Electrochemical properties of SnO₂ nanoparticles immobilized within a metal–organic framework as an anode material for lithium-ion batteries

Abstract

SnO₂ nanoparticles have been immobilized inside MIL-101(Cr) crystals achieving the novel SnO₂@MIL-101(Cr) anode material with a multiple-core/shell structure. Under the protection of MIL-101(Cr) shell, the pulverization and aggregation of SnO₂ during cycling have been diminished. Both good cyclability and rate capability with a reversible capacity of ∼510 mA h g⁻¹ at 0.1C after 100 cycles have revealed.

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Nowadays lithium-ion batteries (LIBs) have been widely applied in portable devices such as laptops, cell phones and wearable devices.¹ However, their further applications in areas such as electric vehicles and grid energy storage have been restricted by the low specific capacity and energy density of current LIB technologies.^2,3 It is urgent to explore new practicable LIB systems with high specific capacity and energy density. Metal oxides (MOs) have been extensively investigated as anode materials due to their high theoretical capacity, natural abundance and low toxicity.^4,5 Among them, SnO₂ is considered to be one of the most promising anodes because of its relatively high specific capacity (790 mA h g⁻¹, which is more than twice that of the currently used graphite⁶) and low voltage plateau vs. Li⁺/Li.⁷ However, like most conversion-reaction based anode materials, SnO₂ suffers large volumetric variation (>300%) during lithiation/delithiation process.⁸ The volumetric variation will induce pulverization and aggregation of the anode material and result in fast capacity fading.⁹ Thus before practical application of SnO₂ anode, fast capacity fading is a crucial problem to overcome. Towards this, mainly two effective strategies have been pursued. One is to design nanostructural SnO₂ such as nano tubes,¹⁰ nano powder,¹¹ and nano arrays,¹² in which the local empty space can effectively buffer the volumetric variation. Another strategy is to coat electrochemically stable protecting shells on SnO₂, such as amorphous carbon,⁶ graphene¹³ and TiO₂,¹⁴ which can provide physical barrier for the pulverized particles and keep their conductivity. Based on these, it is reasonable to expect improved electrochemical properties to combine both strategies, i.e. small particle size with protecting shells.

Herein, we employed metal–organic frameworks (MOFs) as protecting shell and SnO₂ nano particles as cores achieving a novel SnO₂@MOF multiple-core/shell structure. MOFs are a kind of highly porous crystalline materials constructed by organic linkers and metal ions/cluster nodes.¹⁵ Their physical and chemical properties are highly tunable and they have shown potential applications in variety of areas such as gas storage and separation,^16,17 catalysis,^18–20 detection,²¹ photochemical^15,22 and biomedicine.²³ Recently, our group's work has shown the advancement of MOFs as sulfur host in lithium–sulfur batteries^24,25 and now we have fond their promising applications in SnO₂ based Li-ion battery. As SnO₂ protecting shell, compared with porous carbon, graphene and TiO₂, MOFs feature well-defined pores which can be designed through changing different metal ions and organic ligands. Such well-defined pores of MOFs can serve as templates for the growth of SnO₂ nano particles inside. Moreover, the high porosity will buffer the volume change of SnO₂ particles during cycling as well as increase ion diffusion and electrolyte infiltration. Small particle size, with highly porous shell buffering volume change and confining pulverized particles, such SnO₂ anode is expected to show improved battery performance. However, this promise has never been realized. We report here the first example of SnO₂@MOF composite anode material which has demonstrated both good cyclability and rate capability with a reversible capacity of about 510 mA h g⁻¹ at 0.1C after 100 cycles.

Served as SnO₂ protecting shell in Li-ion battery, firstly, MOFs should be electrochemically stable ensuring permanent porosity to protect and confine pulverized particles. Secondly, MOFs should have large surface area and cage-like pores to provide templates for SnO₂ nano particles. Based on these, MIL-101(Cr) is a good candidate. It is built up by Cr(III) ions and terephthalic acid (BDC) with large BET surface area (4100 m² g⁻¹) and giant cage-like pores (2.9 and 3.4 nm).²⁶ Moreover, it is one of the most stable MOFs and its electrochemical stability has been proved in previous work in which MIL-101(Cr) was employed as sulfur host in Li–S battery.^27,28 The preparation of SnO₂@MIL-101(Cr) has been illustrated in Fig. 1. MIL-101(Cr) crystals were synthesized following the procedure reported elsewhere.²⁹ In order to fully open its pores, MIL-101(Cr) was further treated with NaOH solution to remove “free” BDC ligands before use. Sn⁴⁺ was introduced into MIL-101(Cr) crystals by soaking them into SnCl₄ solutions. Then the Sn⁴⁺ in MIL-101(Cr) was transformed into Sn(OH)₄ by stirring Sn⁴⁺@MIL-101(Cr) in NaOH solution.³⁰ After heat treatment, Sn(OH)₄ in MIL-101(Cr) was finally turned into SnO₂ nano particles.¹³ The phase purity of MIL-101(Cr) was confirmed by PXRD pattern in Fig. 2 which is consistent with the simulated one. No recognisable reflection peak of SnO₂ was observed in SnO₂@MIL-101(Cr), which is probably caused by the poor crystallinity of SnO₂ nano particles inside MOF crystals. Similar phenomenon was also seen in previous reported nano-particle@MOF composites.^24,31 In these studies, the nano-particles cannot be recognized in PXRD patterns due to their poor crystallinity or their tiny particle size. The sharp MIL-101(Cr) reflection peaks in SnO₂@MIL-101(Cr) have demonstrated the crystalline MIL-101(Cr) was intact after encapsulation of SnO₂ nano particles. After encapsulating of SnO₂, the BET surface of SnO₂@MIL-101(Cr) has been reduced to 157 m² g⁻¹ (Fig. S1†) because the pores of MOF have been filled with SnO₂ nano particles. Such moderate BET surface area will help improve ion diffusion efficiency in battery performance. The FTIR spectra (Fig. S2†) clearly show the presence of the vibrational bands characteristic of the –(O–C–O)– groups between 1620 and 1400 cm⁻¹ confirming the presence of the dicarboxylate within SnO₂@MIL-101(Cr).

Fig. 1 Scheme illustration of preparation procedure of SnO₂@MIL-101(Cr).

Fig. 2 PXRD patterns of the synthesized MIL-101(Cr) (black), SnO₂@MIL-101(Cr) (red) and simulated MIL-101(Cr) (blue).

SEM and TEM morphologies of MIL-101(Cr) and SnO₂@MIL-101(Cr) have been displayed in Fig. 3. Fig. 3a shows that the as-synthesized MIL-101(Cr) is octahedron crystals with a diameter of 200–300 nm. Such morphologies have been maintained after SnO₂ encapsulation as indicated by SEM of SnO₂@MIL-101(Cr) in Fig. 3b. However, after heat treatment, the surface of octahedron has become concave due to the loss of solvent in pore. The surface of SnO₂@MIL-101(Cr) crystals are clear and smooth demonstrating no SnO₂ exists on the surface of MOF shell. SnO₂ “clouds” can be recognized from SnO₂@MIL-101(Cr) TEM morphologies in Fig. 3c and d, which heterogeneously distributed in MOF crystals. The SnO₂ “clouds” with diameter less than 50 nm are built up by many smaller SnO₂ nano particles which have different contrast with “blank” region in MOF crystals. However, as shown in Fig. 4, consistent distribution of Sn, C, O and Cr demonstrates there is also SnO₂ distributed in the “blank” region in MIL-101(Cr). So based on the results of TEM morphologies and EDS elemental mappings, the distribution of SnO₂ in MOF shell has been confirmed: some of the SnO₂ nano particles gathered together into “clouds” while the rest distributed in the pores of MOF shell which could hardly be distinguished in TEM morphologies due to the small particle size. The SnO₂ loading amount in SnO₂@MIL-101(Cr) has been precisely determined by ICP as 21.3 wt%.

Fig. 3 SEM morphologies of (a) MIL-101(Cr) and (b) SnO₂@MIL-101(Cr) and TEM morphologies of (c and d) SnO₂@MIL-101(Cr) with magnification of the selected region in the inset.

Fig. 4 Dark filed TEM morphologies of SnO₂@MIL-101(Cr) and EDS elemental mappings of the selected region.

In order to study the influence of MOF protecting shell on SnO₂ battery performance, a series of electrochemical tests has been performed on SnO₂@MIL-101(Cr). Fig. 5a displays the first 3 and the 50^th galvanostatic charge/discharge profiles of SnO₂@MIL-101(Cr) at 0.1C (1C = 790 mA g⁻¹) with cut-off voltage between 0.02–2.5 V. It shows typical SnO₂ reaction mechanism which can be concluded as follows:³²

4Li⁺ + SnO₂ + 4e⁻ → Sn + 2Li₂O (1)

xLi⁺ + Sn + xe⁻ ↔ Li_xSn (2)

Fig. 5 (a) Galvanostatic charge/discharge profiles at 0.1C, (b) cyclic voltammograms at 0.1 mV s⁻¹, (c) cycle performance with coulombic efficiency at 0.1C and (d) rate capability of SnO₂@MIL-101(Cr) at different charge/discharge rate.

Due to the irreversible reaction from SnO₂ to Sn (eqn (1)) large capacity loss was observed after the first cycle. During the following cycles, discharge plateau below 0.5 V and charge plateau at about 1 V can be assigned to the reversible reaction of eqn (2). Such reaction mechanism has been further elucidated by its CV curves in Fig. 5b. Two oxidation peaks appear around 0.51 V and 1.25 V, respectively. The 0.51 V oxidation peak corresponds to the de-alloying of Li_xSn (eqn (2)), while the weak oxidation at 1.25 V which gradually disappears in the following cycles, is caused by the partly reversible reaction of the eqn (1).^32,33 The large anode peak located near 0.02 V is attributed to the alloying reaction of eqn (2). The small one located at about 1 V is probably caused by the formation of the solid electrolyte interphase (SEI) layer.³² EIS spectra of the half cell with the SnO₂@MIL-101(Cr) anode after the 1^st and 100^th cycle have been collected and compared, as shown in Fig. S3.† The intercept at the real axis Z′ of the EIS spectra corresponds to the contact resistance (R_e), and the diameter of the semicircle in high frequency represents the charge-transfer resistance (R_ct).³⁴ The R_e (∼65 Ω) doesn't change significantly while R_ct has increased from 210 Ω to 320 Ω after 100 cycles. Such phenomenon is speculated caused by the loss of electronic path due to the pulverization of SnO₂ during cycling. Featuring small particle size with porous shell buffering volume change and confining pulverized particles, SnO₂@MIL-101(Cr) has demonstrated both cyclability and rate capability. As shown in Fig. 5c, despite of the large capacity loss after the first cycle which is caused by the irreversible reaction of eqn (1), it reveals a reversible capacity of about 510 mA h g⁻¹ after 100 cycles at 0.1C. Rate performance in Fig. 5d also indicates that anode material of SnO₂@MIL-101(Cr) features both good rate capability and capacity recoverability. In order to investigate the influence of MOF shell on battery performance, as comparison, electrochemical performance of bare SnO₂ particles has been tested as displayed in Fig. S4.† It suffered rapid capacity fading and only 200 mA h g⁻¹ was left after 20 cycles. Without the porous structure of MOF as template, bare SnO₂ grew into large particles immediately after it was synthesized. Moreover, during cycling, without the protection of MIL-101(Cr) shell, the pulverized particles aggregated causing poor cyclability. The superior properties of SnO₂@MIL-101(Cr) have demonstrated the significant improvement of battery performance benefited from MOF protecting shell. Although the SnO₂ content in electrode is still too low (17 wt%) for practical battery applications, our work has demonstrated a novel approach to the protection of MOs anodes using porous MOFs shell. Our study about increasing MOs loading amount will be carried out soon.

Conclusions

In summary, we have proposed a novel strategy using porous MOFs as protecting shell and templates to synthesize SnO₂@MIL-101(Cr) multiple-core/shell anode material. The SnO₂ growing in the pores of MIL-101(Cr) features very small particle size. Served as protecting shell, the electrochemically stable MIL-101(Cr) buffers the volumetric variation of SnO₂ and confines pulverized particles. Such chemical and structural advancement has resulted in good battery performance. Our work has developed new application of MOFs in Li-ion batteries.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (No. 51272229, 51272231 and 51472217), Zhejiang Provincial Natural Science Foundation of China (No. LR13E020001 and LZ15E020001), Qianjiang Talent Project (No. QJD1302009) and Fundamental Research Funds for the Central Universities (No. 2015QNA4009, 2015FZA4008 and 2014XZZX005).

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Footnote

† Electronic supplementary information (ESI) available: Details of synthesis and characterization and additional figures. See DOI: 10.1039/c5ra16587c

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