Carbon Lithium-Ion Battery

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Mesoporous Ge/GeO2/Carbon Lithium-Ion Battery Anodes with High Capacity and High Reversibility )

Jongkook Hwang,†,^ Changshin Jo,†,^ Min Gyu Kim,‡ Jinyoung Chun,† Eunho Lim,§ Seongseop Kim,† Sanha Jeong,† Youngsik Kim,*, and Jinwoo Lee*,†,§ †

)

Department of Chemical Engineering, Pohang University of Science and Technology (POSTECH), Pohang, Kyungbuk 790-784, Republic of Korea, ‡Beamline Division, Pohang Accelerator Laboratory, Pohang, Kyungbuk 790-784, Republic of Korea, §School of Environmental Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang, Kyungbuk 790-784, Republic of Korea, and School of Energy and Chemical Engineering, Ulsan National Institute of Science and Technology (UNIST), Ulsan 689-798, Republic of Korea. ^These authors contributed equally.

ABSTRACT We report mesoporous composite materials (m-GeO2, m-GeO2/C, and

m-Ge-GeO2/C) with large pore size which are synthesized by a simple block copolymer directed self-assembly. m-Ge/GeO2/C shows greatly enhanced Coulombic efficiency, high reversible capacity (1631 mA h g
1), and stable cycle life compared with the other mesoporous and bulk GeO2 electrodes. m-Ge/GeO2/C exhibits one of the highest areal capacities (1.65 mA h cm
2) among previously reported Ge- and GeO2-based anodes. The superior electrochemical performance in m-Ge/GeO2/C arises from the highly improved kinetics of conversion reaction due to the synergistic effects of the mesoporous structures and the conductive carbon and metallic Ge. KEYWORDS: germanium oxide . mesoporous materials . block copolymer . catalytic function . lithium-ion batteries

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echargeable lithium-ion batteries (LIB) are the most widely used energy sources for portable electronic devices and electric vehicles. The growing demand for LIBs with enhanced energy and power density has motivated development of anode materials with high specific capacity and good rate capability to replace the conventional graphite anode (372 mA h g
1).1 As potential high capacity anodes, a wide variety of materials such as metal oxides and group-IV elements (e.g., Si, Ge, Sn) have been introduced. Recently, Ge has attracted remarkable attention because of its high theoretical capacity (1600 mA h g
1) and excellent electronic conductivity and ionic diffusivity.2
7 Despite the relatively high price of Ge, these features make Ge-based materials promising for use as anodes in high power LIBs. The oxide form of Ge (GeOx) is also an attractive candidate due to its high theoretical capacity. In principle, GeO2 can deliver the specific capacity high up to 2152 mA h g
1 if it reversibly stores 8.4 Liþ. GeO2 has several advantages over pure Ge such as lower cost, better chemical stability, and better cyclability.8
11 GeO2 undergoes a HWANG ET AL.

conversion reaction with Li (eq 1) to form Ge and Li2O, followed by an alloying reaction of Ge (eq 2) at the first discharging.11
13 GeO2 þ 4Li f Ge þ 2Li2 O

(1)

Ge þ 4:4Li T Li4:4 Ge

(2)

To develop superior GeO2 anodes with high reversible capacity and long cyclability, three problems should be solved: (i) a drastic volume change upon Liþ insertion/ extraction rapidly pulverizes the electrodes and thereby degrades their electrochemical capacity after a few cycles, (ii) low electrical conductivity limits the electron transfer and rate capability, and (iii) the conversion reaction is usually irreversible and restricts the maximum capacity to 1126 mA h g
1 (storing 4.4 Liþ). A substantial amount of Li remains as Li2O, which leads to low initial Coulombic efficiency (ICE). With respect to problems (i) and (ii), considerable improvements have been accomplished by introducing various nanostructures (e.g., nanoparticles,6 nanowires,3 nanotubes,4 and porous structures14) and by making their composites with conductive matrix (carbonaceous materials10,15 or metal VOL. XXX

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* Address correspondence to [email protected] (Y.K.), [email protected] (J.L.). Received for review February 4, 2015 and accepted April 13, 2015. Published online 10.1021/acsnano.5b00817 C XXXX American Chemical Society

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This superior electrochemical performance could be attributed to synergy between mesoporous composite structure and catalytic function of Ge. We used ex situ X-ray diffraction (XRD) and in situ X-ray absorption spectroscopy (XAS) to investigate the phase transition and the changes in local structures of the electrodes and found direct evidence for reversible utilization of Li2O in m-Ge/GeO2/C.

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substrates16,17). However, established synthesis methods usually require complicated, tedious procedures. Furthermore, sometimes the actual mass loading of active material is too low (600 C or prolonged heating time were undesirable because they led to highly aggregated Ge particles embedded in mesostuctures (Supporting Information, Figure S1). Transmission electron microscopy (TEM) images showed well-developed mesoporous structures of m-GeO2/C (Figure 2a) and m-Ge/GeO2/C (Figure 2b). They exhibited short-range-ordered nanochannels or wormhole-like mesostructures. After partial reduction, the parent mesostructures were preserved without structural collapse and aggregation of nanoparticles (Figure 2b and Supporting Information, Figure S2). These structural characteristics were further confirmed by small-angle X-ray scattering (SAXS). All SAXS patterns had a single primary peak with broad and unstructured higher-order peaks, which are typical for wormhole-like or short-range ordered hexagonal structure (Figure 2c).30,31 From the primary scattering VOL. XXX

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Figure 1. (a
d) Schematic representation of synthesis method for m-GeO2/C and m-Ge/GeO2/C.

Figure 2. TEM images of (a) m-GeO2/C and (b) m-Ge/GeO2/C. (c) SAXS patterns of m-GeO2/C and m-Ge/GeO2/C. (d) N2 physisorption isotherms (inset: BJH pore size distribution) of m-GeO2/C and m-Ge/GeO2/C.

vector (q*), the calculated d-spacing was 49.9 nm in m-GeO2/C and 52.4 nm in m-Ge/GeO2/C. Notably, the wormhole-like structure could facilitate wettability and accessibility of electrolytes due to its three-dimensionally interconnected nature.32 The mesoporosity was studied by nitrogen gas adsorption
desorption analysis. The isotherms correspond to type IV curves with a sharp adsorption near 0.95 P/P0, which suggests that uniform large mesopores are dominant (Figure 2d). From the HWANG ET AL.

adsorption branch, pore size was estimated using the Barrett
Joyner
Halenda (BJH) method. Both m-GeO2/C and m-Ge/GeO2/C had narrow pore size distributions with peak pore diameters of 38 and 44 nm, respectively (inset of Figure 2d). The specific surface area of m-GeO2/C and m-Ge/GeO2/C was 77 m2 g
1 and 89 m2 g
1, respectively. CHN elemental analysis estimated the carbon content to be 15.6% in m-GeO2/C and 19.8% in m-Ge/GeO2/C. From the XRD fitting result, the ratio of Ge VOL. XXX

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HWANG ET AL.

analysis was performed at current densities of 0.1 A g
1 (Figure 3a). Details of the charge/discharge profiles and reversibility of the electrodes are discussed in next section. The first discharge capacity of m-GeO2 was 2849 mA h g
1, but was only 1140 mA h g
1 during charging (ICE: 40%). The discharge and charge capacities of m-GeO2/C were 2464 mA h g
1 and 1393 mA h g
1, respectively. The m-GeO2/C had higher ICE value (56.5%) than m-GeO2. These results indicate that the amorphous carbon inside the pore wall could reduce the Liþ loss during SEI formation and enhance the reversibility of conversion reaction. The initial discharge and charge capacities of m-Ge/GeO2/C were 2250 and 1631 mA h g
1, which correspond to ICE of 72.5%. The reversible capacity of m-Ge/GeO2/C was much higher than those in m-GeO2 and m-GeO2/C. Considering that the resol-derived carbon stores ∼600 mA h g
1,21 the capacity contribution of carbon in m-GeO2/C and m-Ge/ GeO2/C is 94 mA h g
1 and 119 mA h g
1, respectively. Mesoporous GeO2 electrodes exhibited more stable cycle performance than the bulk-GeO2 electrode (Figure 3b). The cycle performance was measured at 0.5 A g
1 after three precycles at 0.1 A g
1. The m-GeO2, m-GeO2/C, and m-Ge/GeO2/C electrodes showed stable cycle performance during 90 cycles, with 66.2, 75.0, and 76.1% capacity retention, respectively (capacity retention plot, Supporting Information, Figure S6), but the bulk-GeO2 and bulk-Ge exhibited severe capacity fading after a few cycles, mainly due to a lack of reversibility and to large particles (Supporting Information, Figure S7). Notably, the m-GeO2 electrode gradually stabilized during the first few cycles; this trend indicates that the mesoporous structure improves the cycle life of a pure GeO2 electrode. The large mesopores buffer volume change during the conversion and alloying reactions. In addition, a few nanometers thick pore walls restrain the agglomeration of Li and Ge components during Liþ insertion/extraction.20 The intimate contact among components in m-Ge/GeO2/C could lead to highly reversible Liþ storage by distributing the electrically insulating Li2O in well-blended discharge products. It should be noted that the pristine mesostructures of m-Ge/GeO2/C remain intact with minor structural deformation even after 50 cycles at 0.5 A g
1 (Supporting Information, Figure S8). For comparison, we used the aggregated Ge nanoparticles containing m-Ge/GeO2/C (Supporting Information, Figure S1) as an anode; it showed poor reversibility and poor cycle life (Supporting Information, Figure S9); these results further emphasize the importance of a homogeneous matrix in electrode performance. m-Ge/GeO2/C electrode exhibited much higher rate capability than m-GeO2 and m-GeO2/C (Figure 3c). At 8 A g
1, the m-Ge/GeO2/C delivered 428 mA h g
1. When the current density was changed from 8 to 1 A g
1, the reversible capacity recovered to 1215 mA h g
1, demonstrating the stable Liþ insertion/extraction processes. Electrochemical impedance spectroscopy (EIS) VOL. XXX

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to GeO2 was 11.3:88.7 (Supporting Information, Figure S3). The weight fraction of each component in m-Ge/GeO2/C was Ge 9.1%, GeO2 71.1%, and carbon 19.8%. For comparison, pure m-GeO2 was prepared by calcination of preformed m-GeO2/C in O2 atmosphere for 5 h at 450 C. After removal of carbon to produce m-GeO2, the mesostructure partially collapsed but retained large mesopores (Supporting Information, Figure S4). The specific surface area slightly decreased to 54 m2 g
1 and pore size increased to 45 nm; these changes are attributed to the removal of carbon filled in the pores. These observations indicate that carbon contributes to maintenance of the original nanostructures upon heat treatment. The synthesis method used in this work is simple and straightforward compared to the conventional hard template method that requires complicated, laborious multisteps. Typically, hard templating consists of preparation of hard template, impregnation of precursors, heat treatment, and template etching. In addition, on the contrary to previously reported mesoporous GeO2,26
28 the synthesized m-GeO2/C has high thermal stability and maintains its mesostructures even after carbonization/reduction at high temperature. This stability could be attributed to use of PEO-b-PS and additional carbon precursor (resol). The PS block is mostly decomposed to leave mesopores, but some of the PS is converted to carbon that lines the mesopores.33,34 Resol is mixed with GeO2 sol and PEO block on the nanoscale and is also converted to carbon.22 This characteristic was confirmed by electron energy loss spectroscopy (EELS), which clearly visualized the distributions of carbon, germanium, and oxygen in m-Ge/GeO2/C (Supporting Information, Figure S5). All elements were well-mixed and homogeneously distributed within the mesostructured matrix. The carbons from PS and resol form rigid supports that preserve mesostructures during carbonization and that confine GeO2 nanocrystals inside the walls without aggregation. The carbon coating also significantly increases the electrical conductivity of semiconducting GeO2, forms a protective layer that prevents direct reaction between Ge and electrolyte, and suppresses mechanical fracture during cycling. Previously, hydrothermal treatment,35,36 sol
gel reaction,37,38 and chemical vapor deposition (CVD)5 have been used to introduce carbon layers to encapsulate anodes, but these chemical/ thermal post-treatments often resulted in structural deformation/collapse of pristine nanostructures and sometimes require expensive equipment and toxic chemicals. In contrast, our approach is relatively inexpensive and requires neither sophisticated coating steps nor harsh synthesis conditions. Electrochemical Performance of m-GeO2-Based Anodes. To investigate the anode performance of mesoporous GeO2-based electrodes, galvanostatic charge/discharge

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ARTICLE Figure 3. (a) Galvanostatic charge/discharge profile at 0.1 A g
1 current density, (b) Cycle performance of mesoporous and bulk electrodes and (c) rate capability of m-GeO2, m-GeO2/C, and m-Ge/GeO2/C electrodes. (d) Areal capacity vs gravimetric capacity plot for Ge- and GeO2-based electrodes. Ge NW on stainless steel,16 Ge film on CNT/stainless steel,17 Ge NWgraphene,5 macroporous Ge,14 Ge NP/graphene,39 GeO2/Ge/C NPs,9 Ge NP-carbon,40 Ge NW,41 Ge/C cluster,42 alkanethiol-Ge NWs,43 Ge-graphene,44 GeO2 NP/rGO,11 and GeO2/C core
shell15

data expressed as Nyquist plots represent that the m-Ge/GeO2/C electrode has the smallest charge transfer resistance (semicircle at high frequency region) and total resistance after first discharge and charge (Supporting Information, Figure S10). The resistances increase in the order of m-GeO2 > m-GeO2/C > m-Ge/ GeO2/C. The EIS results also support that the addition of conductive carbon and Ge and well-mixed structure of the components reduced the internal resistance of battery systems, resulting in good cycle and rate performance in m-Ge/GeO2/C electrode. We compared the areal capacity of m-Ge/GeO2/C with reported values (Figure 3d). Recently, thin film electrodes fabricated by direct growth of active materials onto metal substrates have been widely adopted.3,5,16,17,45 This approach allows intimate electrical contact between active species and the current collector, realizing exceptionally high rate capability and cycle performance. However, their performance must be stabilized using thick inactive substrates, so the actual mass loading is 1 mg cm
2 are required to produce sufficient areal capacity (>1 mA h cm
2) for most applications. Because of the limited space in LIB systems, the areal HWANG ET AL.

capacity is one of the most important factors for development of commercial LIB electrodes.46 The areal capacities of m-Ge/GeO2/C and previously reported electrodes were calculated from their maximum capacity values (mA h g
1) and active material loading (mg cm
2) and compared to their gravimetric capacities (Figure 3d). The thin-film electrodes fabricated by vapor
liquid
solid (VLS) growth or CVD showed much lower areal capacities than did the m-Ge/GeO2/C electrode (1.65 mA h cm
2). Moreover, m-Ge/GeO2/C electrode showed the highest areal capacity value among the reported Ge or GeO2 electrodes. These results indicate that mesoporous structure gives high areal and gravimetric capacity, stable cycle life, and reasonable rate capability when used with Ge-based active materials. The Origin of Highly Enhanced Reversible Capacity in m-Ge/ GeO2/C Anode. We conducted cyclic voltammetry (CV) and plotted galvanostatic discharge/charge profiles to investigate the lithium storage behavior in electrodes based on m-GeO2 (Figure 4). The CV profiles of each electrode showed similar tendencies but different behavior in their reversibility, which was greater in m-Ge/GeO2/C than in m-GeO2 and m-GeO2/C (Figure 4a
c). m-GeO2 showed irreversible cathodic peaks at 1.7 and 1.3 V and an intense peak near 0.05 V, which indicate formation of solid electrolyte interphase (SEI), conversion, and alloying reaction of VOL. XXX

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ARTICLE Figure 4. Cyclic voltammetry (CV) graphs of (a) m-GeO2, (b) m-GeO2/C, and (c) m-Ge/GeO2/C electrodes (sweep rate: 0.1 mV s
1). Galvanostatic charge/discharge profiles of (d) m-GeO2, (e) m-GeO2/C, and (f) m-Ge/GeO2/C electrodes. All data were obtained in the potential range of 0.001
3.0 V (vs Li/Liþ).

GeO2, respectively (Figure 4a). During the reverse sweep, the electrode exhibited peaks at 0.4 and 1.2 V, which indicate dealloying and reoxidation of Ge, respectively. However, all cathodic and anodic peaks of m-GeO2 suffered from severe decline during the cycles; this behavior suggests that Liþ insertion/extraction processes are not effective in m-GeO2. In m-GeO2/C, cathodic peaks at 1.0 and 0.7 V were observed for the initial discharge (Figure 4b), and the subsequent peak near 0.1 V was much broader than that of m-GeO2. Anodic peaks were more reversible in m-GeO2/C than in m-GeO2. The redox peaks in m-Ge/GeO2/C were similar to those in m-GeO2/C (Figure 4c). In the anodic scan of m-Ge/GeO2/C, two broad humps were clearly observed near 1.2 and 1.7 V; these can be related to the reoxidation of Ge metal to GeO2. The humps were well maintained in m-Ge/GeO2/C, whereas they gradually faded with cycling and almost disappeared after 10 cycles in m-GeO2 and m-GeO2/C. The corresponding charge/discharge profiles (Figure 4d
f) further support the results from CV experiments (Figure 4a
c). The reversibility was improved in both discharge and charge processes in the order of m-Ge/GeO2/C > m-GeO2/C > m-GeO2. This difference was most noticeable at charge profiles >1.0 V (Figure 4d
f, red circles). The charge capacity contribution above 1.0 V during the first cycle was 400 mA h g
1 for m-GeO2, 620 mA h g
1 for m-GeO2/ C, and 910 mA h g
1 for m-Ge/GeO2/C. The charge capacity of m-GeO2 and m-GeO2/C gradually decreased upon cycling, possibly due to incomplete Li2O decomposition (Figure 4d,e). In contrast, the m-Ge/GeO2/C showed excellent reversibility during five cycles (Figure 4f). Furthermore, m-Ge/GeO2/C showed 95.7% CE during the second cycles, whereas m-GeO2 and HWANG ET AL.

m-GeO2/C attained only 81.5, and 92.6%, respectively. The CV and galvanostatic charge/discharge analyses suggest that the capacity and CE of each electrode are closely related to the reversibility of the electrochemical Ge oxidation that occurs upon charging above 1.0 V. To clarify the origin of these distinct electrochemical performances, we used ex situ XRD and in situ XAS to investigate the phase transformation process and chemical environment change during discharging/charging. Ex situ XRD analyses were conducted at different potentials to investigate the crystal structure change in sol
gel derived m-GeO2 electrodes (Figure 5). To avoid the signal from Cu foil, the electrodes were fabricated by mixing active materials with conductive carbon and polytetrafluoroethylene. The pristine GeO2 and Ge component are assigned to hexagonal-phase GeO2 (JCPDS no.: 85-0473) and cubic phase Ge (JCPDS No.: 04-0545), respectively. The relatively weaker intensity in m-GeO2/C and m-Ge/GeO2/C than in pure m-GeO2 is attributed to the composite structure with amorphous carbon.47 The average crystallite size of m-GeO2/C was 15.7 nm. After carbon removal by calcination, the crystallite size of m-GeO2 increased slightly to 16.5 nm because of sintering. The crystallite size of m-Ge/GeO2/C was 20.0 nm for GeO2 phase and 20.2 nm for Ge phase. The average size was increased during H2/N2 reduction at 600 C. Upon the first discharging (lithiation) of a m-GeO2 electrode, the pristine hexagonal GeO2 pattern broadened and weakened due to the conversion of GeO2 to Ge and Li2O matrix (Figure 5a). The main peak at 26 was maintained until 0.25 V and completely disappeared when discharging was complete. The LixGe and Li2O peaks were not observed, so they are either VOL. XXX

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Figure 5. Ex situ XRD patterns for (a) m-GeO2, (b) m-GeO2/C, and (c) m-Ge/GeO2/C electrodes.

amorphous state or their crystals are too small to be detected by XRD. In m-GeO2/C (Figure 5b) and m-Ge/GeO2/C (Figure 5c) electrodes, the hexagonal GeO2 phase disappeared before discharging to 0.5 V, which is much higher than in the m-GeO2 electrode (near 0.001 V). This accelerated decomposition of GeO2 could be attributed to the conductive carbon layer and the cubic-phase Ge metal, which promote electron transfer and Liþ ion diffusion, thereby facilitating electrochemical reduction. The cubic phase Ge in m-Ge/GeO2/C electrode remained unchanged until 0.25 V, then transformed to amorphous LixGe alloys, not to crystalline LixGe phase, during which the alloying reaction proceeds (Figure 5c). At the end of discharging (0.001 V), all the electrodes lost their original crystalline phases and were transformed to a similar, newly developed crystalline phase that cannot be matched to any known crystalline LixGe alloy (Figure 5). Those phases were sustained even after full charge to 3.0 V. Some of the XRD peaks at 32.7, 36.0, and 51.1 can be assigned to metastable tetragonal ST12 Ge metal phase. The peaks at 38.5 and 44.7 have never been reported in previous Ge-based anode studies and remain unidentified. After three cycles of discharge/charge, these phases maintained their crystalline nature and became dominant, while no other phases were recovered. Previously, similar peaks at 32.6 and 35.8 have been observed during de(lithiation) of bulk Ge, GeO2, and GeS2 electrode but the exact phases could not be identified.48 Formation of metastable tetragonal ST12 Ge phase in GeS2, GeO2, and Ge nanocrystal electrodes has been reported recently.6 This phase became dominant after a few cycles and persisted for 100 cycles. First-principle calculations to explain how this phase was generated and sustained upon cycling indicated that it was produced by Li intercalation into Ge in the form of LixGe12 (x g 3). The authors calculated that the Li-intercalated ST12 LixGe12 structures are thermodynamically more HWANG ET AL.

stable than other possible phases (e.g., cubic Ge). At high Li concentration (x g 8), ST12 LixGe12 phases were transformed to amorphous LixGe alloys, but some portion still preserved its crystalline structures because of high thermodynamic stability. The intercalated Li atoms cannot be completely extracted during delithiation. Therefore, after the Li-intercalated ST12 structures emerged, they persisted during the whole de(lithiation) process, so the unknown peaks at 38.5 and 44.7 might signify a metastable hexagonal Ge metal phase49 with similar formation/maintenance mechanism related to Li intercalated structures. Although the pristine crystal phases were different from each other, all electrode materials were transformed to identical crystalline phases after repeated cycles but had distinct electrochemical behaviors. Thus, it can be inferred that the different amorphous matrices that originate from different pristine crystalline phases significantly affected the electrochemical performance of m-GeO2-based electrodes rather than the commonly formed metastable crystalline Ge phases. Therefore, we used atomic-selective X-ray absorption spectroscopy (XAS) to trace the chemical environment of the amorphous matrix, which cannot be identified by XRD patterns. We obtained the in situ Ge K-edge X-ray absorption near-edge structure (XANES) spectra and corresponding Fourier transforms of extended X-ray absorption fine structures (EXAFS) spectra of m-GeO2 and m-Ge/GeO2/C electrodes in the first cycle (Figure 6). At open circuit voltage (3.0 V), the m-GeO2 electrode had an intense peak near 11110 eV, which corresponds to the typical XANES spectrum of GeO2 with tetravalent state (Figure 6a).19 As discharging (lithiation) proceeded, the intensity of this main peak decreased and a new peak feature developed at lower energy region due to the conversion reaction (a peak at lower energy suggests a reduction of the initial atomic element). Similarly, the Fourier-transformed (FT) amplitude of metallic Ge
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ARTICLE Figure 6. In situ XAS obtained during the first cycle: XANES spectra of m-GeO2 for the first (a) discharge and (b) charge, and (c) corresponding Fourier transform of the EXAFS spectra, XANES spectra of m-Ge/GeO2/C for the first (d) discharge and (e) charge, and (f) corresponding Fourier transform of the EXAFS spectra.

in the radial distribution function (RDF) of EXAFS spectra, while the FT amplitude of Ge
O decreased (Figure 6c). These facts mean that increase in Liþ insertion resulted in formation of Ge metal phase and in further alloying. Substantial quantities of Ge
O bonds remained even after full discharge, suggesting incomplete conversion reaction. During charging (delithiation), both XANES (Figure 6b) and EXAFS (Figure 6c) spectra showed only a minor change and were hardly recovered back to pristine state. The XANES profiles for the second cycle also showed a similar pattern with very small peak shifts (Supporting Information, Figure S11). All results imply that Liþ extraction from the lithiated m-GeO2 electrode is highly irreversible due to Li2O formation on the electrode surface and that this irreversibility prevents reoxidation of metallic Ge to GeO2 during charging. In contrast, the m-Ge/GeO2/C electrode showed distinguishable and reversible XANES profiles (Figure 6d,e), although the tendency of peak features variation was similar to that in the m-GeO2 electrode. The distinct peak features of metallic Ge phase in lower energy region (Figure 6d) were associated with greater generation of metallic Ge phase and the formation of Li
Ge bonding than those in m-GeO2.50 These trends were further confirmed by EXAFS spectra, which showed rapid decrease of the FT peak intensity for Ge
O bonding, HWANG ET AL.

increase of the FT peak intensity for pure metallic Ge
Ge bonding, and new formation of FT peak for Ge
Li in the higher r space (Figure 6f). Upon charging, the Ge metallic-like XANES spectra steadily shift to higher energy state and GeO2-like XANES spectra gradually developed; these changes indicate reversible reformation of the GeO2 phase (Figure 6e). The direct evidence for Ge oxidation was also supported with the RDFs of EXAFS spectra, which showed the clear development of Ge
O bonds during charging (Figure 6f). The efficient phase transition between GeO2 and Ge/Li metallic alloy in m-Ge/GeO2/C electrode can be considered as a structural driving force to enhance their electrochemical properties such as CE and specific capacity. This phenomenon can be understood by the accelerated kinetics of conversion reaction due to coexistence of a conductive carbon network and additional metallic Ge, which could act together as an electron/ion transfer pathway and provide intimate contact with Li2O matrix.18,19 Furthermore, the metallic Ge embedded in the GeO2 matrix could have an important function as an electrocatalyst to promote Li2O decomposition.9 In summary, mesostructure mitigated the stress caused by volume change and extended cycle life of the anode significantly. In addition, the Li2O could be highly dispersed inside the mesoporous conductive matrix (Ge metal, carbon), affording effective VOL. XXX

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CONCLUSIONS We reported a relatively simple and direct method to fabricate mesoporous GeO2-based composite materials

METHODS Synthesis of Mesoporous GeO2 Samples. We synthesized PEO-b-PS (34,600 g mol
1 with 14.5 wt % PEO block and a polydispersity index of 1.21) and employed it as a structure directing agent. Resol was prepared according to the previously reported method.25 In a typical synthesis, 0.20 g of PEO-b-PS was dissolved in 10 g of THF. After 1 h of stirring, 0.14 g of resol, 1.01 g of Ge(OEt)4, and 0.18 g of concentrated HCl (35%) were added to the polymeric solution in sequence. The homogeneous mixture was stirred for 1 h and cast on a glass dish. The solvent was evaporated at 40 C, and then the dish was placed into an oven at 100 C for overnight. The brownish film was recovered and ground into fine powders. For m-GeO2/C, these as-made powders were carbonized at 600 C under Ar atmosphere for 2 h. For m-Ge/GeO2/C, the same as-made samples were carbonized at 600 C under Ar atmosphere for 1 h 40 min, and at the same temperature, the Ar gas was changed to 4% H2/Ar gas for 20 min for partial reduction. After in situ thermal reduction, 4% H2/Ar gas was again changed to Ar gas during cooling process. To obtain m-GeO2, the carbonized m-GeO2/C was calcined at 450 C under O2 for 5 h. The heating rate was 1 C/min. Material Characterizations. The morphology of samples was observed using a TEM (JEM-1011 microscope, Jeol LTD) and SEM (S-4200 field-emission SEM, Hitachi). SAXS patterns were measured using 4C SAXS beamline at the Pohang light source (PLS-II). Nitrogen physisorption was conducted using a Tristar II 3020 instrument at 77 K (Micromeritics Instrument Co.). Powder XRD data were collected using a Bruker D8 Advance X-ray diffractometer (Cu KR radiation). Electrochemical Characterization. The electrochemical test was conducted using a two-electrode coin cell system (CR2032). The electrode was prepared by mixing active materials with a conducting carbon (Super P), a carboxymethyl cellulose, and a poly(acrylic acid) in 70:15:7.5:7.5 weight ratio using distilled water as a solvent. The slurry was coated on Cu foil, followed by drying process for 12 h (110 C in vacuum oven). The electrodes were pressed and cut. The active material loading was controlled in the range of 1
1.35 mg cm
2. The lithium foil was used as both counter and reference electrodes. The electrolyte

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(m-GeO2, m-GeO2/C, and m-Ge/GeO2/C) by using BCP directed coassembly. The mesostructured walls are constructed of small GeO2 (Ge) nanoparticles that are homogeneously mixed with carbon species. The m-Ge/ GeO2/C electrode exhibited greatly increased CE, high reversible capacity, and excellent cyclability compared with the m-GeO2 and m-GeO2/C. Moreover, m-Ge/ GeO2/C showed one of the highest areal capacities among reported electrodes based on Ge or GeO2. The origin of the superior electrochemical performance of m-Ge/GeO2/C was thoroughly studied; it results from synergy between mesoporous structures and the catalytic function of Ge, both of which increased the reversibility of the conversion reaction. We have demonstrated that the designed nanostructures and composition control of the GeO2-based anodes can increase their applicability as high capacity anodes in LIBs. This work emphasizes the importance of initial material composition and its contribution to the formation of an amorphous matrix which considerably affects the electrochemical performance of the conversion anodes. Therefore, this work provides a method to develop various high capacity oxide-based conversion anodes that have high reversibility and long cycle stability.

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decomposition of Li2O. From CV and charge/discharge analyses, we confirmed that the highly improved reversibility of m-Ge/GeO2/C is closely related to the electrochemical reaction (Ge reoxidation) that occurs upon charging at voltages >1.0 V. In the ex situ XRD studies, we observed faster GeO2 decomposition (conversion reaction) in m-Ge/GeO2/C electrodes than in m-GeO2 electrode during the first lithiation. In situ XAS experiments were further conducted to trace the changes in local structures of the amorphous matrix. m-Ge/GeO2/C produced more amorphous Ge metal and less insulating Li2O phase than did m-GeO2. We also confirmed the reversible formation of Ge
O bonding (direct evidence for Ge oxidation) during the charging process in m-Ge/GeO2/C electrode. Therefore, we conclude that the combination of interconnected mesostructures, conductive carbon and additional Ge metal improve the kinetics of the conversion reaction, thereby increasing the CE and achieving high reversible capacity in m-Ge/GeO2/C electrode.

was 1.3 M LiPF6 in mixture of ethylene carbonate/ethyl methyl carbonate (EC/EMC, 3:7 volume ratio) þ 10 wt % fluoroethylene carbonate (Panaxetec Co., Korea). All cells were assembled inside the glovebox. The galvanostatic charge/discharge and cyclic voltammetry were conducted using a WBCS-3000 battery cycler (Wonatech., Korea). For ex situ XRD analysis, the electrode was fabricated by mixing active materials with a polytetrafluoroethylene (PTFE) binder and the Super P (70:15:15 in weight ratio). After charge/discharge, the cells were disassembled inside the glovebox. The electrodes were washed by dimethyl carbonate (DMC) to remove the lithium salt. The EIS analysis was conducted after first discharge/charge cycle at 100 mA g
1. The frequency range was 105 to 10
2 Hz under ac stimulus with 5 mV amplitude using a Reference 600 potentiostat (Gamry Instrument, USA). In Situ XAS Characterization. In situ Ge K-edge XAS for m-GeO2 and m-Ge/GeO2/C during the electrochemical reactions were collected on the BL10C beamline (wide-energy XAFS) at the Pohang light source (PLS-II) with top-up mode operation under a ring current of 300 mA at 3.0 GeV. From the high intensity X-ray photons of the multipole wiggler source, monochromatic X-ray beams could be obtained using a liquidnitrogen-cooled double-crystal monochromator (Bruker ASC) with available in situ exchange in vacuum between a Si (111) and Si (311) crystal pair. The Si (111) crystal pair was used for Ge K-edge XAFS measurements (absorption edge energy, 11103 eV) for better energy resolution. Real-time Ge K-edge X-ray absorption spectroscopic data during discharging and charging processes were recorded for m-GeO2 and m-GeO2/C electrodes assembled in a homemade in situ electrochemical cell with polyimide film windows (Swagelok-type cell). In transmittance mode, the incident and transmitted X-ray photon intensities were detected using N2 gas-filled ionization chambers (IC-SPEC, FMB Oxford). Higher-order harmonic contaminations were eliminated by detuning to reduce the incident X-ray intensity by ∼30%. Energy calibration was simultaneously carried out for each measurement with reference Ge bulk metallic foil placed in front of the third ion chamber. The data reductions of the experimental spectra to normalized XANES and

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Supporting Information Available: TEM image, XRD fitting result, SAXS, N2 phsysorption data of mesoporous GeO2 samples and SEM image of bulk-GeO2 sample. Capacity retention and cycle plots for mesoporous GeO2 electrode, and K-edge XANES spectra of m-GeO2 electrode (during second cycle). This material is available free of charge via the Internet at http://pubs.acs.org. Acknowledgment. This research was supported by Basic Research Program through the National Research Foundation (NRF) funded by the Ministry of Science, ICT and Future Planning (2012R1A2A2A01002879 and 2013R1A1A2074550). This work was further supported by a part of the project titled “Technology Development of Marine Industrial Biomaterials”, funded by the Ministry of Oceans and Fisheries, Korea.

REFERENCES AND NOTES 1. Park, C.-M.; Kim, J.-H.; Kim, H.; Sohn, H.-J. Li-Alloy Based Anode Materials for Li Secondary Batteries. Chem. Soc. Rev. 2010, 39, 3115–3141. 2. Park, M.-H.; Kim, K.; Kim, J.; Cho, J. Flexible Dimensional Control of High-Capacity Li-Ion-Battery Anodes: From 0D Hollow to 3D Porous Germanium Nanoparticle Assemblies. Adv. Mater. 2010, 22, 415–418. 3. Chan, C. K.; Zhang, X. F.; Cui, Y. High Capacity Li Ion Battery Anodes Using Ge Nanowires. Nano Lett. 2008, 8, 307–309. 4. Park, M.-H.; Cho, Y.; Kim, K.; Kim, J.; Liu, M.; Cho, J. Germanium Nanotubes Prepared by Using the Kirkendall Effect as Anodes for High-Rate Lithium Batteries. Angew. Chem., Int. Ed. 2011, 50, 9647–9650. 5. Kim, H.; Son, Y.; Park, C.; Cho, J.; Choi, H. C. Catalyst-Free Direct Growth of a Single to a Few Layers of Graphene on a Germanium Nanowire for the Anode Material of a Lithium Battery. Angew. Chem., Int. Ed. 2013, 52, 5997–6001. 6. Cho, Y. J.; Im, H. S.; Kim, H. S.; Myung, Y.; Back, S. H.; Lim, Y. R.; Jung, C. S.; Jang, D. M.; Park, J.; Cha, E. H.; et al. Tetragonal Phase Germanium Nanocrystals in Lithium Ion Batteries. ACS Nano 2013, 7, 9075–9084. 7. Xue, D.-J.; Xin, S.; Yan, Y.; Jiang, K.-C.; Yin, Y.-X.; Guo, Y.-G.; Wan, L.-J. Improving the Electrode Performance of Ge through Ge@C Core
Shell Nanoparticles and Graphene Networks. J. Am. Chem. Soc. 2012, 134, 2512–2515. 8. Wang, X.-L.; Han, W.-Q.; Chen, H.; Bai, J.; Tyson, T. A.; Yu, X.-Q.; Wang, X.-J.; Yang, X.-Q. Amorphous Hierarchical Porous GeOx as High-Capacity Anodes for Li Ion Batteries with Very Long Cycling Life. J. Am. Chem. Soc. 2011, 133, 20692–20695. 9. Seng, K. H.; Park, M.-h.; Guo, Z. P.; Liu, H. K.; Cho, J. Catalytic Role of Ge in Highly Reversible GeO2/Ge/C Nanocomposite Anode Material for Lithium Batteries. Nano Lett. 2013, 13, 1230–1236. 10. Lv, D.; Gordin, M. L.; Yi, R.; Xu, T.; Song, J.; Jiang, Y. B.; Choi, D.; Wang, D. GeOx/Reduced Graphene Oxide Composite as an Anode for Li-Ion Batteries: Enhanced Capacity via Reversible Utilization of Li2O along with Improved Rate Performance. Adv. Funct. Mater. 2014, 24, 1059–1066. 11. Son, Y.; Park, M.; Son, Y.; Lee, J.-S.; Jang, J.-H.; Kim, Y.; Cho, J. Quantum Confinement and Its Related Effects on the Critical Size of GeO2 Nanoparticles Anodes for Lithium Batteries. Nano Lett. 2014, 14, 1005–1010. 12. Pe~ na, J. S.; Sandu, I.; Joubert, O.; Pascual, F. S.; Areán, C. O.; Brousse, T. Electrochemical Reaction between Lithium and β-quartz GeO2. Electrochem. Solid State Lett. 2004, 7, A278–A281. 13. Feng, J.; Lai, M.; Lu, L. Influence of Grain Size on Lithium Storage Performance of Germanium Oxide Films. Electrochim. Acta 2012, 62, 103–108. 14. Jia, H.; Kloepsch, R.; He, X.; Badillo, J. P.; Gao, P.; Fromm, O.; Placke, T.; Winter, M. Reversible Storage of Lithium in Three-Dimensional Macroporous Germanium. Chem. Mater. 2014, 26, 5683–5688.

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15. Ngo, D. T.; Kalubarme, R. S.; Chourashiya, M. G.; Park, C.-N.; Park, C.-J. Electrochemical Performance of GeO2/C Core Shell Based Electrodes for Li-ion Batteries. Electrochim. Acta 2014, 116, 203–209. 16. Kennedy, T.; Mullane, E.; Geaney, H.; Osiak, M.; O'Dwyer, C.; Ryan, K. M. High-Performance Germanium NanowireBased Lithium-Ion Battery Anodes Extending over 1000 Cycles Through in situ Formation of a Continuous Porous Network. Nano Lett. 2014, 14, 716–723. 17. Wang, X.; Susantyoko, R. A.; Fan, Y.; Sun, L.; Xiao, Q.; Zhang, Q. Vertically Aligned CNT-Supported Thick Ge Films as High-Performance 3D Anodes for Lithium Ion Batteries. Small 2014, 10, 2826–2829. 18. Kang, Y.-M.; Kim, K.-T.; Lee, K.-Y.; Lee, S.-J.; Jung, J.-H.; Lee, J.-Y. Improvement of Initial Coulombic Efficiency of Co3O4 by Ballmilling Using Ni as an Additive. J. Electrochem. Soc. 2003, 150, A1538–A1543. 19. Kim, C. H.; Jung, Y. S.; Lee, K. T.; Ku, J. H.; Oh, S. M. The Role of in Situ Generated Nano-Sized Metal Particles on the Coulombic Efficiency of MGeO3 (M = Cu, Fe, and Co) Electrodes. Electrochim. Acta 2009, 54, 4371–4377. 20. Ye, Y.; Jo, C.; Jeong, I.; Lee, J. Functional Mesoporous Materials for Energy Applications: Solar Cells, Fuel Cells, and Batteries. Nanoscale 2013, 5, 4584–4605. 21. Hwang, J.; Woo, S. H.; Shim, J.; Jo, C.; Lee, K. T.; Lee, J. One-Pot Synthesis of Tin-Embedded Carbon/Silica Nanocomposites for Anode Materials in Lithium-Ion Batteries. ACS Nano 2013, 7, 1036–1044. 22. Jo, C.; Hwang, J.; Song, H.; Dao, A. H.; Kim, Y. T.; Lee, S. H.; Hong, S. W.; Yoon, S.; Lee, J. Block-Copolymer-Assisted One-Pot Synthesis of Ordered Mesoporous WO3
x/Carbon Nanocomposites as High-Rate-Performance Electrodes for Pseudocapacitors. Adv. Funct. Mater. 2013, 23, 3747–3754. 23. Jo, C.; Kim, Y.; Hwang, J.; Shim, J.; Chun, J.; Lee, J. Block Copolymer Directed Ordered Mesostructured TiNb2O7 Multimetallic Oxide Constructed of Nanocrystals as High Power Li-Ion Battery Anodes. Chem. Mater. 2014, 26, 3508–3514. 24. Lim, E.; Kim, H.; Jo, C.; Chun, J.; Ku, K.; Kim, S.; Lee, H. I.; Nam, I.-S.; Yoon, S.; Kang, K.; et al. Advanced Hybrid Supercapacitor Based on a Mesoporous Niobium Pentoxide/ Carbon as High-Performance Anode. ACS Nano 2014, 8, 8968–8978. 25. Hwang, J.; Jo, C.; Hur, K.; Lim, J.; Kim, S.; Lee, J. Direct Access to Hierarchically Porous Inorganic Oxide Materials with Three-Dimensionally Interconnected Networks. J. Am. Chem. Soc. 2014, 136, 16066–16072. 26. Zou, X.; Conradsson, T.; Klingstedt, M.; Dadachov, M. S.; O'Keeffe, M. A Mesoporous Germanium Oxide with Crystalline Pore Walls and its Chiral Derivative. Nature 2005, 437, 716–719. 27. Armatas, G. S.; Kanatzidis, M. G. Mesostructured Germanium with Cubic Pore Symmetry. Nature 2006, 441, 1122– 1125. 28. Sun, D.; Riley, A. E.; Cadby, A. J.; Richman, E. K.; Korlann, S. D.; Tolbert, S. H. Hexagonal Nanoporous Germanium through Surfactant-Driven Self-Assembly of Zintl Clusters. Nature 2006, 441, 1126–1130. 29. Soler-Illia, G. J. d. A.; Sanchez, C.; Lebeau, B.; Patarin, J. Chemical Strategies to Design Textured Materials: from Microporous and Mesoporous Oxides to Nanonetworks and Hierarchical Structures. Chem. Rev. 2002, 102, 4093– 4138. 30. Tanev, P.; Pinnavaia, T. A Neutral Templating Route to Mesoporous Molecular Sieves. Science 1995, 267, 865– 867. 31. Lee, J.; Yoon, S.; Oh, S. M.; Shin, C.-H.; Hyeon, T. Development of a New Mesoporous Carbon Using an HMS Aluminosilicate Template. Adv. Mater. 2000, 12, 359–362. 32. Kang, E.; Jung, Y. S.; Kim, G.-H.; Chun, J.; Wiesner, U.; Dillon, A. C.; Kim, J. K.; Lee, J. Highly Improved Rate Capability for a Lithium-Ion Battery Nano-Li4Ti5O12 Negative Electrode via Carbon-Coated Mesoporous Uniform Pores with a Simple Self-Assembly Method. Adv. Funct. Mater. 2011, 21, 4349– 4357.

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Fourier-transformed radial distribution functions (RDFs) were performed through the standard XAFS procedure. Conflict of Interest: The authors declare no competing financial interest.

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33. Lee, J.; Orilall, M.; Warren, S.; Kamperman, M.; DiSalvo, F.; Wiesner, U. Direct Access to Thermally Stable and Highly Crystalline Mesoporous Transition-Metal Oxides with Uniform Pores. Nature Mater. 2008, 7, 222–228. 34. Hwang, J.; Kim, J.; Ramasamy, E.; Choi, W.; Lee, J. Easy Access to Highly Crystalline Mesoporous Transition-Metal Oxides with Controllable Uniform Large Pores by Using Block Copolymers Synthesized via Atom Transfer Radical Polymerization. Microporous Mesoporous Mater. 2011, 143, 149–156. 35. Ding, S.; Zhang, D.; Wu, H. B.; Zhang, Z.; Lou, X. W. D. Synthesis of Micro-Sized SnO2@Carbon Hollow Spheres with Enhanced Lithium Storage Properties. Nanoscale 2012, 4, 3651–3654. 36. Ding, S.; Wang, Z.; Madhavi, S.; Lou, X. W. D. SBA-15 Derived Carbon-Supported SnO2 Nanowire Arrays with Improved Lithium Storage Capabilities. J. Mater. Chem. 2011, 21, 13860–13864. 37. Wang, J.; Li, W.; Wang, F.; Xia, Y.; Asiri, A. M.; Zhao, D. Controllable Synthesis of SnO2@C Yolk
Shell Nanospheres as a High-Performance Anode Material for Lithium Ion Batteries. Nanoscale 2014, 6, 3217–3222. 38. Jiang, H.; Hu, Y.; Guo, S.; Yan, C.; Lee, P. S.; Li, C. Rational Design of MnO/Carbon Nanopeapods with Internal Void Space for High-Rate and Long-Life Li-Ion Batteries. ACS Nano 2014, 8, 6038–6046. 39. Yuan, F.-W.; Tuan, H.-Y. Scalable Solution-Grown HighGermanium-Nanoparticle-Loading Graphene Nanocomposites as High-Performance Lithium-Ion Battery Electrodes: An Example of a Graphene-Based Platform toward Practical Full-Cell Applications. Chem. Mater. 2014, 26, 2172–2179. 40. Ngo, D. T.; Kalubarme, R. S.; Le, H. T. T.; Fisher, J. G.; Park, C.-N.; Kim, I.-D.; Park, C.-J. Carbon-Interconnected Ge Nanocrystals as an Anode with Ultra-Long-Term Cyclability for Lithium Ion Batteries. Adv. Funct. Mater. 2014, 24, 5291– 5298. 41. Chockla, A. M.; Klavetter, K. C.; Mullins, C. B.; Korgel, B. A. Solution-Grown Germanium Nanowire Anodes for LithiumIon Batteries. ACS Appl. Mater. Interfaces 2012, 4, 4658– 4664. 42. Seng, K. H.; Park, M.-H.; Guo, Z. P.; Liu, H. K.; Cho, J. Self-Assembled Germanium/Carbon Nanostructures as High-Power Anode Material for the Lithium-Ion Battery. Angew. Chem., Int. Ed. 2012, 51, 5657–5661. 43. Yuan, F.-W.; Yang, H.-J.; Tuan, H.-Y. Alkanethiol-Passivated Ge Nanowires as High-Performance Anode Materials for Lithium-Ion Batteries: The Role of Chemical Surface Functionalization. ACS Nano 2012, 6, 9932–9942. 44. Ren, J.-G.; Wu, Q.-H.; Tang, H.; Hong, G.; Zhang, W.; Lee, S.-T. Germanium
Graphene Composite Anode for HighEnergy Lithium Batteries with Long Cycle Life. J. Mater. Chem. A 2013, 1, 1821–1826. 45. Huang, X.; Yu, H.; Chen, J.; Lu, Z.; Yazami, R.; Hng, H. H. Ultrahigh Rate Capabilities of Lithium-Ion Batteries from 3D Ordered Hierarchically Porous Electrodes with Entrapped Active Nanoparticles Configuration. Adv. Mater. 2014, 26, 1296–1303. 46. Bogart, T. D.; Oka, D.; Lu, X.; Gu, M.; Wang, C.; Korgel, B. A. Lithium Ion Battery Peformance of Silicon Nanowires with Carbon Skin. ACS Nano 2013, 8, 915–922. 47. Cao, Q.; Zhang, H.; Wang, G.; Xia, Q.; Wu, Y.; Wu, H. A Novel Carbon-Coated LiCoO2 as Cathode Material for Lithium Ion Battery. Electrochem. Commun. 2007, 9, 1228–1232. 48. Kim, Y.; Hwang, H.; Lawler, K.; Martin, S. W.; Cho, J. Electrochemical Behavior of Ge and GeX2 (X= O, S) Glasses: Improved Reversibility of the Reaction of Li with Ge in a Sulfide Medium. Electrochim. Acta 2008, 53, 5058–5064. 49. Eremenko, V. G. New Crystalline Phase of Germanium. Soviet Physics;Solid State 1975, 17, 1647–1648. 50. Lim, L. Y.; Liu, N.; Cui, Y.; Toney, M. F. Understanding Phase Transformation in Crystalline Ge Anodes for Li-Ion Batteries. Chem. Mater. 2014, 26, 3739–3746.

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