In situ Dealloying of Bulk Mg2Sn in Mg-Ion Half ... - ACS Publications

(FEM) to demonstrate a cost-effective route to nanostructured Sn for high ... electrochemical experiments using a Mg half-cell consisting of bulk Mg2S...
1 downloads 0 Views 5MB Size
Download PDF
Article Cite This: Chem. Mater. 2018, 30, 1815−1824

pubs.acs.org/cm

In Situ Dealloying of Bulk Mg2Sn in Mg-Ion Half Cell as an Effective Route to Nanostructured Sn for High Performance Mg-Ion Battery Anodes Hooman Yaghoobnejad Asl, Jintao Fu, Hemant Kumar, Samuel S. Welborn, Vivek B. Shenoy, and Eric Detsi* Department of Materials Science & Engineering, University of Pennsylvania, Philadelphia Pennsylvania 19104-6272, United States S Supporting Information *

ABSTRACT: Nanostructured Sn as negative electrode material in Mg-ion batteries suffers from very slow magnesiation kinetics when its nanoscale feature sizes are not in the sub-100 nm range. Herein, we use electrochemical experiments in combination finite element modeling (FEM) to demonstrate a cost-effective route to nanostructured Sn for high performance Mg-ion battery anodes. Using FEM we found that antagonistic stresses developed during dealloying of Mg2Sn induce pulverization of the dealloyed material and formation of nanostructured Sn with the characteristic feature size in the sub-100 nm range. These results were further confirmed through electrochemical experiments using a Mg halfcell consisting of bulk Mg2Sn particles with an average characteristic size larger than 10 μm as the working electrode, cycled versus Mg metal as counter and reference electrodes, and all-phenyl complex (APC) electrolyte. Ex situ electron microscopy and diffraction techniques were used to study the working electrode material in the pristine, demagnesiated, and remagnesiated forms. The results suggest that the starting micrometer-sized Mg2Sn particles are converted into nanostructured β-Sn with characteristic sizes ranging from 10 to 50 nm during the first demagnesiation. Electrochemical performance of the in situ formed nanostructured Sn was further investigated during subsequent (de)magnesiation cycles in combination with electrochemical impedance spectroscopy (EIS). EIS studies suggest the formation of passive films on the Mg2Sn electrode. A reversible capacity of 300 mAh/g was demonstrated over 150 cycles at the rate of C/5 after application of a combined sequence of regular galvanostatic cycling with an oxidative pulse to control the passive film formation. This work is expected to open new avenues for cost-effective routes to high performance alloytype Mg-ion battery anodes without complex nanosynthesis steps.

■

INTRODUCTION The increasing popularity of Li-ion batteries (LIBs) in portable electronics and electric vehicles is associated with an evergrowing concern about the scarcity of raw lithium resources to maintain the continuous growth of the lithium-ion battery market.1,2 As a consequence, research into promising secondary battery technologies that utilize earth-abundant elements to replace lithium have emerged, among which Mg-ion batteries (MIBs) are widely being investigated as alternatives to LIBs.3−7 Nevertheless, progress toward practical MIBs has been halted partly due to a lack of suitable electrolytes that are compatible with magnesium metal.8−11 In particular, most of the simple salts and/or organic solvents containing unsaturated carbon atoms yield a Mg-ion-blocking passive film on metallic magnesium.12−15 Therefore, only a handful of electrolytes, including the all-phenyl-complex (APC), are compatible with Mg metal as the negative electrode.16−20 Despite its huge success in reversible Mg plating,8,11,18,21,22 the APC electrolyte exhibits several drawbacks including the reactive nature of the salt, the high volatility of the THF solvent, and the relatively limited operating voltage window which makes APC not suitable for high-voltage cathodes.3,20,23 By replacing the © 2018 American Chemical Society

magnesium metal anode with a magnesium alloy that is compatible with conventional electrolytes, these issues can be circumvented. Currently, Bi and Sn are two promising candidate materials for reversible Mg storage in the form of Mg3Bi2 and Mg2Sn, respectively.24−31 The Bi system has been studied in detail, and several reports demonstrate the effectiveness of Mg storage in a Bi host at fast rates, thanks to the high ionic conductivity of Mg2+ ions in the Bi matrix sublattice.29−32 However, the theoretical capacity of Mg3Bi2 is only 330 mAh/g and (de)magnesiation of Bi takes place at a relatively high voltage of 0.28 V (vs Mg2+/Mg). In contrast to Mg3Bi2, Mg2Sn exhibits a much higher theoretical capacity of 903 mAh/g, and the reversible storage of Mg through the halfcell reaction in eq 1 only requires ∼0.15 V (vs Mg2+/Mg),24 making Sn an attractive high energy density anode material for MIBs. Sn + 2Mg 2 + + 4e̅ ↔ Mg 2Sn

(1)

Received: September 30, 2017 Revised: January 7, 2018 Published: January 9, 2018 1815

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

the Results and Discussion section, various aspects of phase, compositional, and microstructural features of the dealloyed and realloyed Mg2Sn electrodes are investigated and the electrochemical cycling is reported. Besides the critical size of the active material, passive film formation on the negative electrode, in particular Mg metal, is known to also play a very critical role in the performance of MIBs.12 Our materials system favored the formation of a passive film in the APC electrolyte as demonstrated by EIS analysis. In our work we successfully demonstrate the use of reversible potential-driven removal of passivating films, to control the passive film formation in MIB and achieve stable cycle life behaviors. This initial work is of primary importance, as it justifies the advantage of starting with Mg2Sn as an active electrode material instead of pure Sn.

Unfortunately, the outstanding features of Sn have been impeded by the slow electrochemical alloying reaction of Sn with Mg at room temperature, so that down-sizing Sn into nanoparticles to enhance the (de)magnesiation kinetics becomes unavoidable.25,26 One notable example is the work of Singh et al. which demonstrated full capacity in Sn nanoparticles with average size smaller than 150 nm. Even with such a small particle size, the full magnesiation of Sn was achieved at a very slow rate of C/500 and the process was only partially reversible with roughly 1/3 of the capacity recovered.24 Besides the reversible magnesiation of pure Sn, several researchers have investigated the suitability of binary Snbased alloys for Mg storage. For example, Kitada et al. studied Sn−X (where X= Cu, In, Pb)33 and showed that low melting Sn alloys (In0.5Sn0.5 and Pb0.6Sn0.4) demonstrate superior performance in (de)magnesiation kinetics compared to the pure Sn in cyclic voltammetry tests, while Cu−Sn intermetallics (Cu6Sn5 and Cu3Sn) are essentially inactive. Furthermore, Parent et al. investigated the Sn−Sb binary system and reported on the utilization of β-SnSb alloy as a precursor for in situ formation of nanosized Sn for effective Mg storage.25 The authors have demonstrated that the β-SnSb alloy effectively pulverizes during the first few cycles, forming Sn-rich and Sbrich magnesiated subdomains with a characteristic size around 33 nm, with only the Sn-rich phase actively participating in subsequent reversible magnesiation. The authors therefore concluded that pure Sn nanoparticles with a characteristic size below 40 nm are essential for fast and fully reversible storage of Mg in the form Mg2Sn. In a follow-up work, Cheng et al. used the proposed β-SnSb alloy methodology to make a highperformance nanoparticulate Sn electrode for MIB, which delivers a specific capacity of 450 mAh/g at C/18 with a capacity retention of ∼74% over 200 cycles, and an excellent high rate capacity retention.26 Despite the highly promising results of binary SnSb alloy as the MIB anode, achievement of a similar electrochemical performance in pure Sn remains as a challenge. In the present work we propose an effective approach to overcome that challenge and reversibly store Mg at practical Crates (C/5 and C/2) for over one hundred cycles in pure nanostructured Sn formed in situ from the Mg2Sn precursor. While stress formation in LIB alloy-type anodes typically leads to their failure and is therefore undesireable,34−39 herein we take advantage of stresses arising during demagnesiation of micrometer-sized bulk Mg2Sn electrode materials to produce nanostructured Sn with a characteristic size in the range of 10− 50 nm. The critical feature of our approach is the chemomechanical coupling associated with the tensile stresses arising during demagnesiation, which promote crack nucleation and propagation in the electrode material. Hence, by starting with bulk Mg2Sn we can effectively produce nanostructured Sn in situ during the first dealloying cycle, without a sacrificial element in the starting Mg2Sn electrode material since the Mg removed from Mg2Sn is reused in subsequent cycles. A key point in our approach is the fact that the sacrificial element is the working ion of the cell, so that dealloying and “half-cell charging” happen concurrently during the first cycle, unlike the usual (electro)chemical dealloying methods where removal of the sacrificial element is merely a method of material preparation for the ultimate electrode. In the first part of the Results and Discussion section, we present a comprehensive study of stress evolution during (de)magnesiation of Sn using a numerical method based on FEM. In the subsequent parts of

■

RESULTS AND DISCUSSION Micromechanics Study of Sn Anode during (De)magnesiation. We used FEM to simulate the mechanical fate of reaction products formed as a result of (de)alloying the Mg2Sn/Sn system. It should be emphasized that previous theoretical calculations on magnesium transport kinetics aimed to extract the activation energy of Mg2+ ions diffusion through bulk Sn using DFT-based routines such as Nudged Elastic Band (NEB).40,41 The results of such calculations do not necessarily reflect the real physics involving the alloying reaction of β-Sn with Mg to form Mg2Sn, since either a hypothetical cubic α-Sn has been used as the host model or the tetragonal β-Sn crystal structure has been modified by forcefully inserting Mg ions in the lowest energy sites. The ion dynamics in intercalation-type electrode materials based on solid-state diffusion of the working ion could be fundamentally different from alloy-type anode materials where the oxidized (i.e., β-Sn) and reduced (i.e., Mg2Sn) forms of the anode take completely different crystal structures, namely tetragonal and cubic for β-Sn and Mg2Sn, respectively. In real situations where the crystal structure of the active electrode material changes during alloying and dealloying, it seems more appropriate to consider (de)magnesiation processes as a surface-limited reaction, where the ratedetermining step is the electrochemical redox reaction of the material at the progressively moving phase boundary, as has been demonstrated previously for the well-studied case of silicon lithiation.42,43 In the present work we are aiming to demonstrate that chemo-mechanical coupling-induced crack nucleation and propagation is a viable mechanism for successful cell operation, which can overcome the very slow diffusion of Mg2+ in the Mg2Sn host lattice. It is to be emphasized that cracking of the electrode active material is the major cause of capacity decay, especially in alloy-type anode materials where the volume change due to cell cycling can reach several hundred percent. Yet, it has to be shown here that this very same destructive phenomenon could in principle open up fast Mg-ion transfer pathways across the electrode−electrolyte interface when bulk Mg2+ ion diffusion in Mg2Sn is very slow and inefficient. Indeed, such cracks obviously contribute to a persistent contact between the unreacted electrode material and the electrolyte by establishing fast Mg2+ ion transport routes to reaction sites at the Mg2Sn−electrolyte interface. Cheng et al. analytically studied crack formation in LIB electrodes by assuming elastic stresses during (de)lithiation reactions.44,45 Here, we extend that concept of stress-induced material cracking to MIB anodes, particularly β-Sn, by accounting for plastic deformations. Furthermore, we assume a biphasic alloytype reaction associated with the (de)magnesiation of Sn 1816

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

Figure 1. Distribution of von Mises stress (left column), hoop stresses (middle column), and radial stresses (right column) during demagnesiation and magnesiation processes in Mg2Sn/Sn particles. (a) Both Mg2Sn and Sn have been modeled as perfectly elastic material. (b and c) Sn has been modeled as plastic material while Mg2Sn has been modeled as elastic material during demagnesiation and magnesiation, respectively.

It is evident from these simulations that the hoop stress near the surface remains tensile but maximum stress is much smaller compared to perfectly elastic deformations considered in the previous case. The critical size (dcr) for the fracture to be energetically favorable in plastic materials is given by42

particles as the ultimate case for a solid-solution problem in which ionic diffusion is infinitesimally small and there is a sharp concentration change at the phase boundary, rather than a concentration gradient. Under these assumptions, progressive demagnesiation is modeled as the propagation of a sharp interface between the growing demagnesiated shell of Sn (gray) and shrinking magnesiated core of Mg2Sn (purple) in a core− shell structure, as shown in Figure 1. The distributions of stresses obtained under conditions of mechanical equilibrium for the core−shell structure with a 5% radially demagnesiated shell are shown in Figure 1. Principal stresses along radial and hoop directions, together with von Mises stress (defined as

σv =

3 σ ′σ ′ 2 ij ij

dcr = ΓE /Zσ Y2

(2)

where Γ is the fracture energy of the material, E is the elastic modulus, σY is the yield strength, and Z is a dimensional parameter uniquely determined by the material and depth of the crack whose value varies between 0 to 1. In the case of Sn particles, we use the values of 0.3 J/m2 for Γ, 42 GPa for E, 800 MPa for σY,46 and 1.0 for Z.47 Substituting these values in eq 2 gives a critical size of dcr = 22 nm, which is in agreement with our experimentally observed particle size of ∼10−50 nm. Next, we investigate the stress evolution in Sn particles during the magnesiation process. Our simulation predicts that large compressive hoop stresses are produced in the magnesiated shell due to the insertion of Mg-ions (see Figure 1c). Such compressive stresses can impede the reaction of Mg ions with Sn particles on account of subsequent insertion of Mg2+ being very slow, as previously observed during lithiation of silicon.48 In summary, the chemo-mechanical coupling in our materials system impacts the (de)magnesiation kinetics in two different ways: On the one hand, compressive stresses arising in the Mg2Sn shell of the Sn active electrode material during magnesiation do not favor further magnesiation. This may partly justify the very slow kinetics reported during the conversion of Sn into Mg2Sn, even for particles as small as ∼150 nm.24 On the other hand, tensile stresses arising in the Sn shell during demagnesiation of Mg2Sn induce cracking of the active electrode material, allowing the electrolyte to flow through these cracks for further demagnesiation. Hence, these tensile stresses favor the demagnesiation step. In other words, it should be possible to make Mg-ion battery anodes starting from micrometer-sized Mg2Sn particles as active electrode materials instead of pure Sn. Demagnesiation of such relatively large Mg2Sn particles is mechanically favorable and should result in nanometer-sized Sn structures. The in situ formed Sn

where σ′ij are the components of deviatoric stress

tensor) are illustrated. The compressive stress due to electrochemical demagnesiation is accommodated by the elastic-plastic deformations in the core−shell structure. To understand the nature of stress distributions, we first consider a simplified picture and model both Sn and Mg2Sn as linear elastic materials. The model predicts that excessive tensile hoop stresses up to 34 GPa are generated during demagnesiation of Mg2Sn in this case, as shown in Figure 1a. The elements at the surface of a demagnesiating Mg2Sn particle accommodate compressive stress by shrinking along the radial direction, but undergoes large tensile hoop stresses. The tensile hoop stress of such magnitude will cause immediate pulverization. A simple analysis based on the linear elastic fracture mechanics theory gives a typical size of 3 Å for fractured particles. However, as it will be shown in a latter section, the typical size of a particle after demagnesiation observed from our experiments is in the range ∼10−50 nm. To understand this discrepancy between our model (∼3 Å) and our experiment (∼10−50 nm), we observe that the mechanical behavior of the Sn is best described as a plastic material with a Yield strength (σY) of ∼800 MPa. Once we include the perfectly plastic behavior in our simulations, the Sn shows tensile plastic yielding after von Mises stress reaches the yield stress of Sn. The stress distribution for the same demagnesiation radius but with plasticity for Sn is shown in the Figure 1b. 1817

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

along with the electrochemically cycled electrodes and the corresponding voltage profiles, respectively (see the Experimental Section for electrode preparation). The electrochemical tests have been carried out between 0.6 V and 5 mV vs Mg2+/ Mg at a relatively slow rate of C/50 to maximize phase conversions. The as-prepared material is composed mainly of cubic Mg2Sn as shown by the XRD pattern (i) in Figure 2a (ICSD col. code: 151368). Following charging the electrode to an upper cutoff potential of 0.6 V (see Figure 2b point (ii)), the material dealloys effectively into β-Sn as depicted by the XRD pattern (ii) in Figure 2a (ICSD col. code: 52269). A subsequent reduction down to a lower cutoff potential of 5 mV vs Mg2+/ Mg (see Figure 2b point (iii)) causes the diffraction peaks of Mg2Sn to grow back again, with some Sn left unreacted (see XRD pattern (iii) in Figure 2a). The unreacted Sn could be assigned to the effect of pulverization upon the first demagnesiation as discussed above, which causes some of the particles to be electrically isolated from the current collector. Also the variation of Mg2Sn phase relative peak intensities for (111): (200) reflections (at 2θ values 22.76° and 26.34°, respectively) among the as-synthesized material (Figure 2a(i)) and electrochemically grown Mg2Sn (Figure 2a(iii)) could be explained based on the preferred orientation and perfect cleavage of Mg2Sn fluorite-type crystal along the (h00) family of planes as a result of ball-milling. Note from the inset of Figure 2a that the Bragg peaks at high diffraction angles associated with electrochemically formed Mg2Sn are broadened compared to the as-prepared Mg2Sn from the XRD pattern (i). Application of Scherrer’s equation to the (333) reflections at about 2θ = 73° indicates an average crystal domain size of 36 nm (calculations provided in the SI). This implies that (i) the resulted average crystallite size value is very close in size to both the FEM calculations and TEM images and (ii) significant particle size reduction occurs during the first cycle.

nanostructured can then be cycled subsequently. The advantage of such a top-down nanosynthesis starting from bulk Mg2Sn is at least twofold: (i) Processing micrometer-sized Mg2Sn powder for large scale applications is more affordable than synthesizing nanostructured Sn using conventional bottom-up routes.49−52 (ii) In a full battery configuration, Mg removed from Mg2Sn can be reversibly stored in the cathode, which means that the loss of material in the process can be avoided or minimized. Hence, starting from Mg2Sn could be more attractive than previous methodologies where sacrificial Sb is used to produce nanostructured Sn in situ from β-SnSb electrode materials.25,26 In the following sections, we present the characterizations and electrochemical performances of Mg2Sn. Crystallography and Microstructural Studies. Figure 2a and 2b show the XRD patterns of the as-synthesized Mg2Sn

Figure 2. (a) XRD pattern of as-prepared Mg2Sn electrode (i), the first demagnesiated (ii), and the first magnesiated (iii) states. (b) The corresponding voltage−capacity profile. Inset in panel (a) shows a magnified view of the high angle region.

Figure 3. SEM micrographs of the Mg2Sn electrode at various states during the first electrochemical cycle versus Mg metal. (a) Pristine Mg2Sn in a mixture with PVDF and carbon nanofibers. (b) Electron micrograph image of the Mg2Sn electrode demagnesiated up to 0.6 V (vs Mg2+/Mg). (c) Electron micrograph image of the remagnesiated electrode down to 5 mV (vs Mg2+/Mg). Panels (d−e) and (f−g) show the elemental distribution of Mg (green) and Sn (red) for the electron micrographs shown in (b) and (c), respectively. 1818

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

Figure 4. SEM micrographs showing the microstructural features of the Mg2Sn composite electrode at low (a) and high (b) magnifications following the first demagnesiation from open circuit voltage up to 0.6 V versus Mg/Mg2+. Nanoporosity can be observed at high magnifications on dealloyed Sn grains (b).

Figure 5. TEM micrographs of the demagnesiated (a−c) and remagnesiated electrode (e−g). Panels (d) and (h) show the Fourier transform of the enclosed areas of panel (c) and (g), respectively.

essential to discriminate between aggregates and absolute particle size. For example, the EDX map in Figure 3c shows that the elemental distribution of Sn and Mg over large aggregates is very uniform and chemically indistinguishable (cf. Figure 3b), which is a clear indication that the aggregates have been formed from very small particles with sizes below the resolution limit of the EDX. The possibility to effectively remagnesiate sub-100-μm-sized Sn particles (Figure 3b) formed in situ to Mg2Sn (Figure 3c) may initially appear counterintuitive and self-conflicting when considering the strict requirement of 30−40 nm particle size for effective magnesiation of Sn. However, detailed investigation of the morphological characteristic of the Sn using high-resolution SEM (Figure 4) reveals that the Sn particles are formed as nanoporous structures incorporating interpenetrating Sn nanostructured (i.e., ligaments) and open nanochannels (i.e., pores). Indeed, in Figure 4b the darker contrast of the enclosed areas with the dashed lines indicate void spaces or a material of considerably lower density (i.e., Z-contrast) like carbonaceous materials, which in turn proves the existence of nanometric voids or “pores” within the inhomogeneous material particle (brighter contrast). The characteristic size of both ligaments and pores are on the order of a few tens of nanometers. This demonstrates for the first time that porosity spontaneously formed in situ during demagnesiation in Mg cells, as previously reported by several authors in Li-based systems.53,54 Ex situ formation of porosity in Sn from Mg−Sn alloys has been

Electron microscopy can provide comprehensive information regarding the morphology and compositional evolution of the alloy-type anode material. SEM images of pristine Mg2Sn as cast in the electrode with binder and conductive additives are shown in Figure 3a, and the corresponding EDX spectrum is provided in Figure S1. Note that the fluorine (F) and carbon (C) signal from the EDX originate from the binder and conductive additive in the composite electrode. The pristine Mg2Sn is a polydispersed collection of micrometer-sized particles in the range 1−50 μm. Figure 3b shows the SEM images of the composite electrode after first demagnesiation. Based on this micrograph, the sample is mainly composed of large Sn particles and some unreacted Mg2Sn. Based on these results, Mg2Sn with a typical size as big as tens of μm can be demagnesiated up to roughly 67% of the initial Mg content at the rate of C/50, in contrast to the very slow kinetics of the reversed process. Several of our attempts to magnesiate microand submicrometer-sized pure Sn at similar and slower rates were unsuccessful. Singh et al. have shown that full magnesiation of pure Sn nanoparticles with a characteristic size of ∼150 nm was achieved at the slow rate of C/500.24 As discussed in the previous section, tensile stresses arising in the active material during demagnesiation enhance the process, making it possible to demagnesiate micron-sized particles, while compressive stresses induced by magnesiation retard the process. Figure 3c suggests that remagnesiation of the in situ formed Sn creates aggregates of 1 μm or larger. However, it is 1819

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials demonstrated previously by chemical dealloying.55,56 The nanoporous Sn architecture formed allows for facile subsequent magnesiation of the Sn ligaments to Mg2Sn, as they are within the effective size range to be active for electrochemical alloying. It should be emphasized that even though SEM images have been used in this work for microstructural studies, we have relied heavily on TEM for particle size determination. SEM has been used merely as a tool to adjudicate the maximum possible particle sizes, spatial elemental distribution, and degree of particles agglomeration. To further investigate the effect of electrochemical (de)magnesiation on the pristine Mg2Sn electrode, TEM has been used to analyze the particle sizes at the highest magnifications. Figure 5 shows representative micrographs of the electrochemically demagnesiated and remagnesiated specimens. The morphology of the in situ formed Sn after the first demagnesiation (Figure 5a−c) depicts pulverized Sn particles with maximum diameters around 50 nm. By comparison with SEM results (Figures 3b and 4) it is possible to conclude that, as a result of dealloying, Sn forms as a composite mixture of highly porous structures with fractured particles, altogether trapped in the binder matrix. Furthermore, Figure 5e−g shows the morphology of a mixture of Sn and Mg2Sn particles, which follow the same size distribution trend as those of Sn formed during the first demagnesiation. The Fourier transformation of high resolution images containing lattice fringes unequivocally identifies Sn and the Sn−Mg2Sn mixture for the demagnesiated and remagnesiated samples, respectively, a result also observed in the XRD patterns (Figure 2a(iii)). According to the TEM results, it is possible to deduce the “critical size” for magnesiation to be in the range of 1050 nm, which corroborates the finite element simulations elaborated on above, and is consistent with reports on β-SnSb electrode material.25 It should be emphasized that the particle size reduction in our approach was achieved at a relatively fast rate without using a sacrificial element since Mg removed from the starting electrode material is reused in the cell during subsequent cycling. Electrochemical Performance. The electrochemical performance of the synthesized Mg2Sn and the subsequently in situ derived nanostructured Sn were investigated in a coin-cell configuration. Further details on the Mg2Sn slurry electrode preparation and coin-cell assembly are provided in the Experimental Section. Figure 6 shows the typical first four (de)magnesiation voltage profiles obtained during galvanostatic cycling at the rate of C/5 in the voltage range between 0.6 V

and 5 mV vs Mg2+/Mg. Demagnesiation and magnesiation occur at the expected average potentials of 0.26 and 0.15 V vs Mg2+/Mg, respectively, except for cycle #1 where they take place at average potentials of 0.38 and 0.12 V vs Mg2+/Mg, respectively. During the first demagnesiation, Mg2Sn used on the working electrode is dealloyed into β-Sn, while magnesium deposition takes place on the Mg metal counter electrode. At C/5, the first demagnesiation process yields a capacity of 400 mAh/g based on the mass of Sn, which corresponds to ∼44% of the theoretical capacity of Sn (903 mAh/g). Note in Figure 2b from the previous section that a much higher capacity of ∼606 mAh/g corresponding to ∼67% of the theoretical capacity of Sn was obtained during the first demagnesiation at a slower rate of C/50. This means that when going from C/50 to C/5, the 10-fold increase in the (dis)charging current density reduces the specific capacity with ∼206 mAh/g. Based on XRD and electron microscopy data, one may relate the observed high overpotential during the first demagnesiation to the large particle size of the starting Mg2Sn material. In subsequent cycles during which the average particle size of Mg2Sn has diminished compared to the pristine one (Figures 2a inset, 3c, 4, and 5e−g), demagnesiation occurs at lower overpotentials. Please note that the same trend is evidently reflected in the impedance behavior of the system, as it will be shown in a later subsection. It should be pointed out that the capacity associated with the first magnesiation represents only ∼60% of the initial demagnesiation coulombic charge. Such an irreversible behavior was observed to be rate-independent, meaning that even cells cycled at slower rates (like C/50 from Figure 2b) only delivered ∼60% of their initial demagnesiation capacity. This behavior can be attributed to the loss of electrical contacts between some pulverized active materials and the conductive additives and/or current collector, resulting in their exclusion from the electrochemical process. Following the first few cycles with stable capacity, we observed a gradual capacity decay which generally leads to an eventual cell deactivation after the ∼20th cycle as shown in Figure 7a. To study the origin of this capacity loss, electrochemical impedance spectroscopy (EIS) in the frequency range of 10 mHz to 1 MHz was used to follow the trend of changes in the cell components. Figure 7b shows the collected EIS data acquired by charging the cells to 0.6 V vs Mg2+/Mg (demagnesiation of Mg2Sn to β-Sn with a state-ofcharge (SOC) of 100%) and following a 60 s relaxation (equilibration period) at the cell OCV. In general, the impedance response in the Nyquist plot is composed of a low frequency curved tail associated with the nucleation of Mg2Sn on Sn, and a high frequency depressed semicircle related to charge transfer and reaction resistance. As discussed above from Figure 6a, the first magnesiation requires a higher polarization overpotential than subsequent cycles. It can be seen from the red trace curve in Figure 7b that the corresponding impedance |Z| is reduced from 3.22 kΩ down to a minimum of 920 Ω at the end of the sixth cycle (blue trace) with eventual constant growth until the 30th cycle (3.20 kΩ, yellow trace). Interestingly, after the sixth cycle while the length of the low frequency tail is increased, the medium-high frequency semicircle remains constant. It is worth mentioning that our attempts to fit the curved low frequency tail of the impedance response with the Warburg element has proved unsuccessful. Therefore, the low frequency curve is probably

Figure 6. Typical first four (de)magnesiation cycles of Mg2Sn|APC|Mg cell at a C-rate of C/5. 1820

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

uncontrolled growth of the passive film or its complete removal by a long oxidative pulse leads to premature cell failure (Figure S3). Note that various additives such as FEC and VC have been proposed and successfully tested in alloy-forming anode materials for Li-ion and Na-ion batteries.57−62 It is suggested that these additives upon contact with the anode active material form stable protecting films which minimizes the exposure of the fresh surface of the anode active material to the electrolyte, thereby enhancing the cycle-life behavior of this class of negative electrodes.63 Despite the fact that such additives are missing or are under investigation for MIB, our results suggest that passive films which may form autonomously on the electrodes could act as a protective layer. The cell performance under regular cycling (i.e., no oxidative pulse) and constant current charging with application of pulses at two different C-rates is compared in Figure S4. It worth mentioning that the formation of passivating films on the Sn electrode in carbonate-based solvents and LiPF6 has been demonstrated and studied previously.64,65 Even though the APC electrolyte system used in the present study is different from that of LIBs, the underlying principles may be the same. Further investigations are currently underway to characterize the nature and chemical features of the passive film and this follow-up work will be submitted elsewhere for publication. Our speculation regarding the formation of the passive film is the chemical reduction of species present in the APC electrolyte on the Mg2Sn electrode, making the corresponding oxidation process and dealloying of Mg2Sn to Sn difficult by blocking Mg2+ ion diffusion out into the electrolyte. Our argument is supported by the trend of the voltage−composition profiles (Figure S5) of Mg2Sn electrode cycled at C/5; following the second cycle the charge transfer due to demagnesiation becomes successively smaller, meaning that at each cycle less Mg2Sn is being dealloyed into Sn compared to the previous cycle, which ultimately leads to a constant capacity decay. Therefore, it seems plausible that if less Sn would exist for reduction in the subsequent cycle, an application of a highpotential pulse would be necessary to force oxidation of Mg2Sn into Sn, as successfully demonstrated above. Figure 8 shows the comparison of capacity retention at two relatively fast rates of C/5 and C/2, following the application of the oxidizing pulses during (de)magnesiation cycles. As discussed above, under these conditions (application of

Figure 7. Initial capacity retention of Mg2Sn|APC|Mg cell (a). The corresponding EIS spectra of cells charged to 0.6 V vs Mg2+/Mg (b). A few representative cycles have been highlighted with different colors.

not a typical bulk solid-state diffusion process; however, due to the low frequency (∼10 mHz) nature, it may be assigned to the nucleation/growth (and dissolution during the oxidative phase of AC wave) of Mg2Sn crystallites, which are considered as relatively slow processes. Passivation of the magnesium electrode/electrolyte interface is known to be a critical issue in MIB;12 therefore, we have hypothesized that a similar mechanism could be affecting the Mg2Sn electrode, so that the gradual increase in the low frequency impedance after the sixth cycle originates from an increase in resistance to interconvert Mg2Sn into Sn, presumably due to passive film formation. In order to verify our hypothesis, we proceeded by testing the performance of our material system under conditions where, in addition to the regular charging/discharging constant current densities applied in the voltage range between 0.6 V and 5 mV, an additional high current density pulse with the upper cutoff voltage of 2.5 V vs Mg2+/Mg was applied for 5 s at the end of each demagnesiation step in order to electrochemically destroy any passive film that might have been formed at the surface of active material during and past magnesiation. The typical oxidation current profile along with the potential response for a single cycle is depicted in Figure S2. Interestingly, the material immediately recovers a considerable fraction of its capacity in the subsequent cycle, meaning that our hypothesis is justified. However, after application of the pulses, on average more than 20 subsequent cycles are required before the capacity stabilizes. These preconditioning cycles are vital for a stable cycle life, as it maintains a thin layer of protective passive films on the surface of electrode active materials. According to our observations,

Figure 8. Capacity retention of Mg2Sn|APC|Mg cells following application of oxidative pulses and past the preconditioning steps at C/5 and C/2. 1821

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials oxidative pulses to destroy the passive film) a gradual increase in capacity is observed in each case. The origin of this increase in capacity cannot be attributed to a further pulverization of the active electrode material since the maximum capacity achieved after this increase is comparable to, if not slightly lower than, the capacity values obtained from the starting material during the first five cycles (see supporting Figures S3 and S4 for the complete capacity spectrum). We speculate that the specific capacities at C/5 and C/2 from Figure 7 grow gradually as a result of stabilization of the SEI surface film, until it reaches a maximum value which corresponds to the optimal thickness of the surface film. Under these conditions the Sn electrode yields a reproducible capacity of ∼300 mAh/g, which is stable for 150 cycles at C/5 (0.18 A/g). The cell cycled at C/2 (0.45 A/g) shows a maximum specific capacity of 180 mAh·g−1, which decays to 150 mAh·g−1 after 150 cycles. Despite the fact that only a fraction of the theoretical capacity has been achieved at each C-rate, these results are particularly important in the sense that they show fast cycling of Sn for an extended number of cycles starting initially from large micrometer-sized particles, utilizing only Sn and Mg and no sacrificial element. In general, the results presented here suggest that Sn may be regarded as not only a high energy density but also a suitable high power density anode material for MIB, provided that the right selection is made on the choice of the starting material form. We believe that such a concept may be expanded and applied to other alloy-type anode materials, since the principle of crack formation and material pulverization during cycling is quite general in high-capacity alloy-type anodes. Note that although APC was used in the present work, the ultimate goal of alloy-type anodes would be to expand the full window of conventional electrolytes to the MIB technology, and we expect our material system to work in conventional electrolytes. Indeed, the concept presented in our work should not differ fundamentally from other reported works on Sn nanoparticles, where it has been shown that these Sn nanoparticles can be reversibly magnesiated using a conventional electrolyte.24

nanostructured Sn as the Mg-ion battery anode, (iii) the possibility to control passive film formation in Mg-ion cells. We anticipated that the new concept demonstrated here (starting from magnesiated alloy) can be extended to other alloy-type anode materials.

■

EXPERIMENTAL SECTION

Theoretical Methods. We adopt a linear elastic and a perfectly plastic material models for Mg2Sn and Sn. Perfect plasticity was modeled using the von Mises yield criterion. (De)magnesiation induced strains were modeled as the thermal expansion/compression of the shell in the core−shell structure. The elastic/plastic material models coupled with intercalation stresses were implemented in the finite element package, COMSOL. Axial symmetry was enforced to reduce computational cost. For Mg2Sn, Young’s modulus E = 61.9 GPa and Poisson’s ratio ν = 0.36. For Sn, E = 42 GPa, ν = 0.24, and σY = 800 MPa. Reagents. Sn shots (Sigma-Aldrich, 99.8%), Mg foil (SigmaAldrich 99.9%) and chip (99.98% Sigma-Aldrich), Polyvinylidene fluoride (PVDF, >99.5%, MTI), N-methyl-2-pyrrolidone (NMP, anhydrous, 99.5%, Acros Organics), carbon nanofiber (Sigma-Aldrich), graphene nanosheets (Sigma-aldrich) and conductive carbon black (Timcal), phenyl magnesium chloride (PhMgCl, 2 M in THF, SigmaAldrich), aluminum chloride (AlCl3, anhydrous, 99.99%, SigmaAldrich), and Cu foil (MTI) were used as received. Tetrahydrofuran (THF, >99.9% anhydrous, Sigma-Aldrich) has been further dehydrated by treatment with molecular sieves (4 Å, Fisher) prior to use. Synthesis. About 3.9 g of Mg2Sn were synthesized by melting 2.84 g of Sn shots slightly off-stoichiometric (Sn excess) and 1.10 g of Mg chips using a graphite boat in a tube furnace for 0.5 h at 900 °C under argon flow, followed by 1 h at 800 °C and naturally cooling down to room temperature. In a variation one may start also by taking excess Mg and remove the extra portion by evaporation at 900 °C so that the final product is slightly rich in Sn on average composition. In either case the Mg2Sn contains excess solidified Sn droplets on the surface (see Figure S6), which is loosely attached to the solidified alloy and can be removed easily. Material Characterization. XRD patterns of the as prepared Mg2Sn and the electrochemically demagnesiated and magnesiated electrodes have been obtained using a Rigaku Geiger Flex horizontal goniometer diffractometer, equipped with a graphite monochromator and using the Kα1 line of a Cu X-ray tube. All patterns have been collected on a Bragg angle range of 10°−90° at steps of 0.1° and a scan rate of 0.5° min−1. SEM images of the various specimens have been studied using a JEOL 7500F HRSEM. The energy dispersive X-ray microanalysis of the samples has been carried out using an Oxford Instruments EDX Li−Si drift detector. Transmission electron microscopy: TEM micrographs of the electrochemically oxidized (demagnesiated) and reduced (magnesiated) phases have been obtained using a JEOL 2010F TEM/STEM at 200 kV acceleration voltage. The electrodes have been extracted from the cycled cell and washed to remove the electrolyte salts, separated from contact with a current collector and loaded on bare copper TEM grids. Electrochemical Measurements. The prepared Mg2Sn sample has been crushed into a powder and ball-milled in an argon-filled sealed Teflon cup for 6 h to reduce the particles sizes using a Spex 8000 mixer/mill machine. The powder thus formed has been mixed and milled further with carbon nanofiber, graphene nanosheets carbon black, and PVDF in a 5:1:1:1:2 ratio, and NMP has been added to the mixture to dissolve the PVDF and form a uniform slurry, which has been casted on a copper current collector as a thin film using a blade film applicator and dried under vacuum. Due to the reactivity of Mg2Sn with atmospheric gases, all the electrode preparation steps have been carried out under an inert atmosphere in an argon-filled glovebox (MBraun) with H2O and O2 content below 0.1 ppm. Circular disks of the Mg2Sn working electrode with active material loading in the range 1.0−2.0 mg/cm2 were cut and assembled into 2032 type coin cells

■

CONCLUSION In summary we have demonstrated the formation of nanostructured Sn in situ in Mg-ion half cells through electrochemical demagnesiation of micrometer-sized Mg2Sn. Using numerical methods, we have shown that material pulverization is highly directional and occurs only during dealloying of Mg2Sn into Sn, while the reverse process is largely restrictive, in agreement with previous reports on the very poor magnesiation kinetics of pure Sn. XRD data confirmed that ∼70% of Mg2Sn can be electrochemically demagnesiated at C/50, and electron microscopy studies show that demagnesiation of Mg2Sn gives rise to Sn nanoparticules as well as nanoporous Sn structures. In situ formed nanostructured Sn was subsequently cycled at high rates for this class of materials, unprecedented in pure bulk Sn electrodes. Stable capacities of 300 and 170 mAh·g−1 over 150 cycles were demonstrated at C/5 and C/2, respectively. We also encountered the possibility of passive film formation on the working electrode, which degrades the capacity retention. Hence a potential-driven oxidative pulse procedure was successfully applied to minimize the effect of passive film, thus enhancing the long cycle life performance of the Sn electrode. The results reported here are the first of their kind showing the following: (i) The in situ formation of nanostructured Sn from Mg2Sn powder prepared by a rudimentary milling process; (ii) the good performance of the synthesized 1822

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials

(9) Gofer, Y.; Pour, N.; Aurbach, D. Lithium Batteries: Advanced Technologies and Applications; John Wiley & Sons, Inc.: Hoboken, NJ, USA, 2013. (10) Yoo, H. D.; Shterenberg, I.; Gofer, Y.; Gershinsky, G.; Pour, N.; Aurbach, D. Mg Rechargeable Batteries: An on-Going Challenge. Energy Environ. Sci. 2013, 6, 2265−2279. (11) Aurbach, D.; Weissman, I.; Gofer, Y.; Levi, E. Nonaqueous Magnesium Electrochemistry and Its Application in Secondary Batteries. Chem. Rec. 2003, 3, 61−73. (12) Lu, Z.; Schechter, A.; Moshkovich, M.; Aurbach, D. On the Electrochemical Behavior of Magnesium Electrodes in Polar Aprotic Electrolyte Solutions. J. Electroanal. Chem. 1999, 466, 203−217. (13) Tutusaus, O.; Mohtadi, R.; Singh, N.; Arthur, T. S.; Mizuno, F. Study of Electrochemical Phenomena Observed at the Mg Metal/ Electrolyte Interface. ACS Energy Lett. 2017, 2, 224−229. (14) Brown, O. R.; McIntyre, R. The Magnesium and Magnesium Amalgam Electrodes in Aprotic Organic Solvents a Kinetic Study. Electrochim. Acta 1985, 30, 627−633. (15) Aurbach, D.; Pour, N. Corrosion of Magnesium Alloys; Woodhead Publishing Ltd.: 80 High Street, Sawston, Cambridge CB22 3HJ, UK, 2011. (16) Aurbach, D.; Schechter, A.; Moshkovich, M.; Cohen, Y. On the Mechanisms of Reversible Magnesium Deposition Processes. J. Electrochem. Soc. 2001, 148, A1004−A1014. (17) Aurbach, D. Magnesium Deposition and Dissolution Processes in Ethereal Grignard Salt Solutions Using Simultaneous EQCM-EIS and In Situ FTIR Spectroscopy. Electrochem. Solid-State Lett. 1999, 3, 31−34. (18) Mizrahi, O.; Amir, N.; Pollak, E.; Chusid, O.; Marks, V.; Gottlieb, H.; Larush, L.; Zinigrad, E.; Aurbach, D. Electrolyte Solutions with a Wide Electrochemical Window for Rechargeable Magnesium Batteries. J. Electrochem. Soc. 2008, 155, A103−A109. (19) Mohtadi, R.; Matsui, M.; Arthur, T. S.; Hwang, S. J. Magnesium Borohydride: From Hydrogen Storage to Magnesium Battery. Angew. Chem., Int. Ed. 2012, 51, 9780−9783. (20) Zhang, Z.; Cui, Z.; Qiao, L.; Guan, J.; Xu, H.; Wang, X.; Hu, P.; Du, H.; Li, S.; Zhou, X.; Dong, S.; Liu, Z.; Cui, G.; Chen, L. Novel Design Concepts of Efficient Mg-Ion Electrolytes toward HighPerformance Magnesium−Selenium and Magnesium−Sulfur Batteries. Adv. Energy Mater. 2017, 7, 1602055. (21) Pour, N.; Gofer, Y.; Major, D. T.; Aurbach, D. Structural Analysis of Electrolyte Solutions for Rechargeable Mg Batteries by Stereoscopic Means and DFT Calculations. J. Am. Chem. Soc. 2011, 133, 6270−6278. (22) Aurbach, D.; Suresh, G. S.; Levi, E.; Mitelman, A.; Mizrahi, O.; Chusid, O.; Brunelli, M. Progress in Rechargeable Magnesium Battery Technology. Adv. Mater. 2007, 19, 4260−4267. (23) Ha, S. Y.; Lee, Y. W.; Woo, S. W.; Koo, B.; Kim, J. S.; Cho, J.; Lee, K. T.; Choi, N. S. Magnesium(II) Bis(trifluoromethane Sulfonyl) Imide-Based Electrolytes with Wide Electrochemical Windows for Rechargeable Magnesium Batteries. ACS Appl. Mater. Interfaces 2014, 6, 4063−4073. (24) Singh, N.; Arthur, T. S.; Ling, C.; Matsui, M.; Mizuno, F. A High Energy-Density Tin Anode for Rechargeable Magnesium-Ion Batteries. Chem. Commun. (Cambridge, U. K.) 2013, 49, 149−151. (25) Parent, L. R.; Cheng, Y.; Sushko, P. V.; Shao, Y.; Liu, J.; Wang, C. M.; Browning, N. D. Realizing the Full Potential of Insertion Anodes for Mg-Ion Batteries through the Nanostructuring of Sn. Nano Lett. 2015, 15, 1177−1182. (26) Cheng, Y.; Shao, Y.; Parent, L. R.; Sushko, M. L.; Li, G.; Sushko, P. V.; Browning, N. D.; Wang, C.; Liu, J. Interface Promoted Reversible Mg Insertion in Nanostructured Tin-Antimony Alloys. Adv. Mater. 2015, 27, 6598−6605. (27) Mohtadi, R.; Mizuno, F. Magnesium Batteries: Current State of the Art, Issues and Future Perspectives. Beilstein J. Nanotechnol. 2014, 5, 1291−1311. (28) Nguyen, D.-T.; Tran, X. M.; Kang, J.; Song, S.-W. Magnesium Storage Performance and Surface Film Formation Behavior of Tin Anode Material. ChemElectroChem 2016, 3, 1813−1819.

with Mg metal as the counter and reference electrode, binder-free glass microfiber as the separator, and APC electrolyte. The APC electrolyte has been prepared following the reported method by reacting PhMgCl and AlCl3 in THF in a 2:1 ratio so that the final complex salt concentration was 0.4 M.18 All electrochemical tests have been carried out using a Bio-Logic VMP-300 multichannel potentiostat/galvanostat/EIS. Samples extracted from coin cells have been washed with dry THF in an argon-filled glovebox and allowed to dry before taking them out of the glovebox for analysis by XRD and electron microscopy.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b04124. Grain size calculation using Scherrer’s equation, EDX spectrum, galvanostatic charge−discharge sequence with the oxidative pulse and capacity retention of the Mg2Sn electrode at C/5 and C/2 with and without application of pulses and the photographs of the prepared Mg2Sn samples (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Hemant Kumar: 0000-0003-4339-5711 Eric Detsi: 0000-0002-4009-7260 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are thankful to Penn Engineering for the financial support through the PI startup. This project was also supported through NSF grant: NSF CMMI1363203.

■

REFERENCES

(1) Larcher, D.; Tarascon, J.-M. Towards Greener and More Sustainable Batteries for Electrical Energy Storage. Nat. Chem. 2015, 7, 19−29. (2) Tahil, W. The Trouble with Lithium Implications of Future PHEV Production for Lithium Demand. http://evworld.com/library/ lithium_shortage.pdf. (3) Muldoon, J.; Bucur, C. B.; Oliver, A. G.; Sugimoto, T.; Matsui, M.; Kim, H. S.; Allred, G. D.; Zajicek, J.; Kotani, Y. Electrolyte Roadblocks to a Magnesium Rechargeable Battery. Energy Environ. Sci. 2012, 5, 5941−5950. (4) Muldoon, J.; Bucur, C. B.; Gregory, T. Quest for Nonaqueous Multivalent Secondary Batteries: Magnesium and beyond. Chem. Rev. 2014, 114, 11683−11720. (5) Saha, P.; Datta, M. K.; Velikokhatnyi, O. I.; Manivannan, A.; Alman, D.; Kumta, P. N. Rechargeable Magnesium Battery: Current Status and Key Challenges for the Future. Prog. Mater. Sci. 2014, 66, 1−86. (6) Byeon, A.; Zhao, M. Q.; Ren, C. E.; Halim, J.; Kota, S.; Urbankowski, P.; Anasori, B.; Barsoum, M. W.; Gogotsi, Y. TwoDimensional Titanium Carbide MXene As a Cathode Material for Hybrid Magnesium/Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2017, 9, 4296−4300. (7) Choi, J. W.; Aurbach, D. Promise and Reality of Post-Lithium-Ion Batteries with High Energy Densities. Nat. Rev. Mater. 2016, 1, 16013. (8) Aurbach, D.; Lu, Z.; Schechter, A.; Gofer, Y.; Gizbar, H.; Turgeman, R.; Cohen, Y.; Moshkovich, M.; Levi, E. Prototype Systems for Rechargeable Magnesium Batteries. Nature 2000, 407, 724−727. 1823

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824

Article

Chemistry of Materials (29) Shao, Y.; Gu, M.; Li, X.; Nie, Z.; Zuo, P.; Li, G.; Liu, T.; Xiao, J.; Cheng, Y.; Wang, C.; Zhang, J. G.; Liu, J. Highly Reversible Mg Insertion in Nanostructured Bi for Mg Ion Batteries. Nano Lett. 2014, 14, 255−260. (30) Murgia, F.; Stievano, L.; Monconduit, L.; Berthelot, R. Insight into the Electrochemical Behavior of Micrometric Bi and Mg3Bi2 as High Performance Negative Electrodes for Mg Batteries. J. Mater. Chem. A 2015, 3, 16478−16485. (31) Arthur, T. S.; Singh, N.; Matsui, M. Electrodeposited Bi, Sb and Bi 1-xSb X Alloys as Anodes for Mg-Ion Batteries. Electrochem. Commun. 2012, 16, 103−105. (32) Barnes, A. C.; Howells, W. S.; Hull, S.; Keen, D. A.; Hayes, W.; Bamest, A. C.; Guot, C.; Hamilton, M. A. Fast-Ion Conduction and the Structure of P-Mg3Bi2. J. Phys.: Condens. Matter 1994, 6, L467− L471. (33) Kitada, A.; Kang, Y.; Uchimoto, Y.; Murase, K. Electrochemical Reactivity of Magnesium Ions with Sn-Based Binary Alloys (Cu-Sn, Pb-Sn, and In-Sn). ECS Trans. 2014, 58, 75−80. (34) Nadimpalli, S. P. V.; Sethuraman, V. a.; Bucci, G.; Srinivasan, V.; Bower, a. F.; Guduru, P. R. On Plastic Deformation and Fracture in Si Films during Electrochemical Lithiation/Delithiation Cycling. J. Electrochem. Soc. 2013, 160, A1885−A1893. (35) McDowell, M. T.; Lee, S. W.; Nix, W. D.; Cui, Y. 25th Anniversary Article: Understanding the Lithiation of Silicon and Other Alloying Anodes for Lithium-Ion Batteries. Adv. Mater. 2013, 25, 4966−4985. (36) Beaulieu, L. Y.; Hatchard, T. D.; Bonakdarpour, A.; Fleischauer, M. D.; Dahn, J. R. Reaction of Li with Alloy Thin Films Studied by In Situ AFM. J. Electrochem. Soc. 2003, 150, A1457−A1464. (37) Dey, A. N. Electrochemical Alloying of Lithium in Organic Electrolytes. J. Electrochem. Soc. 1971, 118, 1547−1549. (38) Komaba, S.; Yabuuchi, N.; Ozeki, T.; Han, Z. J.; Shimomura, K.; Yui, H.; Katayama, Y.; Miura, T. Comparative Study of Sodium Polyacrylate and Poly(vinylidene Fluoride) as Binders for High Capacity Si-Graphite Composite Negative Electrodes in Li-Ion Batteries. J. Phys. Chem. C 2012, 116, 1380−1389. (39) Rhodes, K.; Dudney, N.; Lara-Curzio, E.; Daniel, C. Understanding the Degradation of Silicon Electrodes for LithiumIon Batteries Using Acoustic Emission. J. Electrochem. Soc. 2010, 157, A1354−A1360. (40) Wang, Z.; Su, Q.; Shi, J.; Deng, H.; Yin, G. Q.; Guan, J.; Wu, M. P.; Zhou, Y. L.; Lou, H. L.; Fu, Y. Q. Comparison of Tetragonal and Cubic Tin as Anode for Mg Ion Batteries. ACS Appl. Mater. Interfaces 2014, 6, 6786−6789. (41) Legrain, F.; Malyi, O. I.; Persson, C.; Manzhos, S. Comparison of Alpha and Beta Tin for Lithium, Sodium, and Magnesium Storage: An Ab Initio Study Including Phonon Contributions. J. Chem. Phys. 2015, 143, 204701−204716. (42) Zhao, K.; Pharr, M.; Wan, Q.; Wang, W. L.; Kaxiras, E.; Vlassak, J. J.; Suo, Z. Concurrent Reaction and Plasticity during Initial Lithiation of Crystalline Silicon in Lithium-Ion Batteries. J. Electrochem. Soc. 2012, 159, A238−A243. (43) Li, J.; Xiao, X.; Yang, F.; Verbrugge, M. W.; Cheng, Y. T. Potentiostatic Intermittent Titration Technique for Electrodes Governed by Diffusion and Interfacial Reaction. J. Phys. Chem. C 2012, 116, 1472−1478. (44) Cheng, Y. T.; Verbrugge, M. W. Evolution of Stress within a Spherical Insertion Electrode Particle under Potentiostatic and Galvanostatic Operation. J. Power Sources 2009, 190, 453−460. (45) Cheng, Y.-T.; Verbrugge, M. W. Diffusion-Induced Stress, Interfacial Charge Transfer, and Criteria for Avoiding Crack Initiation of Electrode Particles. J. Electrochem. Soc. 2010, 157, A508−A516. (46) Totten, G. E., Westbrook, S. R.; Shah, R. J. Fuels and lubricants handbook: technology, properties, performance, and testing; ASTM International: West Conshohocken, PA, 2003. (47) Yang, F.; Li, J. C. M. Lead-Free Electronic Solder; Springer: Boston, MA, 2007.

(48) Liu, X. H.; Fan, F.; Yang, H.; Zhang, S.; Huang, J. Y.; Zhu, T. Self-Limiting Lithiation in Silicon Nanowires. ACS Nano 2013, 7, 1495−1503. (49) Erlebacher, J.; Aziz, M. J.; Karma, A.; Dimitrov, N.; Sieradzki, K. Evolution of Nanoporosity in Dealloying. Nature 2001, 410, 450−453. (50) Erlebacher, J. An Atomistic Description of Dealloying. J. Electrochem. Soc. 2004, 151, C614−C626. (51) Kunduraci, M. Dealloying Technique in the Synthesis of Lithium-Ion Battery Anode Materials. J. Solid State Electrochem. 2016, 20, 2105−2111. (52) Detsi, E.; Van De Schootbrugge, M.; Punzhin, S.; Onck, P. R.; De Hosson, J. T. M. On Tuning the Morphology of Nanoporous Gold. Scr. Mater. 2011, 64, 319−322. (53) Chen, Q.; Sieradzki, K. Spontaneous Evolution of Bicontinuous Nanostructures in Dealloyed Li-Based Systems. Nat. Mater. 2013, 12, 1102−1106. (54) Liu, X. H.; Huang, S.; Picraux, S. T.; Li, J.; Zhu, T.; Huang, J. Y. Reversible Nanopore Formation in Ge Nanowires during LithiationDelithiation Cycling: An In Situ Transmission Electron Microscopy Study. Nano Lett. 2011, 11, 3991−3997. (55) Cook, J. B.; Detsi, E.; Liu, Y.; Liang, Y.-L.; Kim, H.-S.; Petrissans, X.; Dunn, B.; Tolbert, S. H. Nanoporous Tin with a Granular Hierarchical Ligament Morphology as a Highly Stable Li-Ion Battery Anode. ACS Appl. Mater. Interfaces 2017, 9, 293−303. (56) Cook, J. B.; Lin, T. C.; Detsi, E.; Weker, J. N.; Tolbert, S. H. Using X-Ray Microscopy to Understand How Nanoporous Materials Can Be Used to Reduce the Large Volume Change in Alloy Anodes. Nano Lett. 2017, 17, 870−877. (57) Darwiche, A.; Marino, C.; Sougrati, M. T.; Fraisse, B.; Stievano, L.; Monconduit, L. Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism. J. Am. Chem. Soc. 2012, 134, 20805−20811. (58) Chen, L.; Wang, K.; Xie, X.; Xie, J. Effect of Vinylene Carbonate (VC) as Electrolyte Additive on Electrochemical Performance of Si Film Anode for Lithium Ion Batteries. J. Power Sources 2007, 174, 538−543. (59) Choi, N.-S.; Yew, K. H.; Lee, K. Y.; Sung, M.; Kim, H.; Kim, S.S. Effect of Fluoroethylene Carbonate Additive on Interfacial Properties of Silicon Thin-Film Electrode. J. Power Sources 2006, 161, 1254−1259. (60) Etacheri, V.; Haik, O.; Goffer, Y.; Roberts, G. A.; Stefan, I. C.; Fasching, R.; Aurbach, D. Effect of Fluoroethylene Carbonate (FEC) on the Performance and Surface Chemistry of Si-Nanowire Li-Ion Battery Anodes. Langmuir 2012, 28, 965−976. (61) Komaba, S.; Ishikawa, T.; Yabuuchi, N.; Murata, W.; Ito, A.; Ohsawa, Y. Fluorinated Ethylene Carbonate as Electrolyte Additive for Rechargeable Na Batteries. ACS Appl. Mater. Interfaces 2011, 3, 4165− 4168. (62) Aurbach, D.; Gamolsky, K.; Markovsky, B.; Gofer, Y.; Schmidt, M.; Heider, U. On the Use of Vinylene Carbonate (VC) as an Additive to Electrolyte Solutions for Li-Ion Batteries. Electrochim. Acta 2002, 47, 1423−1439. (63) Kumar, H.; Detsi, E.; Abraham, D. P.; Shenoy, V. B. Fundamental Mechanisms of Solvent Decomposition Involved in Solid-Electrolyte Interphase Formation in Sodium Ion Batteries. Chem. Mater. 2016, 28, 8930−8941. (64) Lucas, I. T.; Syzdek, J.; Kostecki, R. Interfacial Processes at Single-Crystal β-Sn Electrodes in Organic Carbonate Electrolytes. Electrochem. Commun. 2011, 13, 1271−1275. (65) Lucas, I. T.; Pollak, E.; Kostecki, R. In Situ AFM Studies of SEI Formation at a Sn Electrode. Electrochem. Commun. 2009, 11, 2157− 2160.

1824

DOI: 10.1021/acs.chemmater.7b04124 Chem. Mater. 2018, 30, 1815−1824