Low Interface Energies Tune the Electrochemical Reversibility of Tin

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Low Interface Energies Tune the Electrochemical Reversibility of Tin Oxide Composite Nanoframes as Lithium Ion Battery Anodes Lei Zhang, Jun Pu, Yihui Jiang, Zihan Shen, Jiachen Li, Jinyun Liu, Haixia Ma, Junjie Niu, and Huigang Zhang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b11062 • Publication Date (Web): 08 Oct 2018 Downloaded from http://pubs.acs.org on October 9, 2018

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ACS Applied Materials & Interfaces

Low Interface Energies Tune the Electrochemical Reversibility of Tin Oxide Composite Nanoframes as Lithium Ion Battery Anodes Lei Zhanga, Jun Pua, Yihui Jiangb, Zihan Shena, Jiachen Lic, Jinyun Liud, Haixia Mac, Junjie Niue and Huigang Zhanga,f,* a

National Laboratory of Solid State Microstructures, College of Engineering and Applied

Sciences, Collaborative Innovation Center of Advanced Microstructures, Institute of Materials Engineering, Nanjing University, Jiangsu, China b

c

College of Chemical Engineering, Northwest University, Xi’an 710069, China d

e

Engineering Department, Smith College, Northampton, MA, 01063, USA

Department of Materials Science, Anhui Normal University, Wuhu, China

Department of Materials Science and Engineering, University of Wisconsin-Milwaukee, Milwaukee, WI 53211, USA

f

Key Laboratory of Advanced Energy Materials Chemistry (Ministry of Education), Nankai University, Tianjin, China

Abstract: The conversion reaction of lithia can push up the capacity limit of tin oxide based anodes. However, the poor reversibility limits the practical applications of lithia in lithium ion batteries. The latest reports indicate that the reversibility of lithia has been appropriately promoted by compositing tin oxide with transition metals. The underlying mechanism is not revealed. To

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design better anodes, we studied the nanostructured metal/Li2O interfaces through atomic scale modeling and proposed a porous nanoframe structure of Mn/Sn binary oxides. The first principle calculation implied that because of a low interface energy of metal/Li2O, Mn forms smaller particles in lithia than Sn. Ultrafine Mn nanoparticles surround Sn and suppress the coarsening of Sn particles. Such a composite design and the resultant interfaces significantly enhance the reversible Li-ion storage capabilities of tin oxides. The synthesized nanoframes of manganese tin oxides exhibit an initial capacity of 1620.6 mA h g−1 at 0.05 A g−1. Even after 1000 cycles, the nanoframes anode could deliver a capacity of 547.3 mA h g−1 at 2 A g−1. In general, we demonstrated a strategy of nanostructuring interfaces with low interface energy to enhance the Li-ion storage capability of binary tin oxides and revealed the mechanism of properties enhancement, which might be applied to analyze other tin oxide composites.

Keywords: Manganese tin oxides; Nanoframes; Li-ion batteries; Anodes; First principle calculation

1. Introduction The rapid developments of electric vehicles and portable electronics stimulate the growing demands for high energy-density lithium ion (Li-ion) batteries.1 High capacity materials hold promises to increase the specific energy of Li-ion batteries. Among many previously-reported anodes,2−6 tin oxides have received intensive attentions because of large theoretical capacity and fast charge/discharge capability. The in-situ studies showed that a high capacity can be obtained via conversion reaction of tin oxides and alloying reaction of metallic Sn.7−11 The ACS Paragon Plus Environment

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formation of Li4.4Sn alloy can deliver a capacity of 994 mA h g−1.10 Mössbauer spectroscopy of 119Sn

revealed that the lithiation process of SnO2 first involves a minor intercalation reaction of

LixSnO2 and subsequent formation of Sn(II) and Sn(0) species via the conversion reactions.12−14 The utilization of conversion reactions of lithia for lithium storage is able to further push the capacity limit of SnO2.6,15 However, the poor reversibility of lithia largely limits the practical applications of SnO2 as compared to other intercalation anode materials.16 Due to the irreversibility of lithia, tin oxides usually exhibit rapid capacity decay in a few cycles.3 Tarascon et al. reported that low-valence transition metal oxides (MO, M = Co, Ni, Cu, or Fe) can be used as anodes via conversion reactions.17 A series of composite oxides of transition metals (M) and tin have been reported to promote the conversion reaction of lithia, for example, Co2SnO4, SnO2/Co, NiO/SnO2, and Zn2SnO4.18−21 Their experimental results demonstrat that binary or ternary M/Sn oxides are able to increase either capacity or cycling properties as compared to pristine SnO2. The performance improvement had been attributed to synergic or catalytic effect.22−26 Recently, Hu et al. find that the critical approach to a reversible conversion of lithia is to create a large interface area between lithia and metallic Sn because the large interface area facilitates the atomic interdiffusion.12 On the other hand, Sn is inclined to coarsen upon long-term cycling,27−29 which leads to the formation of crack and electrical isolation of active Sn particles by solid electrolyte interphase (SEI),30,31 as illustrated in Figure 1a. A large Sn/lithia interface area eventually reduces the cyclability of anodes if the interface stability could not be maintained upon cycling. In order to reversibly convert lithia, the nano-sized morphology of Sn particles must be retained in the lithia matrix. The dopant M in a



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tin oxide composite is believed to stabilize the interface or confine Sn grains, as shown in Figure 1b.15 Using in-situ transmission electron microscopy (TEM) technology, Jiang et al. observed that the composite of Sn/M oxides prevent the migration and aggregation of nanoparticles during cycling.32,33 These results indicate that a composite structure can not only enhance the reversibility of lithia but also improve the cycling property of tin oxides. However, the underlying mechanism of the enhancements needs to be further clarified. The size and volume of SnO2 particles play an important role in the reversibility of the conversion step (SnO2 + 4Li+ + 4e ↔ Sn + 2Li2O). A critical size of about 3 nm enables the reversible cycling of SnO2. However, the large surface area of nanoparticles usually results in the excessive SEI formation.11,34,35 Small particle sizes also lead to low packing density. It is desirable to generate a hierarchical structure in which primary particles below the critical size assemble into a secondary structure for the purposes of high packing density. Various porous structures with nano-sized SnO2 have been synthesized to simultaneously take advantage of the high activity of small particles and the high density of large hierarchical particles.36−38 Porous structures also help to alleviate the strains caused by the volume change (about 250%) which is detrimental to the anode cyclability.18,31,39 Thus, to build a better tin oxide based anode, we have to understand the lithia reversibility fundamentally on multiple scales from atoms, interfaces, to particles and apply the underlying mechanisms to design porous nanocomposites with hierarchical structures. Herein, we first focused on the metal/Li2O interfaces and studied how the particle sizes of metallic Mn and Sn change in the composite during lithiation (Figure 1c). Mn is inclined to form


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ultrafine nanoparticles because of low interface energy as compared to Sn in a lithia matrix. Ultrafine Mn can separate Sn nuclei and suppress their aggregation (Figure 1b). To realize such a design strategy, we synthesized porous manganese tin oxide nanoframes (PMNs). Benefiting from its composite structure and porous nanoframe morphology, the PMN anode exhibits an initial capacity of 1620.6 mA h g−1 at 0.05 A g−1 and a high reversible capacity of 912.0 mA h g−1 at 0.2 A g−1 after 100 cycles. The long-term cycling results show that PMN could deliver a capacity of 547.3 mA h g−1 at 2 A g−1 even after 1000 cycles.

Figure 1. Schematic illustrations of (a) coarsening and cracking of large Sn particles and (b) suppressing Sn growing with small Mn particles surrounding. (c) Nucleation and interface models of Sn and Mn particles in Li2O during lithiation. Interface models of (d) Li2O(100)/Sn(010) (Li16O10Sn6, Li16O10Sn10, and Li16O10Sn14) and (e) Li2O(111)/Mn(111) (Li8O5Mn6, Li8O5Mn10, and Li8O5Mn14); (f) Illustrated fabrication and lithiation of PMN. 2. Experimental section



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2.1. Computation methodology and models The critical nucleation radius of metals in a lithia matrix was estimated by the classic nucleation thermodynamics according the reference (see the supporting information (SI) for details).40 The interface energy of Li2O and metals (M=Mn or Sn) was calculated by the density functional theory (DFT). For a given interface of Li2O/M, the interface energy was calculated based on the following expression:41,42 𝛾𝐿𝑖2𝑂/𝑀 =

1 [𝐸 ― 𝑛𝑂𝐸𝐿𝑖2𝑂 ― 𝑛𝑀𝐸𝑀 + (2𝑛𝑂 ― 𝑛𝐿𝑖)𝐸𝐿𝑖 + (2𝑛𝑂 ― 𝑛𝐿𝑖)Δ𝜇𝐿𝑖] 2𝐴 𝐿𝑖2𝑂/𝑀 (1)

Where 𝛾𝐿𝑖2𝑂/𝑀 is the interface energy of Li2O/M, 𝐸𝐿𝑖2𝑂/𝑀 is the total energy of interface model, 𝐸𝐿𝑖2𝑂 and 𝐸𝑀 are the energies of bulk Li2O and M, 𝑛𝐿𝑖2𝑂 and 𝑛𝑀 are the number of Li2O and M in the interfae model, 𝑛𝑂 and 𝑛𝐿𝑖 are the numbers of oxygen and lithium atoms respectively, Δ𝜇𝐿𝑖 is translated to the negative of the voltage against Li/Li+, A is the interface area. The strain energy may exist inside metal slabs for commensurating two phases. It was proportional to the atom number of M (𝑛𝑀) in metal slabs, the interface energy could be further estimated by fitting the following expression. 𝛾𝐿𝑖2𝑂/𝑀 = 𝛾𝑙𝑖𝑚 𝐿𝑖2𝑂/𝑀 + 𝑛𝑀𝜎

(2)

where 𝛾𝑙𝑖𝑚 𝐿𝑖2𝑂/𝑀 is the fitted limit of the interface energy and σ is a constant related to the strain energy in metal slabs. The model building and calculation strategies could be referred to literatures.43−45

2.2 Synthesis of PMNs, bulk MnSnO3, SnO2, and MnO In a typical synthetic route of PMNs, about 0.01 mol MnCl2·4H2O and 0.01 mol SnCl4 were dissolved in 100 mL deionized water under vigorous stirring until the solution turned to be transparent. About 80 mL NaOH (1 M) was slowly dropped into the solution to form brown precipitates and stirred for 1 h. The precipitates were separated by centrifugation and washed ACS Paragon Plus Environment

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with deionized water. After dried in air at 50 °C for 12 h, the obtained powder was annealed in Ar gas at 400 °C for 4h with a ramp rate of 4 °C min−1. Bulk MnSnO3 powders were prepared by hydrothermal treatment of the aqueous solution of 0.225 M NaOH, 0.015 M SnCl4 and 0.015 M MnCl2·4H2O at 120 °C for 2 h. The PMNs powders were dispersed in 30 mL deionized water containing 6 mL HNO3 and then transferred into a 50 mL Teflon-lined stainless-steel autoclave. After heated at 150 °C for 4 h, the resultant SnO2 nanoparticles were centrifuged, washed, and dried at 60 °C. MnO nanoparticles were synthesized by mixing 0.1 M MnCl2 ·4H2O and 1 M NaOH directly. After centrifuged and washed, the precipitates were annealed in form gas (95% Ar and 5% H2) at 400 °C for 4h.

2.3 Materials characterization X-ray diffraction (XRD) patterns were recorded on a Rigaku D/MAX 2500V with Cu Kα radiation. X-ray photoelectron spectra (XPS) spectra were conducted on an ESCALab MKII Xray photoelectron spectrometer with Mg Kα X-ray as the excitation source. A Zeiss Ultra 55 field-emission scanning electron microscope (SEM) was used for morphology observation and energy dispersive X-ray spectroscopy (EDX) measurement. High-resolution TEM (HRTEM) images were recorded on the FEI Tecnai G2 20 microscope at 200 kV. N2 adsorption isotherms were measured on a Micromeritics ASAP 2020M+C instrument.

2.4 Electrochemical Measurement The slurries of 70% active material, 20% acetylene black, and 10% polyvinylidene fluoride (PVDF) in N-methyl pyrrolidone (NMP) was casted on a copper foil to form an anode laminate. The mass loading of active materials was controlled at ~2 mg cm−2. The obtained samples were



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assembled into coin cells (CR2032) with lithium foil as counter electrode in an Ar-filled glove box. The electrolyte was 1 M LiPF6 in a 1:1 (volume ratio) mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC). Cyclic voltammogram (CV) curves were measured in the range of 0.01−3 V at varied scan rates. Electrochemical impedance spectroscopy (EIS) were obtained between 100 kHz to 100 mHz with a potentiostat (BioLogic VSP, France) at 3 V. Galvanostatic charge/discharge curves were measured with a CT2001A test system (Land Electronic Co., China) between 0.01 and 3 V.

3. Results and discussion

According to the classical nucleation theory, the critical nucleation size is proportional to the interface energy of Li2O/metals (see SI). To estimate the interface energy of Li2O/Sn and Li2O/Mn using the DFT calculation, we build the interface models by sandwiching a metallic slab between Li2O slabs to satisfy the periodic boundary conditions as shown in Figure. 1d and e. To commensurate the lattice parameters of metal and oxide, metal slabs were appropriately strained (see details in the SI). The induced strain energy was proportional to the number of metal atoms. Three models were built for each interface with varying the thickness of metal slabs. Figure 2a and d show the linear regression of interface energy with respect to the number of Sn or Mn atoms. From the intercepts, one can deduce that the interface energy of Li2O/Sn with the most stable configuration was estimated to 4.1 J m−2 whereas the Li2O/Mn model has the interface energy of 3.2 J m−2. By using the Equation SI-2, the critical nucleation size for Mn in a lithia matrix was calculated to be 0.18 nm, which is much lower than that of Sn (0.40 nm).


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Figure 2. (a) Calculated interface energies of Li2O/Sn with varied Sn atoms, (b) PDOS of the interface atoms and middle atoms which reside in the middle of Li2O or Sn slabs of Li16O10Sn10, (c) electron difference and charge densities of Li16O10Sn10. (d) Calculated interface energies of Li2O/Mn with varied Mn atoms, (e) PDOS of the interface atoms and middle atoms which reside in the middle of Li2O or Mn slabs of Li8O5Mn10, (f) electron difference and charge densities of Li8O5Mn10. Figure 2b and e present the partial density of states (PDOS) of the middle and interface atoms in Li16O10Sn10 and Li8O5Mn10, respectively. With the O2p orbital in the middle of Li2O slabs as the references, one can see that the interface O2p states in Li16O10Sn10 shift towards the Fermi energy and however, the energy level of interface O2p is lowered in Li8O5Mn10. Figure 2c and f show that electrons are accumulated along M-O bonds and most at the oxygen atoms, indicating the strong ionicity of M-O bonds. Thus, the downshift of the interface O2p orbital in ACS Paragon Plus Environment


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Li8O5Mn10 is induced by the relatively strong electrostatic potentials, indicating the strong Li2O/metal interaction at interfaces as compared to that in Li16O10Sn10. The difference on interface energies can also be inferred from the bond energy. As shown in Table S1, The MnO bond energy (2.46 eV) in MnO is stronger than those of Sn-O bonds in SnO (1.43 eV) and SnO2 (1.54 eV). Such a strong affinity of Mn to O implies a low Li2O/Mn interface energy as compared to that of Li2O/Sn. Table 1 lists the calculated interface properties of Li2O/Sn and Li2O/Mn. The work of separation (Wsep) is the work required to reversibly separate the interface into two free surfaces and thus a measure of the interface bond strength.46,47 The calculated

Wsep of Li2O/Mn is 6.7 J m−2, which is larger than that of Li2O/Sn (4.9 J m−2), indicating the strong bond strength between Li2O and Mn slabs. It can be inferred from the above calculation and analyses that the strong affinity between Li2O and Mn leads to the small nucleation size. Table 1. Calculated interface properties of Li2O and Mn/Sn Interface Work of separation (J m−2) Interface energy (J m−2) Critical nucleation size (nm)

Li2O/Mn

Li2O/Sn

6.7

4.9

3.2

4.1

0.18

0.40


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Based on the calculation results, we design a composite material of manganese tin oxides. Figure S4a shows the TEM image of the uniform precursor MnSn(OH)6 nanocubes. These nanocubes exhibit a uniform rectangular morphology with the side length of 20−25 nm. A lattice fringe spacing of 0.392 nm confirms the (200) crystallographic plane of MnSn(OH)6. The selected area electron diffraction (SAED) of one nanocube in Figure S4c confirms that these nanocubes are single crystals of MnSn(OH)6 (JCPDS card No. 72-0007). After heat treatment, the precursors were dehydrated to form porous MnSnO3 nanoframes (Figure 3a). The phase composition was confirmed by the XRD patterns in Figure 3g. The reduced intensity and increased peak width indicate the low crystallinity of dehydrated PMNs as compared to MnSn(OH)6. Figure 3b presents the HRTEM image of a PMN. It shows that the nanoframe consists of crystalline grains and porous interiors. The multiple rings in the SARD pattern (Figure 3c) were indexed to MnSnO3 (JCPDS card no.85-0615). It confirms that PMNs are polycrystalline. Figure 3d−f present the TEM images of the control samples of bulk MnSnO3, SnO2 and MnO, respectively. The XRD patterns are given in Figure S5. The XPS analysis (Figure 3h) shows that the binding energies of Mn 2p3/2 and Mn 2p1/2 are around 641.9 eV and 653.6 eV, respectively, indicating the existence of divalent Mn.48 The XPS signals of Sn 3d5/2 (486.7 eV) and Sn 3d3/2 (495.1 eV) (Figure 3i) agree with Sn4+ in MnSnO3.20 Figure 3j shows that the BET surface area of the PMNs is 69.1 m2 g−1 and the main pore size is around 4 nm.



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Figure 3. (a) TEM, (b) HRTEM, and (c) SAED images of PMNs. (d–f) TEM images of (d) bulk MnSnO3, (e) SnO2 and (f) MnO. (g) XRD patterns of PMNs. XPS spectra of PMNs: (h) Mn 2p and (i) Sn 3d signals. (j) BET analysis of PMNs (the inset shows the pore size distribution). To test whether the porous structure and nanosized grains of PMNs are favorable for high capacity Li-ion battery anodes, we assembled them with lithium into half cells, tested their electrochemical properties, and compared them with bulk MnSnO3 and SnO2 anodes. Figure 4a shows the CV curves of PMNs. In the first cycle, PMNs exhibit an irreversible peak at 0.95 V which is attributed to the formation of SEI.49 Three reduction peaks at 1.08, 0.5, and 0.06 V in the second cathodic scan are correlated to the multiple-step reduction reactions of metallic Sn/Mn and sequential Sn alloying process upon lithiation.50,51 Two main peaks around 0.55 and ACS Paragon Plus Environment

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1.17 V during the oxidation scan are ascribed to the de-alloying reactions of LixSn and the formation of MnO, respectively.52,53 A broad hump at 1.25−2.0 V implies the re-oxidation of Sn via the reversible reaction of lithia in PMNs. The galvanostatic charge/discharge profiles in Figure 4b show a high initial capacity of 1620.6 mA h g−1. It decreases to 1247.0 mA h g−1 in the second cycle. The overall lithiation profile is a slope without a distinct plateau around 1.0 V that large SnO2 particles usually show.54 The capacity below 1.25 V is around 1164.8 mA h g−1. Upon delithiation, there is about 567.8 mA h g−1 capacity above 1.25 V, indicating the possible formation of tin oxides. The XPS spectra in Figure 4c shows that the Sn 3d signals of a lithiated sample could be deconvoluted to three components which are attributed to Sn(IV/II) oxides and metallic Sn, respectively.9,55,56 The oxides signals may result from the superficial oxidation. After delithiation, there are two typical spin-orbit splitting peaks of Sn 3d5/2 (486.7 eV) and 3d3/2 (495.1 eV), indicating that tin is re-oxidized to Sn(IV). It explains the extra capacity above 1.25 V.



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Figure 4. (a) CV curves and (b) galvanostatic charge/discharge curves of PMNs. (c) XPS spectra of PMNs after lithiation and delithiation. (d) CV curves of bulk MnSnO3. SAED pattern of delithiated (e) PMNs and (f) SnO2. (g) CV curves of nanosized SnO2. Illustrated cycling models of (h) PMNs and (i) SnO2. To understand how the particle size influences reactions above 1.25 V, we intentionally synthesized bulk MnSnO3 (~200 nm) and SnO2 (~5 nm) and measured their CV curves under the same conditions as used for PMNs. Bulk MnSnO3 has a nanocube morphology (Figure 3d) and SnO2 nanoparticles are generally aggregated (Figure 3e). It could be clearly seen from Figure 4d that bulk MnSnO3 delivers a minor current density at 1.25−2.0 V as compared to PMNs and no significant peaks appear above 1.25 V. It implies that only a small portion of bulk


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MnSnO3 is cyclable above 1.25 V probably because the existence of Mn helps to enhance the conversion of tin back to its oxide. The enhancement is not as significant as PMNs because the large particle size (Figure 3d) of bulk MnSnO3 may lead to the incomplete conversion. It is interesting to note that pure SnO2 shows an oxidative peak with a tail at 1.35 V in the first scan (Figure 4g). In contrast to bulk MnSnO3, the small size of SnO2 (5 nm) enhances the delithiation step above 1.25 V and exhibits the peaks in the first oxidative scan. Thus, the lithiated SnO2 is able to be partially delithiated back to its oxide states. However, the two oxidative peaks of SnO2 above 1.25 V disappear in the following cycles and the current density decreases to even lower values than bulk MnSnO3. It indicates that the reversibility of SnO2 above 1.25 V degrades with cycles. By comparing the CV curves and morphologies of three samples, it may be inferred that because of large particle sizes, the long distance for atoms migration limits the delithiation and reformation of Sn/Mn oxides, leading to poor cycling reversibility above 1.25 V. Because PMNs have the advantages of small grain sizes as SnO2 does and the enhancement effect of Mn/Sn binary oxides, PMNs are able to demonstrate good reversibility to oxides and gain extra capacity. Figure 4e presents the SAED pattern and rotational integral of a delithiated PMN, which are indexed to the composite of SnO2 and MnO, instead of MnSnO3. The HRTEM image in Figure S6 shows the phase separation of SnO2 and MnO, which is not re-oxidized back to PMNs. It should be noted that the composite of SnO2 and MnO differs from a simple mechanical mixture because the composite is formed by the electrochemical reaction, yielding a good dispersion of nanosized SnO2 and MnO grains, which is of vital importance to the reversibility of composite



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anodes. The SAED pattern of the delithiated SnO2 anode in Figure 4f shows that the main rings are related to metallic Sn. In conjunction with the CV analysis, it appears that there are two distinct cycling models for PMN and SnO2, respectively. The PMNs anode forms a composite with nanosized SnO2 and MnO grains, which could be cycled via both alloying and conversion reactions as illustrated in Figure 4h. However, the SnO2 anode is mainly charged/discharged via alloying reactions since the second cycle as illustrated in Figure 4i. The difference on cycling models may result from the Mn doping and nanoframe structure, which will be further studied later. Figure 5a shows the rate properties of PMNs, bulk MnSnO3, and SnO2 at varied current densities. The PMNs anode delivers higher capacities at both low and high current densities than bulk MnSnO3 because of the smaller grain size. SnO2 shows higher specific capacities at low current densities than bulk MnSnO3 because of its relatively high theoretic capacities. The decreased capacity retention of SnO2 at high current densities may be because the capacity delivered by the conversion reaction decays significantly with cycling as indicated by the CVs in Figure 4g. The EIS analysis in Figure 5b may also explain the difference of rate capabilities. The semicircle diameter in the Nyquist plots of Figure 5b is related to the resistance of interface charge transfer. A small semicircle diameter of PMNs indicates that the nanosized grain enables the low resistance and facilitates the interface charge transfer, leading to the enhanced rate properties as compared to bulk MnSnO3. Figure 5c presents the cycling properties of three samples at a low current density of 0.2 A g−1. All three samples suffer a capacity decay in the first few cycles. After 100 cycles, PMNs deliver a relatively stable capacity of 912.0 mA h g−1


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while bulk MnSnO3 and SnO2 only retain 383.2 and 223.1 mA h g−1, respectively. To differentiate the conversion and alloying reactions of tin oxides, we separated the total capacity at 1.25 V. Figure 5d presents the separate capacities contributed by conversion and alloy reactions. The PMNs sample shows higher capacities for either conversion or alloying reactions than bulk MnSnO3 and SnO2 because PMNs have the small grain size and porous frame structure. The capacity contributed by the conversion reaction of tin oxides gradually decrease on all the three samples during cycling. Figure 5e shows the long-term cycling properties at 2 A g−1. The PMNs anode shows an initial capacity of 1173.0 mA h g−1 with an initial Coulombic efficiency (CE) of 65%. After 35 cycles, the capacity decreases to 716.0 mA h g−1 and the CE increases to 99.7%. Although bulk MnSnO3 and SnO2 show comparable capacities initially, they degrade rapidly with the cycle numbers. After 1000 cycles, PMNs are still able to deliver a capacity of 547.3 mA h g−1. Figure S7 presents the TEM images and EDX elemental mapping images of pristine and cycled PMN. The uniform distribution of Mn and Sn indicates that there is no segregation of Sn or Mn, which may explain the good cyclability of PMNs. Figure S8 shows the electrochemical properties of MnSn(OH)6. Although the precursor as an anode can provide a high initial capacity, its CEs are lower than PMNs. By comparing the electrochemical properties with bulk MnSnO3, SnO2, and MnSn(OH)6, we can conclude that PMNs are able to increase the cyclability of tin oxides and the mechanism may result from the unique structure.



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Figure 5. Electrochemical properties of PMNs, bulk MnSnO3, and SnO2: (a) rate properties, (b) Nyquist plots, (c) cycling performances at 0.2 A g−1, (d) capacity variation at two voltage ranges with respect to the cycle number, (e) long-term cycling performance at 2 A g−1. The nanoframe structure is able to provide the short diffusion length for Li-ions in both solid and electrolyte phases. A short Li-ion diffusion length will affect the reaction mechanism and kinetics of PMNs. Figure 6a shows the CV curves of PMNs at varied scan rates from 0.1 to 0.6 mV s−1. The current response (i) obeys a power law relationship with respect to the scan rate (v): i(v) = avb.57−59 Both a and b could be determined by the slope and intercept from Figure 6b.


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In the range of 0.1 to 0.4 mV s−1, the b-values for cathodic and anodic peaks are 0.74 and 0.93, respectively, suggesting that the kinetics is mainly surface-controlled. When the scan rate increases over 0.4 mV s−1, the b-values of cathodic and anodic peaks decrease to 0.58 and 0.66, respectively, indicating that the kinetics is diffusion-controlled. The capacity contributions from capacitive and diffusion-controlled charge could be quantified on the basis of the relationship: i(v) = k1v + k2v1/2, where k1v and k2v1/2 correspond to the current contributions from the surface capacitive effects and the diffusion-controlled intercalation process, respectively. The shaded region in Figure 6c shows the capacitive contribution at 0.3 mV s−1. A close examination of capacity contributions at varied scan rates (Figure 6d) allows us to conclude that with the increase of scan rates, the percentage of capacitive contribution increases and the lithiation/delithiation reactions gradually turn to be diffusion-controlled. A large portion of capacitive contribution indicates that nanoframes provide rapid kinetics for Li-ion storage.



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Figure 6. (a) CV curves of PMNs at varied scan rates, (b) log(i) versus log(v) profile, (c) quantitative analysis of capacitive contribution in a CV scan at 0.3 mV s−1, (d) contribution ratio of capacitive and diffusion-controlled charge at varied scan rates. To further elucidate why PMNs improved the cycling properties, we conducted the ex-situ TEM analysis on three partially lithiated anodes of PMNs, SnO2, and MnO, respectively. Figure 7a shows the TEM image of cycled PMNs particles. The diffraction rings in Figure 7g could be indexed to metallic Sn and Mn. The diffusive background is attributed to the amorphous Li2O matrix. The lattice fringes in the HRTEM image of Figure 7d could be identified as the crystallized Sn and Mn as circled by the dashed line. Figure 7b and e show that Sn particles in the lithiated SnO2 anode have a spherical shape with relatively large diameters as compared to the Sn particles in Figure 7d. The TEM images of the MnO anode in Figure 7c and f indicate


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that metallic Mn particles are much smaller than Sn in Figure 7b. The diffraction patterns and rotational integration plots in Figure 7h and i agree with the conversion mechanism in the reactions of SnO2 + 4Li+ + 4e ↔ Sn + 2Li2O and MnO + 2Li+ + 2e ↔ Mn + Li2O.

Figure 7. Ex-situ (a−c) TEM and (d−f) HRTEM images and (g−i) SAED patterns of (a, d, and g) PMNs, (b, e, and h) SnO2, and (c, f, and i) MnO

The differences of metal (Sn and Mn) particle size in Figure 7d−f seem to imply the important mechanism of metals nucleation and growth in a lithia matrix upon the electrochemical cycling. Figure 8 shows the statistical results of the particles size of Sn and Mn upon the HRTEM images. It indicates that Sn usually exhibits large particles with the average diameter of about 25 nm in the partially lithiated SnO2 anode. By contrast, Mn particles show the average diameter ACS Paragon Plus Environment


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around 5 nm in the MnO anode. Grain coarsening should be avoided to build better Sn-based anodes because the large particles lead to poor cycling properties. It is interesting to note that in the lithiated PMNs anode, Mn shows the same particle size as it does in a MnO anode and Sn particles exhibit a reduced average particle size. It indicates that Mn may suppress the Sn particles to grow as large as Sn does in the pure SnO2 anode. The experimental results basically corroborate our first-principle calculation. The thermodynamic mechanism of the good cyclability of PMNs anodes is attributed to the relatively small nucleation size of Mn as compared to Sn and the suppression of Sn growth by ultrafine Mn.

Figure 8. Statistics of the average size of Sn and Mn nanoparticles in cycled PMNs, SnO2, and MnO anodes.

After nucleation, the growth of metal nuclei depends on the mobility of metal atoms at the interface and lithia matrix. Sn has an extremely low recrystallization temperature because of the large atom radius and low melting point (~506 K).9,27,60 Mn migrates slowly in lithia due to the high melting points (1519 K) and the strong M-O affinity. Thus, the nano-sized Sn prefers growing large particles at room temperatures as compared to Mn. Dahn et al. reported that the


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liberated Sn atoms could react with Mn to form Sn2Mn in a Sn-Mn-C anode.61 A reformation of Sn-Mn alloy was not observed in our PMNs anode. The strong Mn-O affinity on the interfaces may impede the Sn-Mn alloying process in lithia. Thus, Mn and Sn nanoparticles precipitated out of the lithia matrix as their own metallic phases. The ultrafine Mn nanoparticles were welldispersed around Sn and limited the coarsening of Sn particles. This synergic effect may explain why the average size of Sn particles in cycled PMNs is smaller than that in cycled SnO2 (Figure 8). In general, the nanoframe and composite structures are able to modulate the nucleation and growth of metal nanoparticles to achieve a homogeneous distribution of re-sized Sn particles, which in turn improves the overall conversion reaction and obtains a high capacity as well as a better cycling capability.q

4. Conclusions Through the first principle calculations of the Li2O/M (M = Sn or Mn) interfaces, we found that because of the low interface energy, Mn in PMNs shows a small critical nucleation size during lithiation. Ultrafine Mn nanoparticles could separate Sn nuclei and suppress the aggregation or growth of Sn particle, leading to ideal homogenous mixtures of metallic Sn/Mn embedded in a lithia matrix. According to the calculated mechanism, we proposed a porous nanoframe structure of Mn/Sn binary oxide via a simple co-precipitation and heat treatment route. Because of Mn/Sn synergic effects and nano-sized grains, the nanoframes are able to activate the conversion reaction of lithia while the porous structures accommodate the volume change. The lithia reversibility and electrode cyclability were significantly improved. The PMNs anode exhibits an initial capacity of 1620.6 mA h g−1 at 0.05 A g−1 and could even deliver a capacity ACS Paragon Plus Environment


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of 547.3 mA h g−1 even after 1000 cycles at 2 A g−1. The results of this work provide a highperformance anode material and help the in-depth understanding of properties enhancement, which may be applied to other composite oxides systems.

Notes The authors declare no competing financial interest. Acknowledgments The authors acknowledge the financial support of the National Science Foundation of China (21776121), Jiangsu Outstanding Youth Funds (BK20160012), “Jiangsu Shuanchuang” Program,

Thousand

Youth

Talents

Plan,

National

Materials

Genome

Project

(2016YFB0700600), and Nantong Fundamental Research Funds (GY12016040). The numerical calculations were done on the computing facilities in the High-Performance Computing Center (HPCC) of Nanjing University. Supplementary material The Supporting Information is available free of charge on the ACS Publications website at DOI:xxxx. Critical nucleation size, Interface energy estimation, Interface model of Li2O/M (M = Mn or Sn) (Figure S1), Top view of interfaces (Figure S2), M-O bond energy calculation, Crystal stucture used for the calculation (Figure S3), Lattice parameters used for calculation and summary of calculated energies (Table S1). Author information Corresponding author E-mail: [email protected] ACS Paragon Plus Environment

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