed by heating CoCl2/P−F gel at 500 °C in an argon atmosphere, as evidenced by powder XRD. The reflections in the XRD pattern were identified as due to the (111), (200), and (220) planes of cobalt (ICDD PDF Card No. 00-0150806). Further heat treatment was carried out between 200 and 500 °C under an oxygen atmosphere to form Co3O4. When the Co/C composite was heated at 200 °C in an oxygen atmosphere, the XRD pattern showed the evolution of the CoO phase (ICDD PDF Card No. 01-078-0431), in addition to the major reflections of the cobalt phase. At 300 °C, CoO was the major phase, and the reflections corresponding to Co became less intense, indicating that most of the cobalt phase was oxidized. The Co3O4 phase (ICDD PDF Card No. 00-0421467) started to appear when the temperature was increased to 400 °C, and finally, the pure Co3O4 phase was obtained at 500
porous structures can be effectively simplified in terms of both the rate of polymerization as well as the sol−gel drying process by using formaldehyde and alcohol-soluble phloroglucinol (Phl) instead of resorcinol.26−33 The higher degree of polymerization of phloroglucinol over resorcinol is due to its enhanced electron density in the 2nd, 4th, and 6th ring positions.32,34 The resulting carbon aerogels possess a higher surface area and uniform porous structure.30−36 However, to the best of our knowledge, there is no report on the synthesis of metal oxide porous structures using the phloroglucinol− formaldehyde (P−F) gel method, which is also a rapid and cost-effective process. In this work, for the first time, we report the synthesis of Co3O4 HMSs by using the P−F gel method. We further show that the Co3O4 HMSs, when tested as an anode, exhibited fast charging capability (50 C rate) with good capacity retention over 500 cycles owing to their hollow and porous morphology. The superior electrochemical performance and the easy synthesis of the Co3O4 HMSs suggest their practical use in LIBs.
2. RESULTS AND DISCUSSION 2.1. Phase Formation and Morphology. Figure 1a shows the powder X-ray diffraction (PXRD) pattern of the product
Figure 1. Powder XRD patterns of (a) cobalt/C composite prepared by calcination of CoCl2/P−F gel in argon atmosphere at 500 °C and (b) Co3O4 HMSs prepared by annealing Cobalt/C composite in oxygen at 500 °C.
formed after initial carbonization of the gel under an argon atmosphere. The reflections in the pattern could be indexed based on the crystalline cubic system of cobalt (ICDD PDF Card No. 00-015-0806). The XRD pattern of the final product obtained after heating in an oxygen atmosphere is shown in Figure 1b, and all of the reflections were indexed based on the standard pattern of cubic spinel Co3O4 (ICDD PDF Card No. 00-042-1467). These results clearly show the formation of the Co/C composite during the initial heating of the cobalt/P−F resin under an argon atmosphere, which was oxidized to Co3O4 during the subsequent oxygen annealing process. The morphology of the as-prepared Co3O4 HMSs was analyzed by FESEM and compared with that of the Co/C composite. Figure 2a shows the spherical morphology of the Co/C composite particles. The FESEM images of the oxygenannealed sample (Co3O4 HMSs) are shown as Figure 2b,c, and they reveal that the final product is composed of porous hollow 7648
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Figure 2. FESEM images of (a) cobalt/C composite (inset shows single sphere of carbon-coated Co/C composite), and (b, c) Co3O4 HMSs at 50 and 200k× magnifications, (d) TEM image of Co3O4 HMSs.
°C. It should be noted that sample soak time was kept at 2 h at each of the above-mentioned temperatures. The formation mechanism of carbon xerogels using the polymer-assisted phenolic−formaldehyde gel method is well understood,22−36 and it can be further extended to the synthesis of nanostructured metal oxides using metal−carbon xerogels. Co3O4 HMS formation can be explained based on the following three steps. In the first step, a gel was formed by the addition of phloroglucinol, triblock copolymer (F127), formaldehyde, and the metal precursor.27,29,31 The second step was drying of the formed gel, and the third step involved carbonization followed by oxidation of the metal−carbon composite.27 So, the first step involves polymerization of phloroglucinol (Phl) and formaldehyde in an ethanol medium in the presence of concentrated HCl catalyst.24,27,29,31−33 Further introduction of triblock copolymer (F127) to the phlorogluconic resin will facilitate the formation of a mesoporous structure through hydrogen bonding.29−33 Meanwhile, the Co2+ ions are chelated with hydroxyl groups (−OH) of the phloroglucinol moiety in the P−F polymer chain. The second step is drying of the wet gel (cobalt-P−F Gel) in a hot air oven at 100 °C for 5−6 h. The final step involves carbonization of the dried gel under an argon atmosphere. The heating rate and carbonization temperature were also found to be important in determining the textural and electrochemical properties of carbon aerogels.26,27 To investigate the effect of temperature on morphology, the dried gel was carbonized between 200 and 500 °C, and the morphologies of the resultant products were examined by FESEM. Figure S5a shows that spherical particles (1−2 μm) of Co/C composite were formed when the CoCl2/P−F precursor was heated at intermediate temperatures of 200−300 °C under an argon atmosphere. Figure S5b shows the FESEM image of the Co/C composite obtained after heat treatment at 500 °C in an argon atmosphere. Further, the Co/C composite was heated in an oxygen atmosphere between 200 and 500 °C to obtain the
Co3O4 phase, and Figure S5c−i shows the corresponding FESEM images. Figure S5c−g shows the initiation and propagation of micron-sized hollow spheres. Further, Figure S5h,i shows the FESEM images of the final product obtained at 500 °C, which exhibited a hollow spherical structure. Carbon acts as a sacrificial template for the formation of Co3O4 HMSs. The approach is based on the initial formation of micrometer sized P−F spherical gels through self-assembly along with the incorporation of the cobalt precursor. Gel formation is essential in obtaining Co3O4 hollow spheres through the creation of a sacrificial carbon template. Carbonization of the dried gel under an argon atmosphere results in the formation of spherical particles (1−2 μm) of Co/C composite. Formation of Co3O4 hollow spheres from the Co/C composite can be explained as follows: during the initial low temperature oxidation, the surface oxidizes to form a mechanically strong outer surface layer. The trench formation on the Co/C spherical wall during the oxidation is due to the development of stress through the formation and sintering of Co3O4 particles. On further heat treatment at 500 °C, such weak regions will rupture, relieving the stress. In the meantime, the interior carbonaceous material decomposes to form gaseous CO2 that upon release, results in a hollow interior. The escape of gaseous product from the interior will further facilitate the propagation of cracks along the weak regions of the shell wall resulting in the formation of a cage-like structure. The escape of CO2 also makes the surface highly porous. Any cobalt species present in the core will move to the surface during sintering. 2.3. Evaluation of Electrochemical Properties. The electrochemical performance of the Co3O4 HMSs as an anode material for Li-ion batteries was investigated by cyclic voltammetry. Figure 3a shows the cyclic voltammetric (CV) data for the initial four cycles with the scan rate of 0.1 mV s−1 in the potential range 0.001−3.0 V versus Li+/Li. The reactivity of Co3O4 with Li during the electrochemical reduction and oxidation process is given as follows: 7649
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Figure 3. (a) Cyclic voltammograms of Co3O4 HMSs in the voltage range of 3.0−0.001 V at 0.1 mV s−1 scan rate, (b−e) discharge/charge profiles of Co3O4 at different rates: (b) 1 C rate, (c) 10 C rate, (d) 25 C rate, and (e) 50 C rate. (f) Rate capability performance of Co3O4 HMSs (0.2−50 C rate). discharge
shoulder peak at around 1.1 V can be assigned to the formation of the partially irreversible solid electrolyte interphase (SEI) layer.45−49 As reported earlier,45,47,49 the organic component of the SEI layer can dissolve/reform reversibly, however, the inorganic component formation is irreversible. Upon further cycling, the cathodic peak exhibited a shift from 0.68 to 1.13 V due to the polarization of the electrode46,49,50 and remained constant, revealing good cycling stability. Further, to establish the cycling stability of the Co3O4 HMSs as a Li-ion battery anode, galvanostatic discharge/charge tests were carried out for several cycles. Figure 3b−e shows the
Co3O4 + 8Li+ + 8 e− XooooooooY 3Co + 4Li 2O charge
Hence, in the CV plot, during the first cycle, a well-defined cathodic peak is observed at 0.68 V and an anodic peak is observed at 2.08 V versus Li+/Li. In the first cathodic scan, a peak at around 0.68 V corresponds to the initial reduction of cobalt oxide to cobalt metal (Co3O4 to Co) accompanied with the formation of amorphous Li2O.45,46 During the anodic scan, a broad peak observed at around 2.08 V can be assigned to the oxidation of Co to Co3O4 and decomposition of Li2O.46 The 7650
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Figure 4. (a−d) Long cycling stability and Coulombic efficiency of Co3O4 HMSs at different rates: (a) 1 C, (b) 10 C, (c) 25 C, and (d) 50 C. The cells were cycled between 0.001 and 3.0 V.
that Co3O4 HMSs exhibit good rate capability, a desirable property for power applications. The cycling performance and the Coulombic efficiency of the Co3O4 HMSs sample are shown in Figure 4a−d. At 1 C rate, the first discharge and charge capacities (Figure 4a) are about 1624 and 920 mA h g−1, and the Coulombic efficiency (the ratio between charge and discharge capacities) for the first cycle is 56.2%. The reason for the initial irreversible capacity loss can be attributed to the irreversible SEI formation and to the large volume change occurring during charge−discharge.51−55 From the second cycle onwards, the capacity of the electrode gradually decreased, and at the 15th cycle, it reached a value of 376 mA h g−1. However, the Coulombic efficiency remained between 97 and 100% for several hundred cycles (Figure 4a).52,55 It is possible that the SEI formation continues for the initial few cycles, accounting for the continuous capacity loss and that afterwards, the electrode demonstrated a stable cyclability.45,54,55 In Figure 4a, it can be noted that from the 15th cycle onwards, the capacity started to increase, and it stabilized at 895 mA h g−1 after the 110th cycle. A similar behavior of increasing capacity during charge/discharge cycles was also reported for a few other oxide materials as well.31,52−61 The initial capacity drop was ascribed to irreversible SEI formation through the degradation of the electrolyte.45,51,53−58 The increasing trend in capacity is usually attributed to the reversible growth of a polymeric gel-like film on the electrode surface.31,51,55,59−63 At 1 C rate, even after 350 cycles, a discharge capacity of 915 mA h g−1 was observed, which is slightly higher than the theoretical
discharge/charge profiles of mesoporous Co3O4 HMSs between 0.001 and 3.0 V versus Li+/Li at 1 C, 10 C, 25 C, and 50 C rates. At 1 C rate, the discharge capacities were 1635, 860, 964, and 915 mA h g−1 for the 1st, 100th, 200th, and 350th cycles, respectively. Even at 10 C rate, the delivered discharge capacities were 692, 429, 457, 466, 467, and 500 mA h g−1 for the 1st, 100th, 200th, 300th, 400th, and 500th cycles, respectively. At 25 C rate, the corresponding discharge capacities were 599, 320, 328, 340, 342, and 350 mA h g−1 and at 50 C rate, the capacities were 475, 234, 232, 230, 232, and 250 mA h g−1 (for the 1st, 100th, 200th, 300th, 400th, and 500th cycles, respectively). For practical applications of porous/ hollow/nanomaterials, the tap density is one of the crucial parameters that impact the volumetric capacity as well as the volumetric energy densities. The tap density of the synthesized Co3O4 HMSs was measured to be around 0.97 g cm−3. Therefore, the volumetric capacities of the Co3O4 HMSs at different current rates were 945 mA h cm−3 (at 1 C), 221 mA h cm−3 (at 10 C), 165 mA h cm−3 (at 25 C), and 113 mA h cm−3 (at 50 C), which correspond to volumetric energy densities of 1134 Wh L−1 (1 C), 266 Wh L−1 (10 C), 198 Wh L−1 (25 C), and 135 Wh L−1 (50 C), respectively. Co3O4 HMSs exhibit excellent rate capability at various discharge rates (0.2−50 C), and the results are shown in Figure 3f. After cycling at higher rates, the cell was subjected to cycling again at lower rates to investigate the stability of the Co3O4 HMSs for varying current rates. It was observed that (Figure 3f) the cell capacity reverts back to the initial capacity, confirming 7651
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capacity value of Co3O4 (890 mA h g−1), and this can be attributed to the activation of the electrode during long cycling.31,53,54 The high reversible capacity beyond the theoretical capacity can probably be attributed to the porous nature of the nanocages. In nanostructured systems, the interfacial lithium storage may play an important role in the overall capacity. The pores within the nanocages can provide additional sites for lithium storage, resulting in a higher capacity compared to the theoretical capacity.45,48,52−55 The other explanation for the additional capacity observed in transition metal oxides is that additional lithium storage is possible between grain boundaries of Li2O and metal particles formed during discharge.20,45,52−55 At the end of the 350th cycle, the capacity of the Co3O4 anode was around 915 mA h g−1 at 1 C rate (Figure 4a). Such a stable capacity retention can be attributed to the hollow porous nature of the Co3O4 HMSs. Similarly, Figure 4d shows that at 50 C rate, after the 500th cycle, the sample exhibited a capacity retention of 87% relative to the 1st cycle capacity. On the basis of the above results, it may be envisaged that Co3O4 HMSs exhibit an ultralong cycle life that would endorse their application as an anode material for LIBs. The superior performance is due to their morphology that provides sufficient buffer space to accommodate volume changes during lithium insertion−deinsertion.21,40,44,55,62,63 In addition, the high surface area and the hierarchical porous network structure render a short path length for lithium diffusion during insertion/ extraction enabling high rate performance.49,55,64−73 To ascertain the excellent performance of the Co3O4 HMSs prepared in the present work, a comparison of the electrochemical performance of the Co3O4 HMSs with that of other previously reported Co3O4 materials was made and is listed in Table S1 (Supporting Information). Further, we compared the electrochemical behavior of the Co3O4 HMSs with that of commercial Co3O4 nanoparticles at 1 C rate, and the results are shown in Figure 5. Although the Co3O4 HMSs showed a
the hollow structure can alleviate the strain developed due to volume change during the charge−discharge process, enabling good structural stability.21,40,44,49,53 The superior transport kinetics of the Co3O4 HMSs was also confirmed by electrochemical impedance measurements. Figure 6 shows the Nyquist plots after completing the 2nd, 50th, 100th, 200th, and 350th discharge cycles. Impedance plots show a semicircle in the mid frequency region, which can be assigned to charge transfer resistance (Rct).21 The inclined straight line in the low-frequency region can be attributed to Warburg diffusion of Li+ ions into the bulk of the active material.43 Even after 350 cycles at 1 C rate, the charge transfer resistance of the electrode had not increased much, confirming the highly reversible lithium storage capability of the Co3O4 HMSs. Further, we compared the charge transfer resistance of the Co3O4 HMS and commercial Co3O4 samples before and after cycling (Figure 7). The charge transfer resistance of the Co3O4 HMS electrode is lower than that of the commercial Co3O4 electrode. The conversion reaction between Li and Co3O4 resulting in the formation of cobalt particles during discharge was qualitatively monitored by measuring changes in the magnetic properties of the pristine and discharged samples (ex situ magnetic studies). The zero field cooled (ZFC) and field cooled (FC) curves of Co3O4 are shown in Figure 8a. With decreasing temperature, the magnetization increased up to 33 K, and then decreased upon further cooling to 5 K, showing a cusp around 33 K. This behavior of temperature-dependent magnetization is typical of antiferromagnetic ordering. The observed Néel temperature of 33 K is similar to the values reported in the case of Co3O4 nanoparticles and double-shell hollow particles.74,75 The M−H data measured at 300, 30, and 5 K are shown in Figure 8b, and the data are linear without any saturation. A small hysteresis curvature observed at 5 K with a coercieve field (Hc) of 146 Oe indicates the antiferromagnetic ground state.75 The magnetization and M−H results of the discharged sample at different temperatures are shown in Figure 8c,d. The temperature-dependent magnetization result of the discharge sample is different from that of the pristine sample. The ZFC and FC data (Figure 8c) show a deviation below 250 K, indicating that the sample behaves as a superparamagnet above that temperature. The M−H curve (Figure 8d) measured at 5 K shows saturation at a field of less than 1 T with a high Hc (986 Oe), and this is similar to the reported value of colloidal Co nanoparticles, revealing the ferromagnetic nature.76 To understand the origin of the long cycle capacity retention, post-mortem SEM analyses of electrodes cycled at 10 C, 25 C, and 50 C rates were carried out after completing 500 charge− discharge cycles. For this, coin cells were dismantled inside an argon-filled glove box and the electrodes were thoroughly washed with fresh electrolyte (ethylene carbonate (EC) and dimethyl carbonate (DMC)). The morphology of the asprepared (Figure 9a) and electrochemically cycled (Figure 9b− d) electrode material was examined using FESEM. It was found that the hollow nanocage-like morphology was retained even after 500 cycles at 10 C, 25 C, and 50 C rates.
Figure 5. Comparison of electrochemical cycling performances of Co3O4 HMSs and commercial Co3O4 nanoparticles at 1 C rate.
capacity retention of 99% at 1 C rate (915 mA h g−1), commercial Co3O4 nanoparticles showed only 2% capacity retention (20 mA h g−1) after 350 cycles. The superior performance of the Co3O4 HMSs can be attributed to the hollow structure that provides sufficient contact between the electrolyte and the active material, and facilitates the rapid electron transfer due to small diffusion length.21,44,45,55 Also,
3. CONCLUSIONS We have synthesized Co3O4 HMSs using a cost-effective phloroglucinol−formaldehyde (P−F) gel method. When used as an anode for Li-ion batteries, the Co3O4 HMSs showed excellent cycling stability and high capacity retention over 350 7652
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Figure 6. Electrochemical impedance spectra (EIS) of Co3O4 HMSs recorded at different discharge cycles (2nd, 50th, 100th, 200th, 350th cycle). Right hand side figure is the enlarged region of the spectra.
Figure 7. (a) Electrochemical impedance spectra of as-prepared electrodes (Co3O4 HMSs and commercial Co3O4). Right hand side figure shows enlarged region of the spectra. (b) Impedance spectra of Co3O4 HMS and commercial Co3O4 electrodes after 350th cycle.
Co upon discharge, as evidenced by changes in the magnetic properties.
cycles at 1 C rate. The material also exhibited good performance at higher rates (such as 10 C, 25 C, and 50 C) with good capacity retention over 500 cycles. The origin of such remarkable behavior can be attributed to the hollow morphology of the particles that provides ample electrode/ electrolyte contact and a shorter diffusion length for Li as well as accommodating volume changes during intercalation/ deintercalation. Further, magnetic measurements on pristine and discharged samples revealed the conversion of Co3O4 into
4. EXPERIMENTAL SECTION 4.1. Synthesis of Co3O4 Hollow Microspheres. The synthesis of Co3O4 HMSs was carried out by adopting the facile phloroglucinol−formaldehyde (P−F) gel method reported elsewhere.30−33 In a typical synthesis, 2 g of CoCl2·6H2O (Merck Chemicals), 1 g of phloroglucinol dihydrate (Across 7653
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Figure 8. Temperature-dependent magnetization of (a) pristine and (c) discharged sample obtained after discharging. The M−H traces of (b) pristine and (d) discharged sample measured at different temperatures, and the insets show a closer view of the results.
were made using a physical property measurement system (PPMS, Dynacool, Quantum Design) as in our earlier work.37 The magnetization was measured under zero field cooled (ZFC) and field cooled (FC) warming conditions in the temperature range between 5 and 380 K with an applied field of 100 Oe. The M−H measurements were made at room temperature and at low temperatures in the magnetic field between −9 and +9 T. For magnetic measurements on the electrochemically discharged sample,37 Swagelok-type cells were opened inside an Ar-filled glove box. The collected samples were washed with dimethyl carbonate (DMC) a few times and the dried samples were packed in a capsule for magnetic measurements. 4.3. Electrochemical Properties . The electrochemical performance of the Co3O4 HMS samples was evaluated galvaniostatically by using a VMP3Z (Biologic) multichannel battery testing system. The working electrode was prepared by mixing active material Co3O4 HMSs, SP carbon, and poly(vinylidene fluoride) as binder in a weight ratio of 70:20:10. NMethyl-2-pyrrolidone was used as a blending solvent. The slurry was coated on Cu foil of thickness 9 μm and dried at 120 °C for 12 h under vacuum. CR2032-type coin cells were assembled in an argon-filled glove box (MBRAUN) with Li as the counter electrode and Celgard-2400 as separator. 1 M LiPF 6 in ethylene carbonate (EC)/dimethyl carbonate (DMC)/diethyl carbonate (DEC) (2:1:2 by volume) was used as electrolyte. The typical active material loading was 0.5− 1.5 mg cm−2. For comparison purpose, cells were also made with commercial Co3O4 powders (Sigma-Aldrich chemicals,
Organics Ltd), 1 g of F127 triblock copolymer (Sigma-Aldrich), and 2−3 mL of 37% formaldehyde (Rankem Chemicals) were dissolved in 25 mL of ethanol. This solution was vigorously stirred for 30 min to get a clear solution. The pH of the solution was adjusted to 1−2 by adding a few drops of concentrated HCl, followed by heating the solution in a water bath at 80 °C for 3 h to obtain cobalt/phloroglucinol− formaldehyde (Co-P−F) gel. The gel was dried in a hot air oven at 100 °C for 5 h. Subsequently, the dried gel was transferred to a tubular furnace and calcined at 500 °C for 2 h under flowing argon followed by heating under an oxygen atmosphere for 2 h at 500 °C to obtain the final product of Co3O4 HMSs. 4.2. Material Characterization. The synthesized products were characterized by powder X-ray diffraction (PXRD) using a Bruker D8 Advance Da Vinci diffractometer equipped with Cu Kα radiation (λ = 1.5418 Å). The morphologies of the products obtained at various stages were analyzed by using field emission scanning electron microscopy (FESEM; Carl Zeiss SUPRA55VP) and transmission electron microscopy (TEM; Technai-20 G2). Raman spectra were recorded in the wavelength range of 200−2500 cm−1 using a Renishaw Laser Raman microscope. Fourier transform infrared (FT-IR) spectra of the samples were recorded using a spectrometer (Bruker, Optick GmbH) by the KBr pellet technique. Thermal analysis of samples was carried out in air (heating rate 5 °C min−1) using a thermogravimetric analyzer (NETZSCH, model STA449F3). The temperature-dependent magnetization and M−H measurements of the pristine and discharged samples 7654
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Figure 9. FESEM images of Co3O4 HMS (a) as-prepared electrode before cycling, (b) after 500 cycles at 10 C rate, (c) after 500 cycles at 25 C rate, and (d) after 500 cycles 50 C rate.
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ACKNOWLEDGMENTS The authors thank the Council of Scientific and Industrial Research (CSIR), India, for supporting this work under the TAPSUN programme. P. Gurunathan and Pedda Masthanaiah Ette thank CSIR for providing financial support under CSIRSRF Fellowship. The XII Five Year Plan project MULTIFUN (CSC-0101) of CSIR, India, is acknowledged for PPMS facility. P. Kamaraj is acknowledged for his help during magnetic measurements. The authors thank Central Instrumentation Facility, CECRI, Karaikudi.
particle size 50−70 nm) under identical conditions. Cyclic voltammetric (CV) analyses were performed at a scan rate of 0.1 mV s−1 between 0.001 and 3.0 V (Li+/Li). Electrochemical impedance spectroscopic studies were carried out with a perturbation of 20 mV amplitude over a frequency range of 400 kHz to 50 mHz.
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ASSOCIATED CONTENT
* Supporting Information S
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsomega.7b01294. XRD patterns and FESEM images for predicting the growth mechanism of Co3O4 HMSs; FT-IR, Raman, and thermogravimetric analysis of Co/C and Co3O4 HMSs; table comparing literature data of various Co3O4-based anode materials for Li-ion batteries (PDF)
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REFERENCES
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
Narayanan Lakshminarasimhan: 0000-0001-5986-2578 Kannadka Ramesha: 0000-0003-3089-0406 Notes
The authors declare no competing financial interest. 7655
DOI: 10.1021/acsomega.7b01294 ACS Omega 2017, 2, 7647−7657
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