Behavior of Germanium and Silicon Nanowire ... - ACS Publications

Behavior of Germanium and Silicon Nanowire Anodes with Ionic Liquid Electrolytes Guk-Tae Kim,†,‡ Tadhg Kennedy,§ Michael Brandon,§ Hugh Geaney,§ Kevin M. Ryan,§ Stefano Passerini,†,‡ and Giovanni B. Appetecchi*,∥ †

Helmholtz Institute Ulm, Karlsruhe Institute of Technology, Helmholtzstrasse 11, 89081 Ulm, Germany Karlsruhe Institute of Technology, P.O. Box 3640, 76021 Karlsruhe, Germany § Materials and Surface Science Institute and the Department of Chemical and Environmental Sciences, University of Limerick, V94 T9PX Limerick, Ireland ∥ ENEA, Italian National Agency for New Technology, Energy and Sustainable Economic Development, Materials and Physicochemical Processes Laboratory, Via Anguillarese 301, 00123 Rome, Italy ‡

S Supporting Information *

ABSTRACT: The electrochemical behavior of binder-free, germanium and silicon nanowires as high-capacity anode materials for lithium-ion battery systems is investigated in an ionic liquid electrolyte. Cyclic voltammetry, cycling tests, and impedance spectroscopy reveal a highly reversible lithium alloying/dealloying process, as well as promising compatibility between the Ge and Si materials and the electrolyte components. Reversible capacities of 1400 and 2200 mA h g−1 are delivered by the Ge and Si anodes, respectively, matching the values exhibited in conventional organic solutions. Furthermore, impressive extended cycling performance is obtained in comparison to previous research on Li alloying anodes in ionic liquids, with capacity retention overcoming 50% for Si after 500 cycles and 67% for Ge after 1000 cycles, at a current rate of 0.5C. This stable long-term cycling arises due to the ability of the electrolyte formulation to promote the transformation of the nanowires into durable porous network structures of Ge or Si nanoligaments, which can withstand the extreme volume changes associated with lithiation/delithiation. Remarkable capacity is exhibited also by composite Ge and Si nanowire electrodes. Preliminary tests with lithium cobalt oxide cathodes clearly demonstrate the feasibility of Ge and Si nanowires in full batteries. KEYWORDS: germanium anodes, silicon anodes, nanowires, ionic liquid electrolytes, lithium batteries

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issues arising from the large degree of expansion/contraction (>300% in the case of Si), accompanying the lithium alloying− dealloying processes.7 The main problems associated with such volume changes are pulverization of the active material, delamination from the current collector, and an unstable solid electrolyte interphase (SEI) layer, which causes continuous electrolyte consumption and unsatisfactory Coulombic efficiencies.8−10 Various material designs have been proposed to alleviate these issues. One of the most noteworthy, introduced in 2007, was an anode architecture based on Si11 or Ge12 nanowires (NWs), grown directly onto current collectors by chemical vapor deposition (CVD). The impressive capacity retention displayed by such anodes was attributed to a combination of

echargeable lithium batteries have become the energy storage system of choice for many technologies owing to their high gravimetric and volumetric energy densities and are therefore likely to remain important for the foreseeable future.1 However, applications such as automotive, electrical storage from renewable sources, smart grids, and consumer electronics are placing ever increasing demands on lithium-ion batteries (LIBs), with respect to charge capacity and energy and power densities. To meet these requirements, considerable research attention has been directed toward highcapacity anode and cathode materials for LIBs. On the anode side, silicon and germanium Li-alloying materials are considered very promising candidates for replacing the currently employed graphite.1−5 In terms of theoretical lithium storage capacity, silicon possesses the highest known value of 3579 mA h g−1, while germanium offers an impressive 1384 mA h g−1, compared to just 372 mA h g−1 for graphite.6 Despite these extremely large capacity figures, commercialization of Si- or Gebased anodes has been frustrated by performance stability © 2017 American Chemical Society

Received: March 10, 2017 Accepted: May 22, 2017 Published: May 22, 2017 5933

DOI: 10.1021/acsnano.7b01705 ACS Nano 2017, 11, 5933−5943

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organic electrolytes, were obtained for both types of anode. We report outstanding extended cycling performance for a Si electrode in an ionic liquid, with a capacity retention of approximately 1000 mA h g−1 after 300 cycles, at a high rate of 0.5C. Ge anodes exhibited even more remarkable long-term cycling ability, by retaining 67% of their initial capacity over 1000 cycles, also at 0.5C. The rate capability of both Si and Ge NW anodes in the IL electrolyte was also noteworthy, with capacities well above those of graphite possible even at a 3C rate. Additionally, the crucially important transformation of Si or Ge NWs into entangled networks of nanoligaments was demonstrated in the IL electrolyte, just as for carbonate-based solutions. Finally, the performance of Ge and Si NWs (in IL electrolytes) was preliminarily tested in composite electrodes and full cells with commercial lithium cobalt oxide (LCO) cathode tapes. The results are discussed in the present paper.

the resistance offered by the one-dimensional nanoscale morphology to elastic strain and the availability of space between adjacent NWs to facilitate expansion.8,11 Since the publication of these works, there has been intense research interest in the further development and improvement of NWbased anodes, with the area being the subject of reviews by ourselves13 and others.4,14−16 We have developed a method for fabricating Si and Ge NW anodes without the requirement for CVD equipment.17,18 Ge NW anodes produced by our technique offered capacities on the order of 900 mA h g−1 after 1900 cycles (1C rate), while also offering exceptional rate capability (745 mA h g−1 at 100C delithiation rate).19 Importantly, we also observed for both Ge NWs20 and Ge− Si heterostructured NWs21 that conversion, with cycling, of the as-grown NWs to a porous network of interconnected nanoligaments is essential for the retention of high capacities over the course of many hundreds of cycles. The success of this morphological transformation depends on a few factors including electrolyte composition. Appealing performance was observed also in mesoporous amorphous silicon electrodes,22 which were seen able to deliver more than 1000 mA h g−1 after 700 cycles. Ionic liquids (ILs), molten salts at room temperature constituted by organic cations and inorganic/organic anions, represent a very interesting class of room-temperature fluids.23−25 It is worth noting that the physicochemical properties of ILs can be finely tuned (through even slight modification of the IL structure and/or introduction of functional groups) to match particular operating conditions. In addition, properly combined IL mixtures show improved properties, often not exhibited by single ionic liquid materials. In the past few years, ILs have been widely investigated as safe electrolyte components to replace the volatile and hazardous alkyl carbonates commonly used in commercial lithium batteries. Particularly, ILs based on N-alkyl-N-methylpyrrolidinium cations, (PYR1A)+ (the subscripts indicate the number of carbon atoms in the alkyl side chains), and bis(trifluoromethanesulfonyl)imide, (TFSI) − , and bis(fluorosulfonyl)imide, (FSI)−, anions, have shown favorable properties.26 Not surprisingly, there has been some convergence in research on these emerging classes of anode materials and electrolytes, with several studies on the performance of both pure Si27−30 and Si composite31−35 electrodes in various IL electrolytes. While some impressive capacity values have been achieved, many of these reports are limited to 100 cycles or less.30,31,33 Where more extended cycling has been detailed, the charge/discharge rates have been moderate or low (≤0.2C rate).34,35 We are not aware of any study on Ge-based anodes in IL electrolytes. In this paper, we focus on the electrochemical performance of directly grown, binder-free Si and Ge NW anodes in an electrolyte formed by suitably combining36 the PYR13TFSI and PYR13FSI IL materials with the LiTFSI lithium salt. For instance, a comparison between Si, i.e., considered as one of the most, if not the most, promising anode material (even because of its wide availability and low cost) for the next generation of Li-ion batteries, and Ge, i.e., addressed as an advanced anodic material despite its cost, is presented. At present, germanium is rather expensive for large-scale applications. However, one of the purposes of the present work is to compare Ge as an advanced anode material from a scientific point of view. It was found that capacities, like those achieved in conventional

RESULTS AND DISCUSSION The synthesized PYR13TFSI and PYR13FSI ionic liquids exhibited moisture, Li+, and Br− content below 2 ppm as determined by Karl Fischer titration, atomic absorbance, and EDX measurements, respectively. In designing the quaternary electrolyte, 0.1LiTFSI− 0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC, we aimed to favorably combine the moderate viscosity, fast ion transport properties, and low melting point of PYR13FSI with the wide electrochemical and thermal stability of PYR13TFSI.26 A previous study indicated that, in terms of compromise between thermal, ion-transport, and rheological properties, the optimal TFSI:FSI anion mole ratio is 2:3,36 whereas the LiTFSI mole fraction was kept equal to 0.1.37 EC (ethylene carbonate) was added because, in combination with the FSI anion, it promotes the formation of a stable, protective SEI onto anodes.38 The presence of this combustible solvent in a rather small amount (organic additive) is not detrimental for the safety of battery systems because it is mostly consumed during the first halfcycle.39 The quaternary electrolyte shows ionic conductivity approaching 10−3 S cm−1 at −10 °C and an electrochemical stability window up to 5 V.40,41 In previous works we demonstrated the possibility of successfully preparing directly grown germanium and silicon nanowire anodes exhibiting large capacities and very good cycling behavior in conventional Li-ion electrolytes.17−21 Motivated by a desire to see if this performance can be replicated in safer, IL-based electrolytes, Cu3Ge-seeded Ge NW anodes and Sn-seeded Si NW anodes were fabricated using the same methodology as in our earlier studies.17,19 Representative scanning electron microscopy (SEM) images of the as-prepared Ge and Si NWs are presented in Figure 1A and B, respectively. It is evident from Figure 1A that the Ge NW anodes are composed of a mixture of tortuous (average diameter 74 nm) and much straighter (average diameter 39 nm) wires. By contrast the Si NWs (Figure 1B) are generally quite straight with an average diameter of 100 ± 30 nm, although some wires with a slightly larger diameter of approximately 200 nm are also apparent. The NWs are crystalline, as previously described in detail for both the Cu3Ge-seeded Ge19 and Sn-seeded Si17 varieties. The reversibility of the Li+ storage process for Ge and Si anodes in ionic liquid electrolyte was initially investigated by cyclic voltammetry (Figure 2). The Ge CVs (Figure 2A) demonstrate very similar features to differential capacity plots that we reported and discussed for the same type of NW anode 5934


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prevalent after the initial cycle, as evidenced by the increasing presence of a broad anodic peak centered at 0.6 V on the second and later delithiation sweeps. The latter feature is associated with the removal of Li+ from a-Li15Ge4.20 The CVs for the Si NWs (Figure 2B) also replicate the features of differential capacity plots recorded for this material in organic-carbonate electrolyte.17 The evolution of these CVs with cycling can be explained in a manner analogous to the Ge data. The sharp increase in the magnitude of the cathodic current on the first cycle below 0.1 V arises from the lithiation of c-Si. The appearance of an additional cathodic peak at 0.15 V on the second and subsequent cycles is symptomatic of the progressive lithiation of the now amorphized Si to form a mixture of a-Li15Si4 and c-Li15Si4 at the lower potential limit, while the anodic peaks at 0.35 and 0.55 V arise from the reverse of these processes.14,44−47 The cathodic peaks at 0.65 V (which is much diminished after the first cycle) and in the region of 0.40 V coincide with features that we showed to be due to lithiation of the Sn seeds in conventional electrolyte.17 Similarly, the anodic peak at 0.65 V is associated with the dealloying of Li+ from Sn; this implies that some of the seed material remains electrically contacted to the Si NWs over the first five cycles at least. For both Ge and Si NW anodes, it is worth noting the reversibility of the lithium alloying process, highlighted by the correspondence between the cathodic and anodic features and the reproducibility of the cyclic voltammetry traces. Impedance measurements carried out on the Ge and Si NW electrodes before and after the cyclic voltammetry measurements are presented and discussed in the Supporting Information (Figure S1). Significantly, no change of the bulk electrolyte resistance (determined as the highfrequency intercept of the Nyquist diagram with the Z′ axis)48 was observed following the CV experiments, suggesting the good electrochemical stability of the ionic liquid-based electrolyte. The voltage vs capacity profiles recorded during galvanostatic cycling tests are reported in Figures 3 and 4. The first discharge (lithiation) (0.02C) behavior (Figure 3) is seen to differ somewhat compared to the following ones. The shoulder observed for both electrodes around 1.6 V fully disappears from the second cycle, and so it may be attributed to SEI layer formation. This feature is followed by two plateaus at 0.35 and 0.2 V for the Ge electrode (Figure 3A) and one flat plateau around 0.15 V (Figure 3B) for the Si anode. In the following discharge half-cycles (run at 0.1C) three plateaus around 0.5, 0.35, and 0.2 V are seen for the Ge electrode, whereas only two plateaus located at 0.25 and 0.1 V are found for the Si one. These results are in good agreement with those obtained by cyclic voltammetry (Figure 2). As discussed for the CVs, the differences between the first and subsequent discharges arise from the fact that Li+ is alloying with crystalline Ge or Si in the initial cycle and thereafter with amorphous phases of these elements. The increase of the current rate from 0.1 to 3C (Figure 4) did not result in a significant change of the voltage profile shape (especially for the Si electrode), but did lead to a substantial reduction of the delivered capacity, particularly at 3C rate. Also, it is worth mentioning that no dramatic rise in ohmic drop with increasing current densities occurred for either the Ge or Si anodes, suggesting moderate diffusive phenomena within the IL electrolyte even at high rates. The electrochemical performance of the Ge and Si electrodes at different current rates is plotted in Figure 5, as capacity

Figure 1. Representative SEM images of the as-prepared nanowire− Cu3Ge-seeded Ge nanowires (A) and Sn-seeded Si nanowires (B).

Figure 2. Cyclic voltammetries run (20 °C) on Li/Ge (A) and Li/Si (B) half-cells in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte. Scan rate: 0.1 mV s−1.

in conventional LiPF6−alkyl carbonate electrolyte.19 A series of peaks between 1.5 and 0.7 V on the first cathodic scan, which do not reappear on subsequent cycles, can be attributed to SEI layer formation. The sharp peak at 0.2 V corresponds to the lithiation of crystalline germanium (c-Ge) and according to the anodic peak around 0.65 V is due to dealloying of Li+ from the crystalline c-Li15Ge4 phase. The absence of the 0.2 V cathodic peak upon further cycling suggests that the active material has become entirely amorphous upon the first delithiation. In its place are observed peaks at approximately 0.53, 0.33, and 0.13 V, which are ascribed to the progressive lithiation of the now amorphous germanium (a-Ge) to a series of increasingly Li-rich phases (a-Ge → a-LixGe → a-Li15Ge4 → c-Li15Ge4).42,43 Formation of the crystalline alloy c-Li15Ge4 becomes less 5935


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Figure 3. Voltage vs capacity profiles of Li/Ge (A) and Li/Si (B) half-cells in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte. The cycling tests were run at 0.1C and 20 °C, whereas the first full cycle was carried out at 0.02C.

Figure 5. Evolution vs cycle number of the discharge/charge capacity (A) and Coulombic efficiency (B) at different current rates, and reversible capacity vs current density dependence (C) of Li/Ge and Li/Si half-cells in 0.1LiTFSI−0.3PYR 13 TFSI− 0.6PYR13FSI + 5 wt % EC electrolyte. T = 20 °C. The current rates are also indicated for comparison purpose.

(Figure 5C). At 0.1C rate the electrodes delivered very large reversible capacities (Figure 5A), i.e., about 1400 and 2200 mA h g−1 for Ge (blue markers) and Si (red markers), respectively. Notably, these values match very closely the performance that we have previously reported, at this rate, for Ge NW19 and Si NW21 anodes in a standard electrolyte of 1 M LiPF6 in ethylene carbonate/diethyl carbonate (1:1 v/v) with 3 wt % vinylene carbonate additive. Relatively low initial discharge/charge efficiencies were observed, with values of 65% for Ge and 61% for Si (Figure 5B). It must, however, be pointed out that low first-cycle Coulombic efficiencies are typical for binder-free Si or Ge NW anodes in regular commercial electrolytes, with the best reported figures in the range of 60−80%.11−13 This arises primarily due to the high surface area-to-volume ratio of such nanoscale electrodes and the inherently large interfacial contact area between the electrolyte and active material, which facilitates extensive SEI layer formation. Importantly, Figure 5B also shows that Coulombic efficiencies approaching or

Figure 4. Voltage vs capacity profiles (20 °C) of Li/Ge (A) and Li/ Si (B) half-cells in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte at different current rates.

(Figure 5A) and Coulombic efficiency (Figure 5B) against cycle number, and capacity vs discharge/charge current density 5936


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after the galvanostatic testing revealed no change in the highfrequency intercept with the Z′ axis, again pointing to the good compatibility of the ionic liquid-based electrolyte with both Ge and Si electrodes (Figure S2). The performance of Ge and Si nanowire anodes was also investigated in prolonged cycling tests conducted at 0.5 C rate. The results are depicted in Figure 6 as voltage vs capacity

exceeding 99% are registered after a few cycles (especially for Si anodes). This indicates that a stable SEI layer forms during the initial cycles in the quaternary electrolyte and renders the active surface substantially inert to further electrolyte decomposition, despite the extreme volume changes experienced by the underlying material during discharge/charge. The Si electrode showed progressively increasing capacities during the initial cycles, most likely due to incomplete lithiation, causing the temporary formation of a c-Si−a-Li15Si4 core−shell structure. As cycling progresses, the crystalline core decreases in diameter upon each successive discharge, resulting in a gradual increase in capacity, as the entirety of the Si material becomes active in the Li+ alloying reaction.14,49 Increasing the discharge/charge current rate (Figure 5C and Table 1) results in a nondramatic capacity decay for both Table 1. Summary of the Discharge/Charge Capacity Values of Ge and Si Anodes in 0.1LiTFSI−0.3PYR13TFSI− 0.6PYR13FSI + 5 wt % EC Electrolyte at Different Current Ratesa reversible capacity/mA h g−1 Ge (IL) Si (IL) Ge (carb)b Si (carb)c

0.01C

0.2C

0.5C

1C

3C

5C

1385 2246 1318 2091

1267 2123 1277 1937

1136 1746 1210 1630

949 1218 1177 1317

455 473 n.a. n.a.

n.a. n.a. n.a. 526

T = 20 °C. Also included for comparison are capacity values from our previous works on the same types of anode in conventional carbonatebased electrolytes (carb), namely, Ge NWs in 1 M LiPF6 in ethylene carbonate/dimethyl carbonate (1:1 v/v) + 3 wt % vinylene carbonate,19 and Si NWs in the same electrolyte except diethyl carbonate substituted for dimethyl carbonate.21 bFrom ref 19. cFrom ref 21. a

anodes up to 1C rate, at which the Ge and Si electrodes retain 69% and 54% of the values exhibited at 0.1C, respectively. A more marked decay in capacity is observed at higher rates; however relevant values (>450 (Ge) and 470 (Si) mA h g−1) are still displayed at 3C (>1 mA cm−2). To put these values in context, rate capability data have been added to Table 1, from our previous studies of Ge NW19 and Si NW21 anodes in conventional electrolytes. The capacity values are similar for both the IL and alkyl carbonate-based electrolytes at moderate rates; in fact higher capacity is observed for Si in the IL, compared to the conventional electrolyte, between 0.1C and 0.5C. At 1C and higher, the performance for the IL begins to lag the commercial electrolyte. This observation may be ascribed to slower ionic diffusion in the more viscous IL compared to the carbonate solvent. However, it is noteworthy that the more limited diffusion rate in the IL becomes apparent only at relatively high current rates (≥1C). From Figure 5A and C it is evident that while Si offers superior capacity to Ge at low-medium current rates, the gap narrows with increasing C rate, until both types of anode exhibit approximately equal performance at 3C. This is a manifestation of the fact that while Si possesses a higher theoretical charge storage capacity than Ge, the latter is characterized by much higher Li diffusivity (× ∼400) and electrical conductivity (4 orders of magnitude greater than Si).50 These properties begin to become as important as intrinsic capacity when the rates of Li+ input and extraction are increased. Impedance measurements conducted before and

Figure 6. Selected voltage vs capacity profiles of Li/Ge (A) and Li/ Si (B) half-cells in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte at 0.5C and 20 °C. (C) Cycling performance for Li/Ge and Li/Si. The first full cycle was carried out at 0.02C.

profiles and summarized in a capacity vs cycle number plot (Figure 6C). The first discharge/charge step (run at 0.02C) displays reversible capacities of approximately 1200 and 2200 mA h g−1 with Coulombic efficiencies of about 60% and 57% for Ge (Figure 6A) and Si (Figure 6B), respectively. No meaningful change of the voltage profile, or drastic increase in the ohmic drop is observed during the extended cycling tests (especially for Ge anodes), despite a progressive increase of the (de)alloying plateau slope, which arises in conjunction with capacity decay. The data presented in Figure 6C indicate that the Ge and Si NW anodes can continue to deliver very large capacities over the course of hundreds of cycles at 100% depth of discharge (DoD) and a relatively high current rate of 0.5 C. For both electrode materials, reversible capacities of over 1000 5937


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ACS Nano mA h g−1 were observed after 300 consecutive cycles, corresponding to 86% (Ge) and 57% (Si) retention of the second cycle (i.e., the first cycle at 0.5C) delithiation capacity, which is in line with values observed in conventional organic solutions.19 It has been pointed out that in addition to offering high capacities advanced anode systems should also display high Coulombic efficiencies if they are ultimately to be suitable for application in full cells.34 From Figure 6 and, especially, from Figure S6, it is evident that, excepting the initial cycle, Coulombic efficiencies tending toward 100% were observed during the entire duration of the prolonged cycling tests, once more highlighting the good reversibility of the Li+ alloying process within the Ge and Si NWs in the IL-based electrolytes. To our knowledge, the retention of a capacity exceeding 1000 mA h g−1 after 300 cycles at a rate as high as 0.5C has not previously been reported for any Si-based anode tested in an IL electrolyte. As was observed and discussed in relation to the data of Figure 5A, the capacity of the Si electrode increases during the initial 20 cycles. However, it then undergoes fading at a higher rate compared to the Ge anode. This suggests faster degradation (probably due to higher lattice stress) of Si than Ge, resulting in the lower capacity retention for the Si material. We have previously reported similar trends for Si and Ge NWs in organic-based electrolytes,17 while others51 have provided detailed discussions of the relative mechanical “toughness” of Ge anodes compared to Si. The durability of Ge as a highcapacity anode material is further emphasized here by the fact that a capacity of 890 mA h g−1 is registered after 1000 cycles at 0.5C (Figure 6A and C), corresponding to 67% of the initial reversible capacity. This result is particularly impressive given that the attainment of stable performance over hundreds of cycles remains challenging for conventional carbonaceous anodes in IL electrolytes.28 The results outlined in Figures 5 and 6 for Si show an improvement in terms of capacity, cycling stability, and rate capability, i.e., more than 830 mA h g−1 (corresponding to more than 50% of the initial capacity) upon 500 consecutive discharge/charge cycles (100% DoD) at 0.5C, compared to the only previous report on binder-free Si NW anodes in an IL electrolyte, which focused on CVD-grown, gold-seeded nanowires, and an electrolyte consisting of 1 M LiTFSI in a commercially available butyl trimethylammonium bis(trifluoromethanesulfonyl)imide ionic liquid with 10% propylene carbonate additive.28 In that earlier study, the Si NW anode exhibited a reversible capacity of approximately 1200 mA h g−1 at the 50th (and final reported) cycle, where charge/ discharge was conducted at a slow 0.05C rate. It is reasonable to suggest that the enhanced performance outlined here can be considerably attributed to the superior properties of our tailormade IL. Specifically, it may be surmised that the presence of FSI− in addition to TFSI− anions adds favorable characteristics to our electrolyte, given that a number of studies have recorded higher capacities for Si-based anodes in FSI− ILs compared to their TFSI− analogues.31,34,52 While Sugimoto et al.31 suggested that this effect arises from the lower viscosity (and related higher ionic conductivity) of FSI−-based ILs, Molina Piper et al.34 proposed, on the basis of experimental and modeling results, that decomposition of FSI− on Si surfaces forms an SEI layer with superior properties to that arising from TFSI− breakdown. An alternative explanation courtesy of Usui et al.52 centers on calculations predicting weaker electrostatic interactions between Li+ cations and FSI− anions relative to

TFSI− anions. This favors desolvation of Li+ at the surface of Si anodes in FSI−-based electrolytes, thereby lowering the energetic barrier to charge transfer. In any case, the comparison of our data to the study on Si NWs in the solely TFSI−-based IL28 suggests a positive role for FSI− in assisting lithiation/ delithiation at Si anodes. In our previous works on Ge and Si NW anodes in carbonate-based electrolytes, we observed by ex situ electron microscopy that the wires undergo a morphological transformation during the first 50−100 cycles, to form a network structure of intertwined ligaments with diameters below 10 nm.19−21 Crucially, it was noted the full formation of the network coincided with a decrease in the rate of capacity fading and the onset of highly stable cycling performance. Similar behavior is apparent in Figure 6C, with decreases in the initial rates of capacity decay for the Ge and Si NW anodes after approximately 100 cycles. To ascertain whether this arises from morphological changes similar to those prevailing in carbonate solutions, Ge and Si NW anodes were cycled 100 times at 0.5C in the IL electrolyte, the cells were then disassembled, and the SEI layer was removed to allow imaging of the anodes. Removal of the SEI layer was achieved by soaking the electrodes in acetonitrile for 24 h and then rinsing them with 0.1 mM acetic acid, deionized H2O, and ethanol in that order. Representative SEM, transmission electron microscopy (TEM), and dark-field scanning TEM images are presented for the Ge NW anode in Figure 7A−C. It is immediately

Figure 7. Electron-microscopy images of the active mass of the Ge anode after 100 cycles at 0.5 C. SEM (A), TEM (B), dark-field scanning TEM (C), and higher resolution TEM (D).

evident that the nanowires have transformed into an entanglement of nanofibers just as reported for the same type of anode in organic-based solvent.19 The amorphous ligaments that constitute the network are depicted at higher resolution in Figure 7D, revealing average diameters of 4.0 ± 0.5 nm. A similar morphological evolution was observed for the Si NW anode, with the nanoligaments also displaying diameters of approximately 4 nm (Figure S3). SEM and TEM images are presented in the Supporting Information, characterizing the Ge anode from Figure 6 following the completion of 1000 cycles (Figures S4 and S5). There is little appreciable difference in the images recorded after 100 and 1000 cycles, indicating the durability of the porous network structure once it forms. We 5938


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mA h g−1. In addition, incorporation of EC results in higher initial Coulombic efficiency (75 vs 70%). This is a further indication of the role played by EC in SEI. Work is being addressed in our laboratory aiming to reduce the EC content and investigate different additives. Impedance measurements were conducted at various stages during the prolonged cycling tests depicted in Figure 6, in order to check the effect of the continuous (de)alloying steps (at 100% DoD) on the anode impedance. The results are reported in Figure 8 for the Ge (Figure 8A) and Si (Figure 8B) cells. In

attribute the stable cycling performance of the Ge and Si NW anodes to the attainment of this network structure, which incorporates sufficient space to permit expansion of the ligaments during Li alloying, while retaining excellent electrical conductivity owing to their interconnection. The higher degree of capacity loss during the restructuring process can be explained by the detachment of a proportion of the active mass from the electrode; the rate of this degradation is evidently much reduced upon complete development of the ligament-based framework. It should be noted that, in our earlier studies, it was necessary to add a stable SEI “former”, such as vinylene carbonate (VC), to 1 M LiPF6 in EC/DMC (dimethyl carbonate) electrolyte in order to achieve the network morphology.20 In the absence of such an agent, Ge NWs transformed, upon cycling, into a compact mass of amorphous Ge, which easily became detached from the current collector, leading to poor anode performance. Therefore, it would seem that EC, which is also utilized as a film-forming additive in our IL electrolyte, is, by itself, unable to provide the surface passivation required to facilitate successful evolution of the nanoligament structure. As already discussed, FSI− anions are believed to contribute to durable SEI formation. It is therefore possible that they effectively play the same role in the IL as VC does in the conventional electrolyte and assist in the development of surface protective layers that prevent the fusion of Ge or Si NWs into nonporous masses. The role of the EC additive on the chemical composition of the SEI layer was investigated by ex situ XPS measurements. Two Si NW electrodes were analyzed, one cycled in the 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI with 5 wt % EC and the other without EC, after 50 charge/discharge cycles (Figure S8). The results are presented in Figure S7. The deconvoluted high-resolution S 2p spectra (Figure S7A) show two prominent peaks, both representative of TFSI− and Li2SO4 (169.0 and 170.2 eV).34,53−58 The peaks at lower binding energies correspond to decomposition products of TFSI− and FSI− such as sulfates, sulfites, sulfides, and polysulfide.53−58 Interestingly, when no EC is present in the electrolyte solution (lower spectra of Figure S7), the relative concentrations of decomposition products of the TFSI− and FSI− anions in the SEI layer are much higher than in the electrode cycled with the additive (upper spectra). This suggests that EC is preferentially reduced during discharging. Evidence of this is also present in the C 1s and O 1s spectra, where peaks corresponding to carbonate bonds (C 1s: 289.8 eV; O 1s: 533.2 eV) are more prominent in the electrode cycled with the EC additive (panels B and C).59,60 The C 1s spectrum shows that the SEI layer predominately consists of hydrocarbon species, evidenced by the dominant peak at 284.8 eV corresponding to the C−C bond. Additional peaks at 286.4 and 288.5 eV are most likely due to Li ethers and carboxylic acid species, respectively.59−63 A full study is currently under way to fully elucidate the effect that alkyl carbonate-based electrolyte additives have on the SEI layer of Si electrodes cycled in an ionic liquid electrolyte and will be published in due course. The beneficial effect of EC is also supported by the results coming from the cycling test (current rate of 0.5C), depicted in Figure S8, performed on Si NW anodes on which ex situ X-ray photoelectron spectroscopy (XPS) measurements were run as reported in Figure S7. About 1500 mA h g−1 was delivered in EC-containing IL electrolyte after 50 consecutive discharge−charge cycles, whereas the capacity value in EC-free electrolyte is seen to not exceed 1000

Figure 8. Impedance plots of Li/Ge (A) and Li/Si (B) half-cells upon selected discharge/charge cycles (see legend) run at 0.5C. T = 20 °C.

agreement with other studies on similar anode materials in both conventional21,64 and IL28,65 electrolytes, the Nyquist diagrams consist of a semicircular feature at medium to high frequencies and a linear section at the lower frequency end. The pristine anodes have not yet undergone any lithiation, so in these cases the projection of the semicircle onto the Z′ axis represents the Li/electrolyte interface resistance and Ge(or Si) charge transfer resistance, while the linear tails, which tend to the vertical, have their origin in the capacitive reactance of the electrodes.66 Following the alloying of Li+, the significance of these features changes with the semicircle now corresponding to the overall interfacial impedance (including that associated with the SEI layer),28 while the linear region arises from the diffusion of Li+ within the active material.64 For both the Ge and Si cells there is a sharp decrease in interfacial resistance (as derived from the difference between the high- and low-frequency intercepts of the semicircle with the Z′ axis) between the first and the 200th cycles. The impedance characteristics subsequently become very stable, with the Nyquist diagrams for the 200th and 300th cycles practically overlapping, suggesting steady interfacial resistances, at this stage, of approximately 150 and 200 Ω for 5939


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the feasibility of Ge and Si NWs in composite anodes with ionic liquid electrolytes, which have up until now been poorly reported. Work is in progress in our laboratories to optimize the electrode formulation as well as to investigate their performance in IL electrolyte full batteries. Finally, in order to verify the feasibility of Ge and Si NWs in practical devices, germanium and silicon electrodes were coupled with LCO commercial cathodes and preliminarily tested in IL electrolyte full cells. However, the germanium and silicon anodes and the LCO cathodes were previously prelithiated in Li/Ge, Li/Si, and Li/LCO half-cells (using 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte), which were subjected to full cycles at 0.1C rate and 20 °C (Figure S9). The SOC of the LCO electrodes were adjusted to 20−30% at the second discharge half-cycle (Figure S9C). After prelithiation the half-cells were disassembled and the Ge, Si, and LCO electrodes were combined in full cells. The results, obtained at a current rate of 0.5C and 20 °C, are reported in Figure 10 in terms of voltage vs capacity profiles. Well-defined,

the Ge and Si cells, respectively. This observation correlates well with the data of Figure 6C, where a stable cycling regime with a low rate of capacity decay has become established by the 200th cycle for the Ge and Si anodes. These properties are, of course, a function of the underlying active material morphology, with the results in Figure 8 indicating that a large decrease in interfacial resistance accompanies the transformation of the original discrete nanowires into the stable porous network structure. This may arise from the larger surface area offered for reaction by the network and/or the accommodation of more facile charge transfer by the particular SEI layer composition that has evolved with the nanoligaments. The performance of Ge and Si NWs was preliminarily investigated in higher mass loading composite anodes. The tests were performed at a current rate of 0.1C with the exception of the first cycle (0.02C). The 0.1LiTFSI− 0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC quaternary mixture was used as the electrolyte. The results, depicted in Figure 9 as

Figure 9. Voltage vs capacity profiles of composite Ge (A) and Si (B) anodes in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte at 20 °C. Current rate: 0.1C (0.02C for the initial cycle). Figure 10. Voltage vs capacity profiles of Ge/LCO (A) and Si/LCO (B) full cells in 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte at 20 °C. Current rate: 0.5C.

charge/discharge voltage vs capacity profiles, are comparable with those reported in Figure 3A for the binder-free Ge NW anodes. For instance, a capacity exceeding 1000 mA h g−1 is still delivered at 0.1C (Figure 9A), but with active material mass loading almost 3 times higher. Conversely, Si composite electrodes (Figure 9B) show lower performance with respect to the binder-free anodes (Figure 3B); however a high capacity of 800 mA h g−1 is still achieved. In our opinion, the results are dependent on both the higher mass loading (in terms of capacity per electrode area unit) of Si than Ge and on the optimized formulation of the composite anodes. This latter issue becomes particularly relevant for silicon anodes because of its lower electronic conductivity with respect to that of germanium. However, the preliminary data of Figure 9 indicate

reproducible plateaus are observed in the 3.5−4.0 V range with a rather stable and highly reversible capacity exceeding 1300 (vs Ge) and 1700 (vs Si) mA h g−1 for the Ge/LCO (Figure 10A) and Si/LCO (Figure 10B) full cells, respectively. Even if progressive decay in capacity is observed upon during the initial 20 cycles, more than 1100 (Ge/LCO) and 1200 (Si/LCO) mA h g−1, respectively, are still detected upon 100 consecutive cycles. This is something not commonly reported for Ge and Si anodes, especially in IL electrolyte full batteries. It should be noted that improvement of the cell performance is expected by 5940


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the Ge NW covered substrate was rinsed several times with toluene to remove any organic residue. Tin-seeded silicon NW anodes were fabricated by a slightly different method, which we have also previously reported.17 In this case, a 20 nm thick Sn layer was thermally evaporated onto stainless steel foil, which was then inserted in a long-necked, Pyrex, round-bottomed flask along with a volume (7 mL) of the high boiling point solvent, squalane (Aldrich, 99%). The flask was placed within an upright furnace and attached to an Ar line via a water condenser. The furnace temperature was ramped to 460 °C, and a volume (0.75 mL) of phenylsilane (Aldrich, 97%) was injected through a septum cap into the system. The reaction was allowed to proceed under reflux for 1 h. After cooling, the Si NW covered substrate was rinsed with toluene. Composite anodes were prepared by mixing Ge or Si NW powder (70% in weight), C45 carbon (Imerys, 20 wt %) as the electronic conductor, and sodium alginate (Sigma-Aldrich, 10 wt %) as the binder in deionized water. The anode slurry was cast onto copper foils (20 μm thick) to obtain tapes, which, after massive water removal at 80 °C, were vacuum-dried at 100 °C overnight. Discs of 12 mm were punched from the anode tapes, having a mass loading of about 0.75 mg cm−2, which corresponds to 1.05 and 1.65 mA h cm−2 for Ge and Si, respectively (i.e., considering a nominal capacity of 1400 and 2200 mA h g−1 for Ge and Si). Cathode tapes were purchased from NEI Corporation (NANOMYTE ALD-BE20E). They were composed of LCO (80 wt %) that has been coated with an oxide using atomic layer deposition. Before use, the cathodes (on 25−50 μm thick Al foils) were vacuum-dried at 100 °C overnight. The mass loading approaches 6.2 mg cm−2, which corresponds to about 0.87 mA h cm−2 (i.e., considering a nominal capacity of 140 mA h g−1 for LCO). Scanning electron microscopy imaging was conducted with a Hitachi SU-70 system operating between 3 and 20 kV. For transmission electron microscopy analysis, the NWs were removed from the substrates by sonication. TEM characterization was performed using a 200 kV JEOL JEM-2100F field emission microscope. X-ray photoelectron spectroscopy measurements were performed using a Kratos ULTRA spectrometer with a monochromatic Al K 1486.58 eV. C 1s at 284.8 eV was used as the charge reference to determinate the core level binding energies. The pass energy of 160 eV was used for the survey spectra and 20 eV for the narrow regions. Construction and peak finding of synthetic features in narrow region spectra used a Shirely-type background. The synthetic features were of a mixed Gaussian−Lorenzian type. Cell Assembly. The electrochemical performance of germanium NW and silicon NW anodes was evaluated in Li/Ge and Li/Si halfcells, respectively. Electrodes (having sizes of about 1 × 0.8 cm) were directly cut from the NW-covered foils and weighted using an ultramicrobalance. For both Ge and Si NW anodes, the active mass loadings were in the range of 0.2 to 0.4 mg cm−2. The (Li/Ge and Li/Si) half-cells as well as Ge/LCO and Si/LCO full batteries were manufactured in a dry-room (moisture content below 10 ppm) by sandwiching a Ge or Si electrode, a glass fiber separator, and a lithium metal electrode (50 μm thick) or a LCO cathode (12 mm diameter disc). A lithium titanium oxide strip (plated onto an aluminum current collector) was used as the reference electrode for the full batteries. The ionic liquid electrolyte was loaded into the separator before cell assembly by applying vacuum for a few minutes. Successively, the cells were housed in soft-pouch envelopes, evacuated for 1 h, and then vacuum-sealed. Electrochemical Tests. Impedance measurements were carried out on two-electrode Li/Ge and Li/Si half-cells (assembled as reported above) at open circuit using a frequency response analyzer (Solartron Schlü mberger model 1260 F.R.A.) coupled to a potentiostat (Solartron Schlümberger model 1287 electrochemical interface). The ac plots were taken in the 100 kHz−1 Hz frequency range by applying a 10 mV amplitude voltage (ΔV) at 20 °C. Cyclic voltammetry was run (20 °C) on Li/Ge and Li/Si half-cells using a VMP3 Bio Logic potentiostat within the 0.01−2.0 (vs Li/Li+)

optimizing the formulation and balance of the electrodes as well as the cell design. Extended results on full cells will be reported in a future publication.

CONCLUSIONS Germanium and silicon nanowires as lithium-ion battery anodes were investigated in combination with a quaternary, ionic liquid-based, 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC electrolyte. Cyclic voltammetry and electrochemical impedance measurements indicate a highly reversible Li alloying process and excellent compatibility between the Ge and Si nanowire anodes and the ionic liquid. Initial reversible capacity values of approximately 1400 and 2200 mA h g−1 were observed at 0.1C rate for the Ge and Si electrodes, respectively, matching those recorded for these anodes in conventional electrolytes. The ionic liquid formulation also facilitates a desirable structural modification of the active material, upon repeated lithium alloying/dealloying, with the nanowires evolving into a porous network of sub-5 nm diameter ligaments. The stability of this morphology underlies the impressive cycling performance of the anodes in the ionic liquid, with the Si nanowires retaining a capacity above 830 mA h g−1 after 500 cycles and the Ge nanowires delivering 890 mA h g−1 after 1000 cycles, where both were cycled at 0.5C (100% DoD). Full cells, manufactured by coupling Ge and Si anodes with LCO cathodes, have exhibited stable capacity values above 1300 (Ge) and 1700 (Si) mA h g−1 (with respect to the anode material), respectively. Therefore, these results show that ionic liquids may be successfully used as the main electrolyte component in combination with nanostructured alloying anodes, for the realization of safer, high energy and power density lithium-ion batteries. METHODS Ionic Liquid Electrolyte. N-Methyl-N-propylpyrrolidinium bis(trifluoromethanesulfonyl) imide, PYR13TFSI, and N-methyl-Npropylpyrrolidinium bis(fluorosulfonyl)imide, PYR13FSI, ionic liquids were synthesized through a procedure reported elsewhere.67 The chemicals N-methylpyrrolidine (Aldrich, 98 wt %), 1-bromopropane (Aldrich, 99 wt %), activated charcoal (Aldrich, Darco-G60), alumina (Aldrich, acidic, Brockmann I), LiTFSI (3M, 99.9 wt %), and NaFSI salts (Solvionic, 99.9 wt %) were used as received. Ethylene carbonate was purchased from Merck (battery grade reagent, 99.9 wt %) and used as received. The electrolyte, 0.1LiTFSI−0.3PYR13TFSI−0.6PYR13FSI + 5 wt % EC (where 0.1, 0.3, and 0.6 are the mole fractions of the three salts, while the EC content is given as weight percent with respect to the overall weight of the salts) quaternary mixture, was prepared by dissolving (in the proper proportion and at 40 °C for a few minutes) LiTFSI (vacuum-dried overnight at 120 °C) in the PYR13TFSIPYR13FSI blend. Finally, the proper amount of EC was added, resulting in fast dissolution. The ionic liquid electrolyte was stored in sealed glass vials within the dry-room. Preparation and Characterization of Germanium and Silicon Nanowire Anodes. Copper germanide (Cu3Ge)-seeded germanium NW anodes were prepared following our previously published procedure.19 Briefly, a 1 nm thick Cu layer was thermally evaporated onto stainless steel foil (SS, 316, Pi-Kem Ltd. UK), which was then placed onto a Stuart CD162 digital hot plate within a dry, Ar-filled glovebox. A cylindrical stainless steel “confiner” tube was pressed onto the foil, enclosing the area where NWs were to be grown, and the hotplate was set to 425 °C. Upon attainment of this temperature, an appropriate volume (12 μL per cm2 target growth area) of the reaction precursor, diphenylgermane (DPG, Gelest, 97%), was dropped onto the foil within the confiner. The reaction continued until the escape of DPG vapor from the confiner tube was complete. Following cooling, 5941


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ACS Nano voltage range at a 0.1 mV s−1 scan rate. All potential values quoted in this work refer to the quasi-reference redox couple Li/Li+. The lithium half-cells were investigated through galvanostatic cycling measurements performed by a Maccor 4000 battery cycler. The tests were run at different current rates (from 0.02C through 3C) and at 20 °C (by climatic chambers) within the 0.01−2.0 (vs Li/Li+) voltage range. The capacity of the electrodes was calculated based on the mass of the Ge or Si NWs. The Ge/LCO and Si/LCO full cells (anode limited) were tested at 0.5C and 20 °C within the 1.6−4.0 (vs Li/Li+) voltage range. The cell capacity was calculated with respect to the mass of Ge or Si anodes.

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ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.7b01705. Additional figures (PDF)

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Phone: +39 06 30483924. ORCID

Kevin M. Ryan: 0000-0003-3670-8505 Giovanni B. Appetecchi: 0000-0002-6623-0373 Notes

The authors declare no competing financial interest.

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