In Situ Dealloying of Bulk Mg2Sn in Mg-Ion Half Cell as an Effective

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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 Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b04124 • Publication Date (Web): 09 Jan 2018 Downloaded from http://pubs.acs.org on January 9, 2018

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Chemistry of Materials

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 PA 19104-6272, USA.

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 characteristic feature size in the sub-100 nm range. These results were further confirmed through electrochemical experiments using a Mg half-cell consisting of bulk Mg2Sn particles with 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

ACS Paragon Plus Environment

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pristine, demagnesiated and re-magnesiated forms. The results suggest that the starting micrometer-sized Mg2Sn particles are converted into nanostructured β-Sn with characteristic sizes ranging from 10‒50 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 alloy-type 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 ever-growing 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 Mgion-blocking passive film on metallic magnesium.12–15 Therefore, only a handful of electrolytes, including 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


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and the relatively limited operating voltage window which makes APC not suitable for highvoltage cathodes.3,20,23 By replacing the 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). On the contrary to Mg3Bi2, Mg2Sn exhibits a much higher theoretical capacity of 903 mAh/g, and the reversible storage of Mg through the half-cell reaction in equation (1) only requires ~0.15 V (vs Mg2+/Mg),24 making Sn an attractive high energy density anode material for MIBs. + 2 + 4 ̅ ↔

(1)

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 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 Sn-based 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


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compared to the pure Sn in cyclic voltammetry tests, while Cu-Sn intermetallics (Cu6Sn5 and Cu3Sn) are essentially inactive. Furthermore, Parent et al. investigated 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 Sb-rich magnesiated sub-domains with 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 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 high-performance nano-particulate 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 C-rates (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 alloytype 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 characteristic size in the range of 10-50 nm. The critical feature of our approach is the chemo-mechanical 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


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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 Discussions 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 the Results and Discussions section, various aspects of phase, compositional and microstructural features of the dealloyed and re-alloyed 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 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 active electrode material instead of pure Sn.


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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 like Nudged Elastic Band (NEB).40,41 The results of such calculations do not necessarily reflect the real physics involving the alloying reaction of βSn 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 rate determining 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 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


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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 alloy-type reaction associated with the (de)magnesiation of Sn 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 growing demagnesiated shell of Sn (grey) 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 = where 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 1(a). The elements at the surface of a demagnesiating Mg2Sn particle accommodate


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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 of ~800 MPa. Once we include the perfectly plastic behavior in our simulations, the Sn shows tensile plastic yielding after von Mises stress reaches to the yield stress of Sn. The stress distribution for the same demagnesiation radius but with plasticity for Sn is shown in the Figure 1(b).

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.


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It is evident from these simulations that the hoop stress near the surface remain tensile but maximum stress is much smaller compared to perfectly elastic deformations considered in previous case. The critical size (dcr) for the fracture to be energetically favorable in plastic materials is given by:42 = Γ/

(2)

Where Γ is the fracture energy of the material, E is the elastic modulus, 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 , 46 and 1.0 for Z.47 Substituting these values in Equation 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 1(c)). 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 Mgion battery anodes starting from micrometer-sized Mg2Sn particles as active electrode materials


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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 nanostructured can then be cycled subsequently. The advantage of such a top-down nanosynthesis starting from bulk Mg2Sn is at least two-fold: (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 βSnSb electrode materials.25,26 In the following sections, we present the characterizations and electrochemical performances of Mg2Sn.

Crystallography and Microstructural Studies Figure 2(a) and 2(b) show the XRD patterns of the as-synthesized Mg2Sn 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.


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Figure 2. (a) XRD pattern of as-prepared Mg2Sn electrode (i); the 1st demagnesiated (ii); and the 1st magnesiated (iii) states. b) The corresponding voltage-capacity profile. Inset in panel (a) shows a magnified view of the high angle region. The as-prepared material is composed mainly of cubic Mg2Sn as shown by the XRD pattern (i) in Figure 2(a) (ICSD col. code: 151368). Following charging the electrode to an upper cut-off potential of 0.6 V (see Figure 2(b) point (ii)), the material dealloys effectively into β-Sn as depicted by the XRD pattern (ii) in Figure 2(a) (ICSD col. code: 52269). A subsequent reduction down to a lower cut-off potential of 5 mV vs Mg2+/Mg (see Figure 2(b) point (iii)) causes the diffraction peaks of Mg2Sn to grow back again, with some Sn left unreacted (see XRD pattern (iii) in Figure 2(a)). 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 2(a)i) and electrochemically grown Mg2Sn (Figure 2(a)iii) could


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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 2(a) 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; (ii) significant particle size reduction occurs during the first cycle. 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 3(a) 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 3(b) 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 micro- and sub-micrometer-sized pure Sn at similar and slower rates were unsuccessful. Singh et al. have shown that full magnesiation of pure Sn nanoparticles with characteristic size of ~150 nm was achieved at the slow rate of C/500.24


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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 re-magnesiated 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.

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 3(c) suggests that re-magnesiation of the in situ formed Sn creates aggregates of 1 µm or larger. However, it is essential to discriminate between aggregates and absolute particle size. For example, the EDX


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map in Figure 3(c) shows that the elemental distribution of Sn and Mg over large aggregates is very uniform and chemically indistinguishable (cf. Figure 3(b)), 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 re-magnesiate sub-100 micrometer-sized Sn particles (Figure 3(b)) formed in situ to Mg2Sn (Figure 3(c)) may initially appear counter intuitive and selfconflicting 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, 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 nanometer. 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 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. The SEM has been used


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merely as a tool to adjudicate the maximum possible particle sizes, spatial elemental distribution and degree of particles agglomeration.

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).


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Figure 5. TEM micrographs of the demagnesiated (a-c); and re-magnesiated electrode (e-g). Panels (d) and (h) shows the Fourier transform of the enclosed areas of panel (c) and (g), respectively. 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 re-magnesiated specimens. The morphology of the in situ formed Sn after the first demagnesiation (Figures 5(ac)) depicts pulverized Sn particles with diameters of maximum around 50 nm. Comparing with SEM results (Figures 3(b) 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 5(e-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


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first demagnesiation. The Fourier transformation of high resolution images containing lattice fringes unequivocally identifies Sn and Sn-Mg2Sn mixture for the demagnesiated and remagnesiated samples, respectively, a result also observed in the XRD patterns (Figure 2(a)-(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 re-used 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 V and 0.15 V vs Mg2+/Mg, respectively, except for cycle #1 where they take place at average potentials of 0.38 V 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).


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

Note in Figure 2(b) 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 10fold 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 2(a) inset, 3(c), 4 and 5(e-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 sub-section. It should be pointed out that the capacity associated with the first magnesiation represents only ~60% of the initial demagnesiation columbic charge. Such an irreversible behavior was observed to be rate-independent, meaning that even cells cycled at slower rates (like C/50 from Figure


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2(b)) 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 7(a). 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 7(b) shows the collected EIS data acquired by charging the cells to 0.6 V vs Mg2+/Mg (demagnesiation of Mg2Sn to β-Sn with an state-of-charge (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

Figure 7. The 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.

from Figure 6(a), the first magnesiation


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requires a higher polarization overpotential than subsequent cycles. It can be seen from the red trace curve in Figure 7(b) that the corresponding impedance |Z| is reduced from 3.22 kΩ down to a minimum of 920 Ω at the end of 6th cycle (blue trace) with an eventual constant growth till the 30th cycle (3.20 kΩ, yellow trace). Interestingly, after the 6th cycle while the length of the low frequency tail is increased, the medium-high frequency semicircle remains constant. It 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 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 6th 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.


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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, an 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 has 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 that such additives are missing or are under investigations 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 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 undergoing 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+ ions diffusion out into the electrolyte. Our argument is supported by the trend of the voltage-composition profiles (Figure


21


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S5) of Mg2sn electrode cycled at C/5; following the 2nd 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 high-potential 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

oxidative pulses

to

destroy the passive film) a gradual

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.

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 5 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 grows 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


22

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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 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 are showing fast cycling of Sn for an extended number of cycles starting initially from large micron sized particles, utilizing only Sn and Mg and no sacrificial element. In general, the results presented here suggest that Sn may be regarded not only as 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 alloytype anodes. Note that although APC was used in the present work, the ultimate goal of alloytype 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

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


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~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: (i) The in situ formation of nanostructured Sn from Mg2Sn powder prepared by a rudimentary milling process; (ii) the good performance of the synthesized nanostructured Sn as Mg-ion battery anode, (iii) the possibility to control passive films 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 the Young’s modulus E= 61.9 GPa and Poisson’s ration ν=0.36. For Sn E= 42 GPa, ν=0.24 and σY = 800 MPa .


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Reagents Sn shots (Sigma-Aldrich, 99.8%), Mg foil (Sigma-Aldrich 99.9%) and chip (99.98% SigmaAldrich), 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, 2M in THF, Sigma-Aldrich), aluminum chloride (AlCl3, anhydrous, 99.99%, Sigma-Aldrich), Cu foil (MTI) were used as received. Tetrahydrofuran (THF, >99.9% anhydrous, Sigma-Aldrich) has been further dehydrated by treating with molecular sieve (4Å, Fisher) prior to use. Synthesis About 3.9 g Mg2Sn was 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.


25


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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 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 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 5:1:1:1:2 ratios and NMP has been added to the mixture to dissolve the PVDF and form a uniform slurry, which has been casted on 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 has been carried out under inert atmosphere in an argon-filled glove box (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 with Mg metal as the counter and reference electrode, binder-free glass microfiber as separator and APC electrolyte. The APC electrolyte has been prepared following the reported method by reacting PhMgCl and AlCl3 in THF at 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


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potentiostat/galvanostat/EIS. Samples extracted from coin cells have been washed with dry THF in an argon-filled glove box and allowed to dry before taking them out of the glove box for analysis by XRD and electron microscopy. Supporting Information The following items are available free of charge as the supporting information: 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 * Correspondence should be addressed to E.D. ([email protected])

ACKNOWLEDGMENT. The authors are thankful to Penn Engineering for the financial support through the PI startup.

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In Situ Dealloying of Bulk Mg2Sn in Mg-Ion Half Cell as an Effective

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