Temperature Dependent Li Storage Performance in Nanoporous Cu

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Temperature Dependent Li Storage Performance in Nanoporous Cu-Ge-Al Alloy Wenqing Ma, Yahui Wang, Yijun Yang, Xi Wang, Zhihao Yuan, Xizheng Liu, and Yi Ding ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b20654 • Publication Date (Web): 11 Feb 2019 Downloaded from http://pubs.acs.org on February 11, 2019

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ACS Applied Materials & Interfaces

Temperature Dependent Li Storage Performance in Nanoporous Cu–Ge–Al Alloy Wenqing Ma,† Yahui Wang,‡ Yijun Yang,+,* Xi Wang,+ Zhihao Yuan,†,‡ Xizheng Liu,‡,* and Yi Ding†,‡ †

School of Materials Science and Engineering, Tianjin University, Tianjin 300350, P. R. China.

‡

Tianjin Key Laboratory of Advanced Functional Porous Materials, Institute for New Energy

Materials and Low-Carbon Technologies, School of Materials Science and Engineering, Tianjin University of Technology, Tianjin 300384, P. R. China. +

School of Science, Beijing Jiaotong University, Beijing 100044, P. R. China.

Corresponding author: [email protected] (X. Liu) or [email protected] (Y. Yang)

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Abstract

Lithium ion batteries (LIBs) process performance fading and safety concerns at low temperature (LT) prohibit their application in cold climates. The alloy-type electrodes demonstrate great potentials in stable and dendrite-free anodes at LT. Herein, we report a temperature-dependent Li-storage performance in Al-based nanoporous alloy anode. The nanoporous-structured Cu-Ge-Al ternary alloys (NP-CuGeAl) have been designed and prepared by selectively etching Al out. The high Al-content NP-CuGeAl (acid etching for 6 h, named as CGA-6) is composed of multi-intermetallic compounds (denoted as MxNy, M, N = Cu, Al, Ge) with bimodal porous architectures. Investigated as anode, at room temperature (RT), the CGA-6 delivers a capacity as high as 479.7 mAh g-1 at 0.5 A g-1 over 1020 cycles. And the low Alcontent ones show improved LT electrochemical performance. At -20 °C, the CGA-48 (acid etching for 48h) shows much better performance as compared with the CGA-6. In situ TEM and ex situ characterizations confirm that the MxNy/LizMxNy couples are highly reversible and the porous structure is durable upon battery cycling.

Keywords: Low temperature, Lithium ion battery, Anode, Al-based alloy, Nanoporous


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1. Introduction Lithium ion batteries (LIBs) are widely accepted as dominant electrical energy storage/release mediators for consumer electronics and booming electric vehicles (EVs).1-6 However, commercial available LIBs are usually suffer from severe performance degradation when they operated in winter and high-latitude (altitude) area.7-10 The LIBs’ poor low-temperature (LT) Li storage performance thus resulted in limited reliability for their application for EVs. Its something we called "temperature anxiety" or temperature-dependent "range anxiety". To overcome this limitation, substantial efforts, such as electrolyte optimizing by developing effective additives/solvents,9-11 heating by self-powered device12,13 and developing new promising alternative electrode materials8,14,15-19 have been devoted for LIB LT performance enhancement. In the view of materials science, as we know, the poor LT electrochemical performance of LIBs are mainly associated with the electrode materials sluggish kinetics.7-10,14,15 Therefore, it is worthwhile to continue the screening for effective electrode materials with decent kinetics for high performance LT LIBs. Among them, porous structured alloy materials have garnered particular attention thanks to their interconnected channels for electrolyte permeation, fast Li+ transportation and (de)lithiation-accompanied volume changes accommodation. Very recently, Passerini et al. reported the porous Cu-Zn intermetallic alloys for LT Li-storage. The porous Cu-Zn alloys show competitive LT performance, as compared with commercial graphite, with a fast Li-(de)alloying kinetics.20 However, to date, the LT LIB anode preferred structure and contents, which undoubtedly useful for materials designing, is still blind to us. Among all alloytype alternatives, Al-based alloy is a special one mainly because its low-cost and environmental friendliness nature.4,21-23 However, so far there is no satisfactory application of Al-based anodes because of the similar problems encountered with all alloy-type LIB anodes, such as rapid 3 ACS Paragon Plus Environment


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capacity decrease and limited cyclability caused by the notorious volumetric changes and unstable solid electrolyte interface (SEI) film upon battery cycling.2-4,21-24 Additionally, the sluggish lithiation/delithiation kinetics of Al-based anode remain a big obstacle during high rate operation.4,25-28 To overcome these issues, several strategies have been developed including: i) alloying with Li-active or/and -inactive elements to boost the lithiation/delithiation kinetics as well as alleviate materials pulverization,29-31 ii) utilizing low-dimension Al films or nanoparticles to improve the Li+ diffusion and the volume changes,21,23,26,32,33 and iii) introducing electrolyte additives to suppress the deterioration of SEI and the consumption of the electrolyte.3 For instance, Rahner et al. fabricated a eutectic mixture of Al-Ni (Al/Al3Ni), which not only facilitates the Li+ transportation but also promotes the mechanical stability of the whole electrode.29 Additionally, some other Li-active/-inactive elements (M = Fe, Mn, Si, Sn, etc.), which can form intermetallic compounds with Al to accelerate electron/ion transfer as well as alleviate the volumetric changes, have also been reported and inspiring breakthroughs have been achieved.4,27,30,31,34 It is widely known that dealloying is an effective method for larger scale fabrication of various catalysts and electrodes that show high surface area porous structures.1,22,35-40 However, previous reports on porous-structured alloy-type anode construction, by dealloying Al-based alloy, are tend to thoroughly etching Li-active Al out because of its sluggish kinetics and unstable SEI film upon battery cycling.41 Herein, we upgrade the nanoporous strategy, by direct dealloying a CuGeAl ternary alloy in acid solution, by which the Al was purposefully (partially) reserved through optimizing the porous structure and components. Upon Al dissolution, the left Al, Cu and Ge atoms selforganize into interconnected network with pores and ligaments. The chemical constituents and porous structure (morphology) of these porous CuGeAl ternary alloys were systematically 4 ACS Paragon Plus Environment

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studied. Moreover, the electrochemical performance of these samples were both evaluated at room temperature (RT) and LT (-20 ºC). Interestingly though, the high and low Al-content samples show a contrary temperature-performance correlation. Furthermore, XRD and in-/ex-situ TEM were used to disclose the composition–structure–property correlation in the lithiation/delithiation processes. This study provides a good example for guiding Li-storage material research for different temperature utilization.

2. Experimental 2.1. Materials preparation The Cu-Ge-Al ternary alloy (Cu15Ge15Al70, at.%) was prepared by weighting the pure ( > 99.99 wt.%) Cu, Ge, and Al as designed, melting in an Ar-filled arc furnace. A single roller meltspinning technique was then adopted to get the Cu-Ge-Al alloy ribbons under Ar atmosphere. The Cu-Ge-Al ternary alloy ribbons were then dealloyed in 0.1 M HCl solution at 60 °C for different time for composition and morphology optimization. The as-obtained samples were rinsed with deionized (DI) water, dried in a vacuum oven at room temperature and named as CGA-X (X represents the dealloying time, by the hour).

2.2. Materials characterizations The compositions of the CGA-X were analyzed by inductively coupled plasma (ICP). The XRD was conducted on a powder diffractometer equipped with Cu-Kα radiation (Rigaku Ultima IV, λ = 1.5405 Å, operated at 40 kV and 30 mA). The N2 adsorption/desorption isotherms were measured on a Quantachrome Autosorb-iQ. Brunauer-Emmett-Teller (BET) and Barrett-JoynerHalenda (BJH) methods were correspondingly adopted for specific surface area and pore size 5 ACS Paragon Plus Environment


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distribution calculation. The morphology was demonstrated by scanning electron microscope (SEM, Verios 460L, FEI) and transmission electron microscopes (TEM, Tecnai G2 Spirit TWIN and Talos F200X, FEI). In situ TEM experiment was conducted using a JEOL-2100F (with Omega filter and Nanofactory Instruments STEM-TEM holder). To fabricate the test cell, a Au rod (attached to the piezo-manipulator) was used as the carrier and current collector of the CGA6 nanoparticle. A small amount of Li on a W rod, which was covered by a thin Li2O (naturally oxidized from Li) solid electrolyte layer, was used as both counter (CE) and reference (RE) electrode. A bias voltage of -1 V (versus Li+/Li) was applied to CGA-6 to drive the lithiation process, and the electron beam irradiation effect was avoided except for image acquiring.

2.3. Electrochemical measurements The Li storage performance of CGA-X was characterized with 2032 coin cells. To prepare the working electrodes, a mixture of CGA-X, ketjen black, and sodium carboxymethyl cellulose (7 : 2 : 1, wt.%) was dispersed into DI water and spread onto a Cu foil and drying in a vacuum oven (60 °C, 12 h). LiPF6 was dissolved into a mixture of dimethyl carbonate and ethylene carbonate (1 : 1, v : v) to form a 1.0 M electrolyte. A pure Li foil and Celgard-2400 were correspondingly utilized as the CE and separator. The battery was assembled in a glove-box (Ar-filled) and tested at different temperatures (-20 and 25 °C). The LT performance was performed by using a lowconstant temperature test box (Shipac DT6040A), which was set to -20 °C. Cycling and rate performance were studied in the voltage window of 0.01–3.0 V (versus Li+/Li) on a Land CT2001A. Cyclic voltammetry (CV) profiles were collected on a CHI 760E electrochemical workstation (scan rate: 0.1 mV s-1). A Princeton ParSTAT MC electrochemical workstation was

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used to collect the electrochemical impedance spectroscopy (EIS) data (frequency range: 105 – 10-2 Hz).

3. Results and discussion The digital photo and SEM image of the Cu-Ge-Al alloy ribbons are shown in Figure S1a and b. The width of the alloy ribbons is approximately 0.5 cm. The XRD pattern (Figure S1c) indicates that Cu-Ge-Al alloy precursor is composed of simple substance Al (JCPDS# 04-0787) and multi intermetallic compounds (AlCu3, CuAl2, Al21Ge8, AlGe0.064, AlGe0.045, Al6Ge5, AlGe0.333, Al5Ge3, etc.). The Cu-Ge-Al alloy precursor was immersed into 0.1 M HCl aqueous solution to construct porous structure by selectively etching some of Al as schematically shown in Figure S1d. For structure and component optimization, the porous Cu-Ge-Al alloys were prepared with different dealloying time and denoted as CGA-X in which X represents the dealloying time (by the hour). The as-obtained CGA-3 exhibits high porosity and continuity with unimodal pores and ligaments both size around 50 nm (Figure S2a). As expected, the characteristic peaks indexed to elemental Al falls obviously as demonstrated by XRD (Figure S2b). Whereas, the Al-based MxNy still remains but evolved to low-Al-content ones (e.g. AlGe2, AlGe0.064, etc.). Interestingly though, during this process, the CumGe alloy phases (CuGe and Cu3Ge) formed which result from the alloying between the exposed Cu0 and Ge0 by dealloying Al-based MxNy compounds.33,36 Previous works have shown that the intermetallic compounds (MxNy) can not only boost the electronic conductivity but also accommodate the volume changes of the alloy-type



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Figure 1. (a) XRD patterns; (b) N2 adsorption/desorption isotherms and BJH pore size distribution; (c, d) different magnification SEM images; (e-g) TEM and HRTEM images; (h) SAED pattern and (i) HAADF-STEM image and (i1-i3) its corresponding EDS elemental mapping results of CGA-6 composite. See Figure S1-3, S5 and Table S1 for more details.


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anodes simultaneously.4,41-44 As dealloying time goes on, the characteristic peaks of metallic Al disappeared, while the Al-based MxNy compounds further evolved to Cu3Al2, Al6Ge5, etc (Figure 1a). Moreover, the intensity of CuGe and Cu3Ge peaks increased. The typical bimodal porous architectures formed as disclosed in Figure 1c and 1d. The pores size around 150 and 25 nm are both clearly observed. Over time, however, the bimodal porous structure disappeared in CGA-24 and CGA-48, and the pores further enlarged to about 300 nm (Figure S2e and S2i). Furthermore, all of the Al-based MxNy compounds disappeared (Figure S2f and S2j). Meanwhile, the intensity of Ge characteristic diffraction peaks increased. The mass fraction of these samples were quantitatively estimated by ICP analysis and summarized in Table S1. The mass fraction of Ge and Al in CGA-6 are 38.39 and 24.56 wt.%, respectively. As we know, only Ge and Al contribute to Li-storage and can be considered as active components. Therefore, the active components weight percentage in CGA-6 can be calculated as 62.95 wt.%, which is comparable to previous results (Table S2).3,21,23,24,45 The BET data was further collected to disclose the porosity of these samples (Figure 1b, S3, Table S1). The specific surface area of CGA-6 (44.75 m2 g-1) is higher than CGA-3 (32.09 m2 g-1), CGA-24 (25.23 m2 g-1) and CGA-48 (22.21 m2 g-1), which means a decent electrolyte/electrode interface. The difference of specific surface area of CGA-X should be related to the evolution of the porosity/continuity of the pores/ligaments upon Al dissolving.35 The pore distribution curve shown in Figure 1b (inset) confirmed the pore size for CGA-6 around 20 and 1 nm. The pores ~ 1 nm indicates that some fine pore structures might exist in the internal of CGA-6. These nanopores provide decent void space for volume accommodation during discharging as well as admirable channels for electrolyte permeation and fast kinetics.22,37


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To clarify the relationships between porous structure and electrochemical performance, the RT Li storage performance of these samples were evaluated by galvanostatic discharge/charge experiments, as displayed in Figure 2. For convenience, the electrochemical performance (e.g. current density, specific capacity) evaluation was based on the mass sum of Cu, Ge, and Al in this work. The initial discharge capacities of CGA-3, CGA-6, CGA-24 and CGA-48, are 1327.8, 1097.3, 980 and 863.15

Figure 2. The room temperature electrochemical performance of the (a) CGA-3; (b) CGA-6; (c) CGA-24 and (d) CGA-48 electrodes. See Table S3 for more details. mAh g-1 as summarized in Table S3, respectively. The first discharge capacity evolution should be arising from the difference of the relative content of Li-inactive Cu in these samples. Moreover, the CGA-X galvanostatic discharge/charge and dQ/dV curves shown in Figure S4 indicates that the difference of the specific surface area of these CGA-X samples should be another reason responsible for the CGA-X initial discharge capacity evolution (see Figure S3 and


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S4 and their captions for more details). Strikingly, the CGA-6 shows admirable rate capability. The discharge capacities at 0.5, 2.0, 4.0, 8.0 and 16.0 A g-1 are 364.43, 180.29, 118.58, 78.21 and 63.58 mAh g-1, respectively. Afterwards, the 121st cycle was adopted as the benchmark for cyclability evaluation. The capacity fading for CGA-6 is ~ 0.0014% per cycle after 121 cycles which is much better than CGA-3, CGA-24 and CGA-48. The fine structural and morphological information of CGA-6 were collected to disclose the origination of superior electrochemical performance. Figure 1e shows that both the large- and meso-pores are open, which benefits the electrolyte permeation and volumetric fluctuation alleviation. Moreover, the HRTEM images reveal the MxNy nanograins are tightly accumulated in the ligaments (Figure 1f, S5). The (202) and (211) crystal faces of CuAl2 nanoparticles can be clearly observed in Figure 1g. Additionally, the lattice fringes of some other MxNy (e.g. Cu3Ge, Al6Ge5, etc.) could be also clearly observed (Figure S5). Moreover, the selected area electron diffraction (SAED) pattern confirms the formation of multi-MxNy compounds and their high crystallinity (Figure 1h). The elemental mapping result shows that the Ge, Al, and Cu are uniformly packed in the ligaments (Figure 1i-i3). We also acquired the SEM images of the CGA-X electrodes after the initial discharging (Figure S6). The pores in CGA-X electrodes are still remained after initial discharging. Confirmed by these results, the 3D bi-continuous and hierarchically porous structure and multiple intermetallic compounds (MxNy) corporately enhanced the rate capability and cyclability.


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Figure 3. (a) Rate capability and cyclability; (b) the initial three galvanostatic discharge/charge profiles; and (c) Nyquist plots after the first cycle of the CGAmix and CGA-6 electrodes. Longterm (d) rate and (e) cycling performance of CGA-6 electrodes (see Figure S7-9 for more details). To further illustrate the advancement of porous structure in CGA-6 electrode composite, by contrast, the pure Al, Ge and Cu powder was also adopted as anode as shown in Figure 3a. CGA6 displays discharge capacities of 544.70, 381.15, 288.41, and 186.35 mAh g-1, at 0.1, 0.5, 1.0, and 3.0 A g-1, respectively. The discharge capacity retentions are about 70.0% (0.5 A g-1), 52.9% (1.0 A g-1), and 34.2% (3.0 A g-1) of the capacity at 0.1 A g-1, respectively. For comparison, the pure Cu, Ge, and Al powders were mixed obey the element content of CGA-6 (analyzed by ICP) and denoted as CGAmix. The CGAmix was also subjected to similar test, and shows relatively high discharge capacity of 636.93 mAh g-1 at 0.1 A g-1, but fading fast in subsequent cycles. In a word, the porous structure obviously improved the rate performance. Figure 3b displays the


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initial three discharge/charge curves of CGA-6 anode (the first one and the latter two cycles were performed at 0.05 and 0.1 A g-1, respectively). The capacities dropped quickly at the first cycle, and there are two main reasons account for this phenomenon. The initial two cycles were performed at different current densities, and the capacity would decrease with current density increasing. Moreover, in the first cycle, the electrolyte would decompose on the electrode with capacity irreversibly decrease. The voltage plateau at ~ 0.7 V in the first discharge can be ascribed to the SEI film formation process.9,14,36 The 2nd and 3rd charge/discharge profiles overlay well, which indicates the lithiation/delithiation processes are highly reversible.23,46 Figure S7 shows the CV curves of CGA-6. The broad irreversible reduction peak (~ 0.7 V versus Li+/Li) in the initial cathodic scan can be attributed to the formation of the SEI film. Moreover, the 2nd cycle CV curve overlaps well with the 10th one. These observations are well consistent with charge/discharge profiles shown in Figure 3b. The EIS data collected after the first cycle further revealed that the rigid Cu/Al conductive framework endows the CGA-6 much better charge transfer and fast kinetics behaviors than CGAmix (Figure 3c and S8).47 Remarkably, CGA6 exhibits superior long-term rate cyclability. As shown in Figure 3d and S9a, the reversible discharge capacities at 0.5, 1.0 and 3.0 A g-1 are 697.8, 555.9, and 224.1 mAh g-1, respectively. When the current densities changed from 3.0 to 0.5 A g-1, the capacities of CGA-6 higher relapsed. A capacity as high as 479.7 mAh g-1 can be retained over 1020 cycles. Strikingly, a discharge capacity of 191.9 mAh g-1 is still retained after 8000 cycles at 4.0 A g-1 (Figure 3e, S9b). Notably, the CGA-6 capacity slightly increased along with battery cycling. The gradual activation of the Li-active materials and permeation of the electrolyte with SEI film irreversible formation should be responsible for this phenomenon. The capacity gradually increasing has


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been previously reported,6,41,45 and should be further investigated to find out the specific reasons. And, this result is among the best Al-based anodes reported recently (Table S2).3,21,23,24,45

Figure 4. In situ TEM observation of the morphology and structure evolutions of an individual CGA-6 nanoparticle upon the initial discharge. Schematic drawing of (a) the test cell, and (b) the CGA-6 lithiation process. Time series evolutions of (c-f) the whole particle and detailed analyses by focused on one of the (c1-f1) particles and (c2-f2) pores. The scale bar is 50nm. The superior cyclability should be originated from the 3D nanoporous feature which endows CGA-6 with facile strain alleviation and electron/Li+ transportation.4,48


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Figure 5. (a) TEM and (b-f) HRTEM images acquired after 160 s lithiation. (c1) The boundary of crystalline and amorphous area of the lithiated CGA-6 electrode (highlighted with yellow dashed lines). To better understand the reaction mechanism, in situ TEM was performed to monitor its morphological and structural evolutions upon the first lithiation. Figure 4a schematically shows the all-solid nanoscale LIBs. The lithiation process is schematically shown in Figure 4b in which the pores accommodated the volume expansion. Figure 4c-f shows TEM images of an individual CGA-6 nanoparticle during its first discharge. The porous could be clearly observed. The length of this particle is 396 and 439 nm before and after 160 s lithiation, respectively. For more details, we monitored the evolutions of one particle (pore) to further understand the volume change alleviation mechanism during lithiation. As shown in Figure 4c1-f1 and c2-f2, the focused particle expands from 127 to 143 nm, while the tracked pore size reduced from 29 to 14 nm before and after Li-alloying for 160 s. As a result, no obvious structural degradation and pulverization observed after lithiation, which indicates that the large, theoretically, volume expansion induced by Li-alloying could be effectively accommodated by nanoporous strategy which means decent mechanical stability.


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The structural information of the lithiated CGA-6 was further studied by HRTEM. Several representative micro-regions in Figure 5a were measured, as displayed in Figure 5b-f. The fabric-type one-/two-dimension lattice fringes of the lithiated nano-grains (LizMxNy, e.g. Li3CuAl5, Cu2GeLi, etc.) could be clearly observed and ascribed to lithiated products of MxNy in the pristine CGA-6. In other words, Al has participated into the lithiation/delithiation processes by alloying/dealloying with Li in MxNy/LizMxNy forms. It should be noted that, the amorphous area between two crystalline particles might be ascribed to the lithiated products and/or some SEI component (Figure 5c1 and S10). Because the crystalline LizMxNy and amorphous regions closely

Figure 6. (a) SEM images of the fresh (left) and retrieved (right, 200 cycles) CGA-6 electrodes. (b) TEM and (c-d1) HRTEM images of CGA-6 composite obtained at 200th discharging state. (e, e1) HRTEM images of CGA-6 acquired at 200th charging state. (see Figure S11-14 for more details) packed in the rigid framework, they can buffer the strain induced by lithiation/delithiation and support for the superior electrochemical performance. Conclusively, the 3D bi-continuous


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structure of the porous CGA-6 composites could improve the ions/electrons transmission and support for the structural integrity during repeated lithiation/delithiation processes. To further clarify the morphological and structural evolutions of CGA-6 along with battery cycling, SEM, HRTEM, and XRD were performed at different cycling stages. As shown in Figure 6a and S11, the CGA-6 electrode after 200 lithiation/delithiation processes (right) exhibited a surface morphology as smooth as the fresh one (left). Figure 6b revealed that both the large- and meso-pores of CGA-6 are still remained. Moreover, most of the nanoparticles maintained their initial morphology, although the volume expansion is inevitable (Figure S11). Moreover, the uniform SEI films are also clearly observed on the ligament surface, which is of great importance to avoid electrolyte further consumption and also promote Li+ transmission (Figure S11 and S12). The lattice fringes after discharged to 0.01 V could indexed well to the LizMxNy (e.g. Li3CuAl5, Cu2GeLi, etc.), which is agree well with the in situ TEM observations (Figure 6c-d1, and S13). The structural evolution of the CGA-6 electrodes can also be further clarified by the XRD patterns (Figure S14). The diffraction peaks of the LizMxNy (e.g. Li3CuAl5, Cu2GeLi, etc.) appeared during discharging, but disappeared after the charge process. However, no GeLix related peaks could be observed, which is mainly because the amorphous property. Meanwhile, as shown in Figure 6e, e1, and S13, the fringe spaces of the MxNy (e.g. CuAl2, Cu3Ge, etc.) nano-grains can be clearly observed at 200th charging state. Based on these analyses, the conversion between MxNy and LizMxNy is highly reversible and accounts for the structural integrity and fast ions migration of the CGA-6 composite.29,30,41,42 Therefore, the high reversibility of the MxNy/LizMxNy conversion couple and the hierarchically porous structure altogether endow CGA-6 a fast kinetics and structural integrity upon lithiation/delithiation.


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Figure 7. The low temperature electrochemical performance of the (a) CGA-6 and (b) CGA-48 electrodes. (c) The room temperature electrochemical performance of the CGA-48 electrode. (d) Schematic illustration of the capacity retention of CGA-6 and CGA-48 electrodes when the temperature is lowered (from 25) to -20 °C (see Figure 3a, S15 and Table S4, S5 for more details). To further explore the electrochemical performance at different temperatures, the batteries utilize CGA-6 and CGA-48 anodes were comparatively tested at -20 and 25 °C. As we know, the material conductivity would decrease and the electrolyte viscosity would increase when the temperature drops.20 As a result, the specific capacity degradation of CGA-6 can be observed at variety current densities (Figure 3a, 7a and 7d; Table S4 and S5). Interestingly, however, the CGA-48 shows improved specific capacity retention in comparison with CGA-6 when lowered the temperature (from 25) to -20 °C (Figure 7b, 7c; Table S4). Strikingly, when the current density was set back to 1.0 A g-1, the CGA-48 still delivers 122.9 mAh g-1, while the value of


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CGA-6 is only 63.6 mAh g-1. That is, the CGA-48 shows better LT (-20 °C) electrochemical performance than the CGA-6 which exhibits superior RT Li-storage performance. The reason account for this might be that the element Al kinetics can be effectively enhanced at RT, but cannot at LT (-20 °C). In consequence, the LT electrochemical performance of CGA-3, CGA-6, CGA-24 and CGA-48 increased accordingly (Figure 7b-d, S15; Table S4, S5).

4. Conclusions In conclusion, 3D nanoporous structured CuGeAl anodes (NP-CuGeAl) with different Alcontent have been prepared and their applicability at different temperatures (-20 and 25 °C) were evaluated. The chemical component, structure and morphology of these CuGeAl anodes can be readily optimized. The high Al-content NP-CuGeAl composite is constructed by multiintermetallic compounds (MxNy) with hierarchically porous structures. Over time, the element Al keeps dissolving with materials morphology evolving. Utilized as anodes, these NP-CuGeAl composites show temperature-dependent Li-storage performance. The high Al-content NPCuGeAl composite shows admirable electrochemical performance as compared to the low Alcontent ones at room temperature. Interestingly, the lower Al-content NP-CuGeAl composite shows the better LT Li-storage performance. The present work of utilizing porous structure and intermetallic compounds to improve the kinetic and volume change accommodation ability provides new guidelines for alloy type anodes engineering especially for low temperature lithium ion batteries. Supporting Information The Supporting Information is available free of charge on the ACS Publications website.


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Morphological and structural information of the Cu-Ge-Al ternary alloy ribbons; SEM images, XRD patterns, N2 adsorption/desorption isotherms and chemical components of the CGA-X composites; initial galvanostatic charge/discharge and typical dQ/dV curves of the CGA-X electrodes; TEM and HRTEM images of CGA-6; SEM images of the first discharged CGA-X electrodes; CV curves and Nyquist plots of CGA-X and CGAmix electrodes; SEM, TEM, HRTEM and XRD analysis of CGA-6 after 200 cycles; RT and LT rate and cycling performance of the CGA-X. The authors declare no competing financial interest.

Acknowledgements This work was financially supported by the National Natural Science Foundation of China (No. 21603162 and 51802013), Tianjin Municipal Science and Technology Commission (17JCYBJC21500), Tianjin Municipal Major Project of New Materials (16ZXCLGX00120) and the Fundamental Research Funds of Tianjin University of Technology.

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Temperature Dependent Li Storage Performance in Nanoporous Cu

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