Feasible Defect Engineering by Employing MOF Templates into One

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Feasible Defect Engineering by Employing MOF Templates into One-Dimensional Metal Oxides for Battery Applications Jun Young Cheong, Won-Tae Koo, Chanhoon Kim, Ji-Won Jung, and Il-Doo Kim ACS Appl. Mater. Interfaces, Just Accepted Manuscript • Publication Date (Web): 04 Jun 2018 Downloaded from http://pubs.acs.org on June 4, 2018

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ACS Applied Materials & Interfaces

Feasible

Defect

Engineering

by

Employing

MOF

Templates into One-Dimensional Metal Oxides for Battery Applications

Jun Young Cheong,† Won-Tae Koo,† Chanhoon Kim, Ji-Won Jung, and Il-Doo Kim*

Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Republic of Korea * E-mail address: [email protected] (I. D. Kim) †

These authors contributed equally to this work.

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Abstract Facile synthesis of rationally designed nanostructured electrode materials with high reversible capacity is highly critical to meet ever-increasing demands for lithium-ion batteries (LIBs). In this work, we employed defect engineering by incorporating metal organic framework (MOF) templates into one-dimensional nanostructures by simple electrospinning and subsequent calcination. The introduction of Co based zeolite imidazole framework (ZIF-67) resulted in the abundant oxygen vacancies, which induce not only more active sites for Li storage but also enhanced electrical conductivity. Moreover, abundant mesoporous sites are formed by the decomposition of ZIF-67, which are present both in and outside of the resultant SnO2Co3O4 NFs. Attributed to the creation of vacancy sites along with the synergistic effects of SnO2 and Co3O4, SnO2-Co3O4 NFs exhibit an excellent reversible capacity for 300 cycles (1287 mAh g–1 at a current density of 500 mA g–1) along with superior rate capabilities and improved initial coulombic efficiency (I.C.E) compared with pristine SnO2 NFs. This is an early report on utilizing MOF structure as the defect formation platform into one-dimensional nanostructures, which is expected to result in superior electrochemical performances required for advanced electrodes.

Keywords: nanofibers, metal-organic frameworks, anodes, Li-ion batteries, SnO2, Co3O4


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1. INTRODUCTION The ever-growing demands for energy consumption and supply propel the development of advanced energy storage system.1,2 In particular, lithium-ion battery (LIB) has been proposed as the feasible alternative to meet such demands, since its first commercialization in the 1990s.3,4 As energy consumption continues to increase, the electrode materials for LIB with higher Li storage are highly sought after; in particular, a number of group IV elements (such as Si, Ge, and Sn) have been researched as potential anode materials for LIB.5–8 Nevertheless, the extreme degree of volume changes induce rapid structural degradation, which results in quick capacity fading.9 On the other hand, a number of conversion-based metal oxides have been constantly researched as electrodes for LIB,10–13 which usually exhibit higher theoretical capacity than currently used graphite (372 mAh g–1) and undergo less volume expansion compared with above group IV elements. Among various candidates, tin (IV) oxide (SnO2) has been considered as the promising anode material as it can reach the theoretical capacity of 1493 mAh g–1 when its reaction becomes completely reversible, along with other merits such as abundance, low cost, and environmental friendliness.14,15 Nevertheless, critical limitations are also present for SnO2: still prevalent volume changes that lead to structural integrity, agglomeration of as-formed Sn, pulverization, low conductivity, and low initial coulombic efficiency. So far, to overcome these issues, the most common approach was to utilize nanostructured SnO2 with carbonaceous materials,16–18 where the nanostructuring and conductive carbon matrix may alleviate the volume changes while allowing better electronic conductivity. However, such adopted approach has a critical disadvantage as it sacrifices the high theoretical capacity that SnO2 can bring by introducing carbon, which leads to limited enhancement in the reversible capacity.



To maximize the loading amount of SnO2 while alleviating the volume stresses, various forms of SnO2-metal oxide composite electrodes were developed.19,20 This approach mainly uses SnO2 as the main active material and another metal oxide to either buffer the volume changes and/or decomposes Li2O to enhance the initial coulombic efficiency, which can bring much higher reversible capacity. Previously, the incorporation of regular SnCo alloy results in enhanced electrochemical performance, where a careful selection of metal oxide composites can bring similar effects.21 Nevertheless, in most of these cases, the conductivity issues of SnO2 had to be compensated by either using conductive materials in addition to the composite or metal oxide that exhibits much lower capacity compared with SnO2,19,20 which may also result in limited electrochemical performance. In recent years, defect engineering of electrode materials was introduced to optimize the overall electrochemical performance.22–24 Careful engineering of microstructures, such as defect engineering, can alter the surficial and interfacial properties, leading to more suitable ionic and electronic transport. As prominent examples, metal–organic frameworks (MOFs) have recently attracted significant attention not only due to their high porosity and specific surface area but also due to the presence of defects.25,26 So far, MOFs can be utilized as a template to synthesize porous metal oxide nanostructures that can be applied in various fields, including electrocatalysts,27,28 gas sensors,29,30 and LIBs.31,32 As MOF-derived materials also have high porosity and large active reaction sites, they also exhibit improved electrochemical performances as anodes for LIBs.33 For instance, Lu et al. synthesized Co3O4@Co3V2O8 hollow structures as an anode material for LIBs by using MOF templates.34 The unique porous structure of Co3O4@Co3V2O8 hollow nanoboxes enhanced the Li storage properties in terms of specific capacity, cycling stability, and rate performance. Moreover, as stated above, the introduction of MOFs can aid in the creation of additional defects. Fang and co-workers previously reported that the Co and Zn based MOF-templated composite metal oxide ACS Paragon Plus Environment

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(Co3O4/ZnO) nanosheets exhibited the enhanced performances as anodes for LIBs. They observed that the oxygen defects were generated in bimetallic metal oxides, which contributed to superior electrochemical properties.35 Nevertheless, although great efforts have been made, it still remains as challenge to easily prepare desirable MOF-derived electrode materials due to their limitations such as low structural stability and poor conductivity.36 In this work, we introduce defect-induced mesoporous SnO2-Co3O4 nanofibers (NFs) by using electrospinning with MOF templates, which can tune the defect levels of electrode materials. Since one-dimensional structures can provide increased electrochemical performances in energy storage devices due to their high surface area to volume ratio and high porosity,37,38 we employed MOF templates into 1D metal oxide nanofibers for defect engineering. Co based zeolite imidazole framework (ZIF-67), MOF template, was electrospun together with the Sn precursor and organic polymer, where the ZIF-67 was homogeneously decorated outside and embedded inside the NFs. Upon calcination, the mesoporous sites were formed on the composite NFs by the decomposition of ZIF-67, leading to various active sites, suitable for ionic and electron transport. In addition, a defect engineering with the incorporation of ZIF-67 formed not only more active sites but also oxygen vacancies in both SnO2 and Co3O4, which are expected to enhance the overall electric conductivity. Attributed to their rationally designed nano-architecture, they deliver a reversible capacity (1287 mAh g–1) after 300 cycles at a current density of 500 mA g–1, along with superior rate capabilities.

2. MATERIALS AND METHODS 2.1. Materials. 2-methylimidazole (mIM, 99.0%) and polyvinylpyrrolidone (PVP, Mw



∼ 1,300,000 g mol−1) were purchased from Aldrich. Tin(IV) acetate (Sn[CH3CO2]4), cobalt nitrate hexahydrate ([Co(NO3)2·6H2O], 98%), N,N-dimethylformamide (DMF, 99.8%), ethanol (99.5%), and methanol (99.9%) were purchased from Sigma-Aldrich. 2.2. Synthesis of SnO2 NFs. To prepare the electrospinning solution for SnO2 NFs, 0.25 g of tin(IV) acetate was dissolved in 1.35 g of ethanol, and 0.35 g of PVP was dispersed in 1.35 g of DMF. Then, two solutions were mixed and stirred at room temperature for 2 h using a magnetic bar. The prepared solution was electrospun at a voltage of 15 kV with a feeding rate of 0.1 mL min−1. The syringe needle was away from the collector by 15 cm. Temperature and humidity were maintained at 25 °C and 40%. Finally, SnO2 NFs were obtained after the calcination of as-spun NFs at 600 °C (5 °C m−1) for 1 h in air atmosphere. 2.3. Synthesis of SnO2-Co3O4 NFs. Firstly, ZIF-67 was prepared by using precipitation reaction. 0.20 g of Co(NO3)2·6H2O and 0.40 g of mIM were dissolved in 20 mL of methanol, respectively. Then, the solutions were rapidly mixed and precipitated at room temperature for 6 h. Then, ZIF-8 was purified by using a centrifugation and dried at 50 °C for 1 day. Subsequently, ZIF-8 was dispersed in 1.35 g of ethanol solution including 0.25 g of tin(IV) acetate. The loading amount of ZIF-8 was controlled to 30, 60, and 120 mg. The suspension was mixed with 0.35 g of PVP dissolved in 1.35 g of DMF, and stirred at room temperature for 2 h. Then, electrospinning was carried out at a voltage of 16 kV with a feeding rate of 0.1 mL min−1. The distance between the syringe needle (25 gauge) and the stainless steel collector was maintained at 15 cm. Then, as-spun NFs were calcined at 600 °C (5 °C min–1) for 1 h in air atmosphere to produce SnO2-Co3O4 NFs. 2.4. Synthesis of SnO2-Co3O4 NFs without MOF templates. As a control sample, we prepared SnO2-Co3O4 NFs by using Sn and Co precursors. 0.25 g of tin(IV) acetate and 0.08 g of Co(NO3)2·6H2O were dissolved in 1.35 g of ethanol, and 0.35 g of PVP was dispersed in


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1.35 g of DMF. Then, the above solutions were rapidly mixed and homogeneously stirred at room temperature for 2 h. The electrospinning was conducted by using same synthetic condition and process. Then, SnO2-Co3O4 NFs without MOF templates were produced after the same calcination step (600 °C for 1 h in air). 2.5. Material characterization. The microstructures and morphologies of the samples were analyzed by scanning electron microscopy (SEM, SU5000, Hitachi) and transmission electron microscopy (TEM, Tecnai G2 F30 S-Twin, FEI). The crystal structure was investigated by X-ray diffraction (XRD, SmartLab, Rigaku) using Cu Kα radiation (λ = 1.5418 Å). The chemical binding state of the samples was verified by X-ray photoelectron spectroscopy (XPS, Sigma Probe, Thermo VG Scientific). The molar ratio of Sn and Co elements in SnO2-Co3O4 NFs was confirmed by the inductively coupled plasma optical emission spectrometry (ICP-OES). To analyze Brunauer-Emmett-Teller (BET) surface area and the pore distribution of the samples, N2 adsorption/desorption isotherm was carried out at 77 K (Tristar 3020, Micromeritics). Thermal behaviors of the samples were investigated by thermogravimetric analysis (Labsys Evo, Setaram). The photoluminescence (PL) analysis was conducted by using an excitation wavelength of 325 nm (LabRAM HR Evolution Visible_NIR, HORIBA). 2.6. Electrochemical measurement. To carry out the electrochemical cell tests, 2032 coin-type half cells were assembled inside the glove box. The electrode materials were composed of 80 wt% of active materials (in our case, SnO2-Co3O4 NFs and SnO2 NFs), 10 wt% of carbon black, and 10 wt% of binder composed of poly(acrylic acid) (PAA)/sodium carboxymethyl cellulose (CMC) (wt%/wt% = 50/50). The weight percentage of active materials (SnO2-Co3O4 NFs and SnO2 NFs) is 80 wt% and the capacity was calculated based on the mass of the active materials. They were slurry casted on the Cu foil and dried for 10



min to be used as a component of the electrochemical cells. Before the electrochemical cell was fabricated, it was further dried in vacuum at 150 °C for 2 h. In the cell assembly, Li metal was used as the counter electrode and Celgard 2325 was used as a separator. Moreover, for an electrolyte, 1.3 M lithium hexafluorophosphate (LiPF6) dissolved in a solvent mixture of ethylene carbonate/diethylene carbonate (EC/DEC, v/v = 3/7) with 10 wt% of fluoroethylene carbonate (FEC) (PANAX ETEC) was used. All of the electrochemical cell tests were carried out at a voltage window of 0.005 to 3 V for formation cycle (current density: 50 mA g–1) and 0.01 and 3 V for other current densities. The cyclic voltammetry (CV) was conducted at a scan rate of 0.1 mV s–1 within a range of 0.01 and 3 V using battery testing device (WBCS4000, Wonatech). Impedance tests were carried out using the AC impedance analyzer (ZIVE SP1, Wonatech).

3. RESULTS AND DISCUSSION The synthetic process of SnO2-Co3O4 NFs derived from MOF templates was illustrated in Figure 1. Firstly, ZIF-67, consisting of Co ions and mIM, was synthesized by using room temperature precipitation, as previously reported.39 SEM image of ZIF-67 shows polyhedron structure of ZIF-67 with an average diameter of 300 nm (Figure 2a), and XRD analysis confirms the crystal structure of ZIF-67 (Figure S1). Then, electrospinning was carried out to fabricate one-dimensional structure. As a template, ZIF-67 was dispersed in the DMF/ethanol (1:1) solution dissolved PVP and tin(IV) acetate. Then, the electrospinning of the above solution produced as-spun tin(IV) acetate/PVP NFs decorated by ZIF-67. After calcination at 600 ºC for 1 h, tin(IV) acetate/PVP/ZIF-67 NFs were transformed to SnO2-Co3O4 NFs. During the calcination, the metal species in the composite NFs were oxidized, and the organic compounds were completely decomposed. In particular, a number of meso-pores (2–50 nm)


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were generated on NFs by the decomposition of organic ligands in ZIF-67. The morphology and microstructure of the sample were analyzed by using SEM and TEM. The SEM image of as-spun tin(IV) acetate/PVP/ZIF-67 NFs revealed the dense fibrous structure with the average diameter of 250 nm (Figure 2b). Notably, some parts were protruded from the surface of as-spun NFs due to exposed ZIF-67 templates (green arrows in Figure 2b). On the other hand, without ZIF-67 templates, tin(IV) acetate/PVP NFs exhibited smooth fibrous structure (Figure S2a). The thermogravimetric analysis of tin(IV) acetate/PVP/ZIF-67 NFs confirmed that the as-spun composite NFs were completely oxidized to SnO2-Co3O4 NFs at 600 °C (Figure S3). Thus, the calcination of the samples at 500 °C showed low crystallinity of SnO2-Co3O4 NFs, and the calcination at 700 °C showed increased crystallinity but caused the grain growth, along with decreased oxygen vacancies and partial collapse of NFs (Figure S4). The as-spun composite NFs were completely converted to SnO2-Co3O4 NFs with an average diameter of 150 nm after the calcination at 600 °C (Figure 2c). The decomposition of organic compounds caused the shrinkage of the NFs. The SEM image showed the rough surface of SnO2-Co3O4 NFs and some protruding parts on NFs. However, the calcination of as-spun NFs without ZIF-67 templates, the calcination at same condition only produced rough surface of pristine SnO2 NFs (Figure S2b). In addition, TEM image clearly exhibited the microstructure of SnO2-Co3O4 NFs (Figure 2d). The dark points which are like polyhedrons in SnO2-Co3O4 NFs were Co3O4 cubes obtained by the calcination of ZIF-67. In addition, a lot of meso-pores were observed in SnO2-Co3O4 NFs in the TEM image. The high-resolution TEM (HRTEM) image revealed a mesoporous structure and a Co3O4 polyhedron embedded in SnO2 NF (Figure 2e). The crystal plane of Co3O4 (220) and SnO2 (110), which were corresponded to the lattice fringe of 2.44 Å and 3.35 Å respectively, was also observed in the HRTEM image (Figure 2f). The fast Fourier transform (FTT) images of red and green dotted boxes in Figure 2f clearly showed the lattice ACS Paragon Plus Environment


fringe of SnO2 (110) and Co3O4 (220) (Figure S5). Selective area electron diffraction (SAED) patterns of SnO2-Co3O4 NFs confirmed the polycrystalline properties of SnO2 with crystal planes of (110), (101), (211), and (301) (Figure 2g). However, the intensity of a crystal plane of Co3O4 (311) and (411) were weak due to the relatively low amounts of Co3O4 in SnO2Co3O4 NFs. In addition, the EDS mapping images of SnO2-Co3O4 NFs revealed the welldispersed Co3O4 polyhedrons in SnO2 NFs (Figure 2h). The EDS line profiles showed two divided peaks of Co intensity in the vertical line of an SnO2-Co3O4 NF (Figure 2i), demonstrating the functionalization of Co3O4 polyhedrons in SnO2 NFs. The crystal structure of the samples was verified by XRD analysis. To analyze more detail, we prepared SnO2-Co3O4 NFs obtained from the electrospinning solution containing 30, 60, and 120 mg of ZIF-67 (Here after, they are denoted as SnO2-Co3O4 NFs_30, SnO2Co3O4 NFs_60, and SnO2-Co3O4 NFs_120). The SEM images of as-spun NFs and calcined NFs with different loading amounts of ZIF-67 showed similar morphologies, respectively (Figure S6 and S7). In addition, we confirmed the exact loading amounts of Co in SnO2Co3O4 NFs by using ICP-OES analysis. The content of Co in SnO2-Co3O4 NFs_30, SnO2Co3O4 NFs_60, and SnO2-Co3O4 NFs_120 was 8.5, 17.2, and 34.1 wt%, respectively. The XRD patterns of the samples revealed the tetragonal SnO2 (JCPDS no. 41-1445) with polycrystalline structure (Figure 3a). In addition, the weak intensity of Co3O4 (311) was observed in the XRD data of SnO2-Co3O4 NFs_60 and SnO2-Co3O4 NFs_120, while the peaks of Co3O4 were not observed in XRD data of SnO2-Co3O4 NFs_30. From the XRD results, we calculated the crystal size of SnO2 by using Scherrer equation. Scherrer equation is defined as D = Kλ/βcos(θ), where D is an average grain size at a specific peak (θ), K is a shape factor (typically 0.9), λ is a wavelength of X-ray sources (λ = 0.154 nm for Cu Kα,), and β is a line broadening at half-maximum of a specific peak. The average grain size of SnO2 obtained from the major peaks of (110), (101), and (211) plane is 51.8 nm for SnO2 NFs, ACS Paragon Plus Environment

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12.4 nm for SnO2-Co3O4 NFs_30, 7.9 nm for SnO2-Co3O4 NFs_60, and 22.6 nm for SnO2Co3O4 NFs_120. The average crystal size of SnO2 in SnO-Co3O4 NFs was largely reduced compared to that of SnO2 NFs, indicating that the ZIF-67 template in Sn precursor/PVP NFs inhibit the crystal growth of SnO2 during the calcination. In addition, the restraint of SnO2 grain growth was not effective when excessive amounts of ZIF-67 were loaded on SnO2 NFs, and SnO2-Co3O4 NFs_60 showed smallest grain size among the samples. Furthermore, we calculated the lattice fringes of the samples from XRD peaks. As a result, the samples showed almost similar lattice distances and small variations (lower than 0.037 Å) for major three planes of SnO2 (110), (101), and (211) (Table S1). However, the lattice spacing of SnO2Co3O4 NFs was slightly decreased compared with that of pristine SnO2 NFs. Since oxygen vacancies in metal oxides can cause lattice distortion by the chemical expansion,40–42 the refinement of the XRD results confirmed the oxygen vacancies in the samples. The chemical binding state of the samples was analyzed by using XPS analysis. The high-resolution XPS data in the victim of Sn 3d for SnO2 NFs exhibited chemical binding states of Sn4+ with characteristic peaks at 495.0 eV for 3d3/2 and 486.5 eV for 3d5/2 (Figure 3b), which correspond to previous studies on the chemical state of SnO2.43 On the other hand, when Co3O4 polyhedrons were decorated on SnO2 NFs, Sn 3d peaks shifted toward low binding energy. The peak of Sn 3d5/2 was located at for 486.0 eV for SnO2-Co3O4 NFs_30, 486.1 eV for SnO2-Co3O4 NFs_60, and 486.2 eV for SnO2-Co3O4 NFs_120. Similarly, the high-resolution XPS analysis of Co 2p revealed the peak shift depending on the loading amounts of Co3O4 (Figure 3c). Co atoms in Co3O4 have two chemical state of Co2+ and Co3+. The binding energy of Co3+ 2p3/2 for SnO2-Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2Co3O4 NFs_120 is 779.9 eV, 780.3, and 780.4 eV, respectively. The peak shift of Co 2p3/2 reveals the change of electronic state of Co, which means that the formation of oxygen vacancies.35 The high-resolution XPS spectra of O 1s showed the chemical state of O2– for ACS Paragon Plus Environment


lattice oxygen, O– for oxygen with defects sites, and O2– for physisorbed oxygen (Figure 3d).44 To compare oxygen vacancies of the samples, the relative area ratio of oxygen deficiency to lattice oxygen (O–/O2–) was calculated for all samples (Figure S8). As a result, SnO2-Co3O4 NFs_60 exhibited higher ratio (0.89) than other samples (0.39 for SnO2 NFs, 0.68 for SnO2-Co3O4 NFs_30, and 0.76 for SnO2-Co3O4 NFs_120). In addition, the oxygen peaks of O 1s were shifted to the lower binding energy, similar to Sn 3d and Co 2p. The XPS results apparently indicate that the functionalization of Co3O4 polyhedrons using MOF templates increase the oxygen vacancies in SnO2-Co3O4 NFs, and the oxygen defect levels in composite NFs are easily tuned by changing the loading amounts of MOF templates. Although the formation mechanism of oxygen vacancies by using MOF templates is not yet known, oxygen defect sites might be caused by the highly porous structure of MOF as previously reported.35 In addition, we carried out PL analysis to further investigate the oxygen vacancies of the samples. The PL spectrum of SnO2-Co3O4 NFs exhibited the dominant broad peak at around 560 nm and the weak peak at 440 nm, while that of pristine SnO2 NFs exhibited only broad peak around 560 nm. The broad peak around 560 nm can be attributed to oxygen vacancies in the band gap of SnO2,45,46 and the peak at 440 nm can be explained by oxygen vacancies in the SnO2 valence band and impurities in the grain boundary layer of SnO2.47 Thus, the PL spectrum of the samples indicated that the incorporation of MOF templates into SnO2 NFs generated additional oxygen vacancies in SnO2. To investigate BET surface area and porous structure of the samples, we carried out N2 adsorption/desorption isotherms at 77 K (Figure S10a). The isotherms of SnO2-Co3O4 NFs exhibited type II isotherms that reveal mesoporous structure, while that of SnO2 NFs revealed non-porous structure. In addition, the specific surface area obtained by BET methods is 9.00 m2 g–1 for SnO2 NFs, 29.27 m2 g–1 for SnO2-Co3O4 NFs_30, 38.34 m2 g–1 for SnO2-Co3O4 ACS Paragon Plus Environment

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NFs_60, and 36.59 m2 g–1 for SnO2-Co3O4 NFs_120. Co3O4-SnO2 NFs exhibited 3- or 4-fold higher surface area than SnO2 NFs. The pore size distribution of each sample was also obtained by Barrett-Joyner-Halenda (BJH) method (Figure S10b). SnO2-Co3O4 NFs showed increased volume of meso-pores in the range of 3–30 nm. In particular, the cumulative pore volume of SnO2-Co3O4 NFs_30 (0.1453 cm3 g–1), SnO2-Co3O4 NFs_60 (0.1431 cm3 g–1), and SnO2-Co3O4 NFs_120 (0.1156 cm3 g–1) is much larger than that of pristine SnO2 NFs (0.0540 cm3 g–1). Since the decomposition of organic ligands in MOF generates numerous mesopores in the composites NFs,29 SnO2-Co3O4 NFs have higher reaction sites and larger pore volume than pristine SnO2 NFs. This meso-porous structure of SnO2-Co3O4 NFs is beneficial for Li ion diffusion and electrolyte penetration. To clearly verify the effect of defect engineering of MOF templates, we prepared SnO2Co3O4 NFs without MOF templates (SnO2-Co3O4 NFs_w/o MOF) by using the electrospinning of Sn precursors and Co precursors. The Co contents in SnO2-Co3O4 NFs were controlled to the same levels of that (17.2 wt%) in SnO2-Co3O4 NFs_60. In contrast with MOF templated SnO2-Co3O4 NFs, the SEM image of as-spun composite NFs and SnO2Co3O4 NFs_w/o MOF showed smooth surface of fibrous structures without protruding parts (Figure S11). The XRD analysis of SnO2-Co3O4 NFs_w/o MOF exhibited the dominant phase of tetragonal SnO2 with the minor phase of spinel Co3O4 (Figure S12a). However, from the XRD result, the grain size of SnO2-Co3O4 NFs was calculated to 32.1 nm, which is larger than that (7.9 nm) of SnO2-Co3O4 NFs_60. In addition, the XRD peaks of SnO2-Co3O4 NFs_w/o MOF were shifted to higher angles than those of SnO2-Co3O4 NFs, due to the decreased amounts of vacancies and interstitial dopants in SnO2.41,42 The peak shift of Sn, Co, and O was also observed in the XPS analysis of SnO2-Co3O4 NFs_w/o MOF (Figure S12b,c,d), while the relative ratio of oxygen deficiency to lattice oxygen (O–/O2–) was decreased to 0.47 compared to that (0.89) of SnO2-Co3O4 NFs_60 (Figure S8), indicating the ACS Paragon Plus Environment


decrease of oxygen vacancies compared with SnO2-Co3O4 NFs derived from MOF templates. In addition, N2 isotherms of the sample without MOF templates revealed the decreased surface area (14.31 m2 g–1) and pore volumes (0.0458 m3 g–1) (Figure S13). These results demonstrated that the MOF templating route easily generated numerous meso-pores as well as defect sites in SnO2 NFs. To understand the effect of introducing MOF templates to modulate the defect levels of the electrode materials, their electrochemical performance in the LIB cell was further investigated. The redox reactions of SnO2-Co3O4 NFs_60, as a practical example, have been analyzed by using cyclic voltammetry (CV) (Figure 4a). In the first cycle, two cathodic peaks were present at 0.5 and 0.1 V, which are ascribed to the initial conversion reaction (formation of solid electrolyte interphase (SEI) layer) and alloying reaction of Sn, in accordance with the previous work.20 During the charge process, three anodic peaks were present at 0.5 V, 1.25 V, and 1.9 V, which are ascribed to the de-alloying of LixSn to Sn, oxidation of Sn to SnO, and subsequent oxidation of SnO to SnO2, in accordance with the previous work.20,48 After the 1st cycle, the broad cathodic peak is shown clearly at 1.25 V, and no sharp peak related to the alloying reaction of Sn is shown. Although no distinct sharp peak is present for the electrochemical reaction of Co3O4 with Li, the broad peak at 1.5 to 0.75 V indicate that both the conversion of SnO2 and Co3O4 takes place. To examine the initial irreversible capacity loss along with the overall reaction pathway, the charge and discharge profiles in the formation cycle (at a current density of 50 mA g–1) were shown for SnO2-Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2-Co3O4 NFs_120. To understand the irreversible capacity loss, initial coulombic efficiency (I.C.E) was calculated for SnO2-Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2-Co3O4 NFs_120. The I.C.E of SnO2-Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2-Co3O4 NFs_120 was 64.9%, 71.1%, and 72.1%, which shows a constantly improved value with respect to the amount of Co3O4. Such a noticeable trend can be ACS Paragon Plus Environment

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attributed to the introduction of Co3O4 that leads to the enhanced reversibility in the initial cycle through the reduction of Li2O through the as-formed Co upon lithiation.49,50 The effect of Co3O4 introduction in enhancing I.C.E becomes minimal when SnO2-Co3O4 NFs_60 and SnO2-Co3O4 NFs_120 are compared, where the suitable MOF concentration exists for enhancing the overall reversibility of Li storage. The cytcle retention characteristics of SnO2Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2-Co3O4 NFs_120 were further examined at a current density of 500 mA g–1 for 300 cycles (Figure 4c). The corresponding charge and discharge profiles for Co3O4 NFs_30, SnO2-Co3O4 NFs_60, and SnO2-Co3O4 NFs_120 were also plotted for the 2nd cycle, 10th cycle, 50th cycle, and 100th cycle (Figure S14). Based on the results, SnO2-Co3O4 NFs_60 exhibits the most stable cycling performance among various SnO2-Co3O4 NFs samples, with good reversibility as evidenced by its coulombic efficiency. Although SnO2-Co3O4 NFs_120 exhibit higher I.C.E in the initial cycle, it does not possess better cycling stability compared with SnO2-Co3O4 NFs_60. It can be suggested that relatively smaller grain sizes along with more active sites for Li storage contributed to the stable cycle retention characteristics of SnO2-Co3O4 NFs_60.51 Moreover, it is possible that oxygen vacancies could trigger local build-in electric field, which promotes Li ion transport by coulombic force.52 Moreover, compared with SnO2-Co3O4 NFs_60, SnO2-Co3O4 NFs_120 exhibits more rapid capacity increase and sudden decay due to the higher loading amount of Co3O4. Steady capacity increase in the initial cycles is observed for SnO2-Co3O4 NFs_60, which can be attributed to the formation of gel-like polymeric film in the surface of Co.53 In terms of rate capabilities (Figure 4d), SnO2-Co3O4 NFs_60 also maintains sustained capacity even at a higher current density (5.0 A g–1) compared with Co3O4 NFs_30 and SnO2-Co3O4 NFs_120. Attributed to the formation of numerous defect sites arising from the oxygen vacancies, SnO2-Co3O4 NFs_60 exhibits not only excellent cycle retention but also good rate capabilities.



SnO2-Co3O4 NFs_60 was chosen as the representative SnO2-Co3O4 NFs sample to compare its superior electrochemical performance compared with pristine SnO2 NFs and SnO2-Co3O4 NFs_w/o MOF (Figure 5). To truly understand the effect in introducing MOF templates, it is important to compare the electrochemical performance of SnO2-Co3O4 NFs_60 not only with SnO2 but also with SnO2-Co3O4 NFs_w/o MOF. Based on the comparison of CV curves of SnO2-Co3O4 NFs_60 and those of SnO2 NFs (Figure S15), it can be seen that the combination of SnO2 and Co3O4 resulted in the broad CV peaks, where different redox reactions with Li simultaneously take place. The CV curves of SnO2-Co3O4 NFs_w/o MOF are also presented in Figure S16, exhibiting sharp peaks. To examine the overall electrochemical characteristics in the initial cycle, the charge and discharge profile of SnO2 NFs, SnO2-Co3O4 NFs_w/o MOF, and SnO2-Co3O4 NFs_60 in the formation cycle (50 mA g–1) was compared (Figure 5a). Although the discharge capacity was almost identical for SnO2 NFs and SnO2-Co3O4 NFs_60, significant difference in the charge capacity was observed, where SnO2 NFs exhibit less reversible capacity (852.2 mAh g–1) compared with SnO2-Co3O4 NFs_60 (1189.3 mAh g–1). More importantly, lower I.C.E (67.8%) was observed for SnO2 NFs. The addition of Co3O4 contributed to the improved capacity, which stems from the synergistic effects of SnO2 and Co3O4. Similarly, SnO2-Co3O4 NFs_w/o MOF also exhibited lower charge capacity compared with SnO2-Co3O4 NFs_60. Two other noticeable differences are observed in the charge and discharge profile; during the discharge process, SnO2-Co3O4 NFs_60 starts storing Li at voltage below 1.5 V whereas SnO2 NFs mainly participate in the reaction with Li at voltage below 1.0 V. Additionally, the introduction of Co3O4 results in additional capacity increase after the 2.0 V, leading to an additional contribution to the reversibility of electrochemical reactions. The cycle retention characteristics were further compared for SnO2-Co3O4 NFs_60, SnO2-Co3O4 NFs_w/o MOF, and SnO2 NFs at a current density of 500 mA g-1 (Figure 5b). For reference, charge and


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discharge profile of SnO2 NFs was also shown for the 2nd cycle, 10th cycle, 50th cycle, and 100th cycle (Figure S17). As can be seen in Figure S17, both the charge and discharge profiles quickly diminish after the 100th cycle, as evidenced by the low capacity that is caused by the rapid capacity fading. Although SnO2 NFs exhibit reversible capacity comparable to the SnO2-Co3O4 NFs_60 up to 40 cycles, it quickly fades after the 40 cycles. Not only capacity but also the coulombic efficiency quickly fades after 40 cycles, attributed to the structural degradation of the overall electrode materials, needing further investigation in the subsequent analyses. Similarly, SnO2-Co3O4 NFs_w/o MOF also undergoes steady capacity fading for 100 cycles, with less reversible reaction with Li, as evidenced by the lower coulombic efficiency. Finally, the rate capabilities of SnO2 NFs and SnO2-Co3O4 NFs_w/o MOF were also further compared with SnO2-Co3O4 NFs_60 at different current densities (expressed in A g–1) (Figure 5c). Although SnO2 NFs exhibit comparable capacity at low current densities, it quickly fades and reaches very low capacity (less than 100 mAh g–1) at high current density (5.0 A g–1). By the introduction of MOF templates to create mesoporous Co3O4 on the inner and outer sites of SnO2 NFs, it can allow faster ionic and electronic transport, leading to higher Li storage at high current density. Compared with SnO2-Co3O4 NFs_w/o MOF, SnO2-Co3O4 NFs_60 also exhibits slightly enhanced capacity at high current density (5.0 A g–1). The overall electrochemical performance of SnO2-Co3O4 NFs_60 was compared with various SnO2-Co3O4 composites in the previous literatures (Table S2), where its performance was comparable or superior to the previously reported SnO2-Co3O4 nanostructures. SnO2-Co3O4 NFs_60 exhibits higher capacity (1287 mAh g–1) for longer cycles (300 cycles) even at more harsh current density (0.5 A g–1) compared with previously reported SnO2-Co3O4, which highlights the advantages of using MOF templates into one-dimensional metal oxides to enhance Li storage characteristics. To understand the effect of MOF templates in enhancing the Li diffusivity, cyclic ACS Paragon Plus Environment


voltammetry at various scan rates (0.1, 0.5, 1, and 10 mV s–1) was conducted for SnO2-Co3O4 NFs_60 and SnO2 NFs (Figure 6a and 6b). Based on the CV curves, the linear fitting can be conducted with respect to the various scan rates for both samples (Figure 6c and 6d). Randles-Sevcik equation54 can be employed to calculate the Li diffusivity, which is written as Ip = 0.4463(F3/RT)1/2n3/2AD1/2Cv1/2, where Ip is the peak current, F Faraday’s constant, R gas constant, T absolute temperature, n number of electrons involved in the redox reaction, A the electrode area, D Li ion diffusion constant, C the shuttle concentration, and v the scan rates. The equation can be further reduced to D1/2 = Ip/(0.4463(F3/RT)1/2n3/2ACv1/2) = K/(0.4463(F3/RT)1/2n3/2AC) when K is considered as a slope of the linear fitting curves, as K = Ip/v1/2. Based on the values of K, it has been shown that the Li diffusivity of the SnO2Co3O4 NFs_60 is 1.11 × 10–11 cm2 s–1, which is higher than that of SnO2 NFs (1.0 × 10–11 cm2 s–1). Clearly, it has been demonstrated that the Li diffusivity has been enhanced by the introduction of MOF templates, which bring a number of oxygen vacancies that lead to enhanced electronic conductivity. Furthermore, based on the CV curves, the Li storage mechanism of SnO2-Co3O4 NFs_60 was further analyzed by calculating the proportion of diffusion-controlled intercalation behavior and surface-induced capacitive behavior. The relationship between current (i) and the scan rate (v) can be expressed in power law (i = avb), where a and b are the arbitrary constants.55 The equation can be expanded to ic = k1v + k2v1/2 and ic/v1/2 = k1v1/2 + k2, where k1v is a capacitive process and k2v1/2 a diffusion-controlled process. Based on these equations, the contribution of diffusion-controlled intercalation and capacitive behavior were calculated and expressed in graphs for various scan rates (Figure S18). With increasing scan rates, the capacitive behavior becomes more dominant, where the high-rate lithiation will induce more capacitive behavior for SnO2-Co3O4 NFs_60. To further delve into the effect of using MOF templates to combine mesoporous Co3O4 into SnO2 NFs, the electronic conductivity arisen from the MOF templates was analyzed. To


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compare the electronic conductivity of SnO2-Co3O4 NFs_60 and SnO2 NFs, the impedance tests were carried out, based on the previously suggested equivalent circuit model (Figure S19a).56 Based on the semicircle areas in the Nyquist plots, the charge transfer resistance can be measured and compared. After the 1st cycle (Figure S19b), it can be seen that both SnO2Co3O4 NFs_60 and SnO2 NFs show minimal charge transfer resistance, which can be calculated by the semicircle areas in the middle frequency regions. However, after the 300th cycle (Figure S19c), there is a difference in the charge transfer resistance between SnO2Co3O4 NFs_60 and SnO2 NFs. The charge transfer resistance of SnO2 NFs is much larger than that of SnO2-Co3O4 NFs_60 – this can be attributed to the introduction of Co3O4 that leads to the reversible decomposition of Li2O and stable cycling characteristics up to 300 cycles. On the other hand, SnO2 NFs undergo severe capacity decay after the 100th cycle, which resulted in high charge transfer resistance. The Nyquist plots clearly demonstrate that the introduction of mesoporous Co3O4 into SnO2 NFs resulted in less charge transfer resistances compared with SnO2 NFs, which resulted in the suitable ionic and electronic transport, helpful for reversible Li storage. Furthermore, it was revealed that not only the improvement in electronic conductivity but also that in structural integrity was achieved by introducing mesoporous Co3O4 into SnO2 NFs. The postmortem analysis (ex situ SEM images) of SnO2-Co3O4 NFs_60 and SnO2 NFs was carried out after the 300th cycle (Figure 7a and 7b). It demonstrates that the structural integrity was maintained for SnO2-Co3O4 NFs_60, which can be attributed to the many mesoporous Co3O4 sites that can alleviate the volume expansion while preventing the agglomeration of as-formed Sn nanoparticles. On the other hand, SnO2 NFs underwent severe structural degradation and subsequent pulverization, which resulted in the rapid capacity fading. The schematic illustration of overall reaction pathway for both SnO2 NFs and SnO2Co3O4 NFs_60 is further presented (Figure 7c and 7d). SnO2-Co3O4 NFs_60 exhibits faster Li ACS Paragon Plus Environment


ion diffusivity compared with SnO2 NFs and maintains the structural integrity, whereas SnO2 NFs undergo agglomeration and pulverization, leading to the loss of electric contact. Based on electrochemical analyses and ex situ analysis, the introduction of Co3O4 led to both enhanced electronic conductivity and structural integrity, which led to excellent electrochemical performance. In addition, the proposed strategy of defect engineering via MOF templating routes can be easily extend to other materials by simply changing types of metal precursors in electrospinning solution, and MOF templates, such as ZIF-8, UiO-66, MIL-101, and MOF-74.

4. CONCLUSION In this work, we have employed feasible defect engineering of metal oxide nanofibers by incorporating MOF templates into one-dimensional nanostructures. Hybrid one-dimensional SnO2-Co3O4 NFs were successfully synthesized with MOF templates by simple electrospinning and calcination. The introduction of MOF templates further produced numerous mesoporous sites upon decomposition and oxygen vacancy sites. By controlling the amount of MOF, the defect levels of SnO2-Co3O4 NFs were modulated. Based on electrochemical and postmortem analyses, both the electronic conductivity and structural integrity were greatly improved by the introduction of MOF templates, arising from the presence of oxygen vacancy sites and mesoporous sites that alleviate the volume changes. Such improvement also led to the enhanced Li diffusivity characteristics, indicating faster ionic transport. The SnO2-Co3O4 NFs_60 exhibits excellent cycle retention (1287 mAh g–1 after 300 cycles at 500 mA g–1), with enhanced rate capabilities and reversibility compared


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with SnO2 NFs. The defect engineering of various composite metal oxides using MOF in one-dimensional nanostructures are simple and feasible, which can also be applied to other fields such as electrocatalysis.

ASSOCIATED CONTENT Supporting Information. Additional characterizations and sensing characteristics. These materials are available free of charge via the Internet at “http://pubs.acs.org.” AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Notes The authors declare no competing financial interest.

ACKNOWLEDGEMENT This work was supported by the Wearable Platform Materials Technology Center (WMC) funded by the National Research Foundation of Korea (NRF) grant of the Korean Government (Ministry of Science, ICT & Future Planning) (No. 2016R1A5A1009926). This work was also supported by Korea CCS R&D Center (KCRC) grant funded by the Korea government (Ministry of Science, ICT & Future Planning) (No. NRF-2014M1A8A1049303) and the National Research Foundation of Korea (NRF), grant no. 2014R1A4A1003712 (BRL Program).



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Figure 1. Schematic illustration of synthesis process of SnO2-Co3O4 NFs.


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Figure 2. SEM images of: (a) ZIF-67 templates, (b) as-spun tin(IV) acetate/PVP/ZIF-67 NFs, and (c) SnO2-Co3O4 NFs. (d) TEM image of SnO2-Co3O4 NFs, (e,f) HRTEM image of SnO2Co3O4 NFs, (g) SAED patterns of SnO2-Co3O4 NFs, (h) STEM image and elemental mapping images of SnO2-Co3O4 NFs, and (i) EDS line profile of SnO2-Co3O4 NFs to yellow line in Figure 2h.



Figure 3. (a) XRD analysis of SnO2-Co3O4 NFs, and XPS analysis using high resolution spectra of SnO2-Co3O4 NFs in the vicinity of (b) Sn, (c) Co, and (d) O.


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Figure 4. (a) CV analysis of SnO2-Co3O4 NFs_60 in the 1st, 2nd, and 3rd cycle, (b) Charge and discharge profile of SnO2-Co3O4 NFs samples in the formation cycle (50 mA g–1), (c) Cycle retention tests of SnO2-Co3O4 NFs samples (500 mA g–1), and (d) Rate capabilities of SnO2-Co3O4 NFs samples (expressed in A g–1).



Figure 5. (a) Charge and discharge profile of SnO2-Co3O4 NFs_60, SnO2-Co3O4 NFs_w/o MOF, and SnO2 NFs in the formation cycle (50 mA g–1), (b) Cycle retention tests of SnO2Co3O4 NFs_60, SnO2-Co3O4 NFs_w/o MOF, and SnO2 NFs (500 mA g–1), and (c) Rate capabilities of SnO2-Co3O4 NFs_60, SnO2-Co3O4 NFs_w/o MOF, and SnO2 NFs (expressed in A g–1).


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Figure 6. CV curves of (a) SnO2-Co3O4 NFs_60 and (b) SnO2 NFs at various scan rates (0.1, 0.5, 1.0, and 10.0 mV s–1). The relationship between the anodic peak currents and square root of scan rate for (c) SnO2-Co3O4 NFs_60 and (d) SnO2 NFs.



Figure 7. Ex situ SEM image of (a) SnO2-Co3O4 NFs_60 and its magnified image and (b) SnO2 NFs and its magnified image. Schematic illustration of (c) SnO2-Co3O4 NF_60 and (d) pristine SnO2 NF before and after 300 cycles.


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Feasible Defect Engineering by Employing MOF Templates into One

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