Amorphous Bimetallic Co3Sn2 Nanoalloys Are Better Than Crystalline

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Amorphous Bimetallic Co3Sn2 Nanoalloys Is Better Than Crystalline Counterparts for Sodium Storage Jian Zhu, and Da Deng J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b05232 • Publication Date (Web): 26 Aug 2015 Downloaded from http://pubs.acs.org on August 27, 2015

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The Journal of Physical Chemistry

Amorphous Bimetallic Co3Sn2 Nanoalloys Is Better Than Crystalline Counterparts for Sodium Storage Jian Zhu, Da Deng* Department of Chemical Engineering and Materials Science, Wayne State University, 5050 Anthony Wayne Dr., Detroit, Michigan, United States, 48202 KEYWORDS: Amorphous, Crystalline, Bimetallic, Nanoalloy, Sodium-Ion Batteries.

ABSTRACT. Sodium-ion batteries are considered as a promising alternative to replace the existing lithium-ion batteries for energy storage due to the benefits of low cost and safety. However, it is still challenging to develop suitable electrode materials for reversible storage of sodium. Metal anodes have high capacity for sodium storage, but suffer the issue of poor cyclability due to pulverization caused by large volume variation and electrode disintegration. To address this issue, amorphous bimetallic active-inactive nanoalloy Co-Sn with Sn acts as high capacity active compound and Co acts as conductive inactive matrix has been explored here. We demonstrated that amorphous nanoalloys exhibited superior electrochemical performances as compared to the low-crystalline nanoalloys and crystalline counterpart nanoalloys as negative electrode materials for sodium ion batteries. The degree of crystallinity has negative effects on electrochemical performances. The improved performance of amorphous nanoalloys could be

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attributed to the easy accessibility for sodium ions, strain accommodation and defect sites to host sodium ions.

Introduction Sodium-ion batteries (SIBs) are attracting much attention again recently. Sodium is a widely available, abundant and relatively low-cost element.1 Arguably, SIBs could be much cheaper than LIBs with comparable energy density.1 Additionally, due to the lower half-reaction potential of sodium (2.7 V for Na/Na+) relative to lithium (3.03 V for Li/Li+), safe electrolyte with lower decomposition potential can be used in SIBs.2 In other words, SIBs could be safer than popular lithium-ion batteries (LIBs). Large-scale battery systems for grid application will require nextgeneration batteries produced at low cost. Although SIB technology is still at germination stage, it is very promising to replace the expensive LIBs. However, due to the larger radius of Na ions as compared to Li ions and their different ionic coordination properties (sodium ions prefer to coordinate at prismatic or octahedral sites), sodium ions have much slower solid-state diffusion kinetics and the Na-insertion/extraction reactions have much poorer reaction kinetics than that of lithium ions.3-7 Compared to Li+ ions, the insertion of Na+ ions requires larger sites for Na+ ions in the matrix and it will create larger distortion to the host lattice. Consequently, crystalline host materials used in LIBs might not be suitable for high-power SIBs and/or their cycling performances will be poor. Recently, amorphous carbon and amorphous phosphorus have been proposed for SIBs and the percolation pathways in amorphous materials might facilitate the diffusion of sodium ions.8-10 However, amorphous carbon and amorphous phosphorus have issues of low capacity and poor conductivity, respectively. Amorphous bimetallic nanoalloys with good conductivity, strain tolerance and large sites for the insertion of Na+ ions are desirable electrode materials for superior SIBs. Amorphous 2 ACS Paragon Plus Environment

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bimetallic nanoalloys as electrodes for SIBs have never been reported so far based on the best of our understanding. We noted that crystalline bimetallic nanoalloys have been investigated for both LIBs and SIBs before.11-13 Given that a large number of bimetallic nanoalloys are available, we will focus on amorphous Sn-Co as an example to illustrate and test our hypothesis that “amorphous is better than crystalline bimetallic nanoalloys as anodes for SIBs”. The selection of Sn is justified due to its high theoretical capacity of 850 mAh/g (based on Na15Sn4). However, the bare Sn electrode suffers from the severe pulverization issue due to volumetric expansion and disintegration leading to poor cycling performance (Figure 1a). There are other metals with similar issues, including Sb at 660 mAh/g (Na3Sb). The benefits of amorphous bimetallic nanoalloys are that the uniform distribution of the electrically conducting M1 (e.g., Co) nanoparticles in close proximity to the high-capacity M2 (e.g., Sn) storage compound could contribute to electrical integration in the electrodes. At the same time M1 could also function as an inactive matrix to inhibit aggregation and size increase the M2 phase to overcome the electrode pulverization issue (Figure 1b). In order words, this combination will offer highly stable electrodes with high capacity for SIBs. The two selected couples of sodium active-inert and active-active systems will provide rational assessment on the practical applications of amorphous bimetallic nanoalloys for high-performance SIBs. The selection of Co is justified due to that it can alloy with Sn, and can function as a good electrically conducting host and an inactive matrix to inhibit aggregation and size increase of Sn phase. It helps to subsequently overcome the electrode pulverization issue of Sn. The sodium active Sn could be finely dispersed into a conducting matrix to achieve homogenous volumetric expansion in the amorphous bimetallic nanoalloys instead of destructive local volume increase in those crystalline ones, or better electrode stability (Figure 1b).



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Figure 1. Illustration to show (a) the pulverization issue observed in crystalline alloys or metals and (b) the enhanced stability of amorphous Sn-Mx nanoalloy electrodes. Amorphous bimetallic nanoalloys have been tested in LIBs but rare for SIBs.14-17 It is particularly interesting to highlight that, using Sn–Co/C as anode materials instead of carbon, Sony Cooperation has already commercialized new-generation LIBs with the trade name of NexelionTM in 2005. Although the composition details were fully not disclosed by Sony, Whittingham at al. analyzed those commercial Sony batteries and concluded that the metallic SnCo phase in NexclionTM is amorphous.15 However, it is challenging to prepare those amorphous Sn-Co based nanoscale materials. Dahn et al. tried to synthesize Sn-Co-C alloys by sputter deposition for high-performance LIBs.14 Xie et al. reported the preparation of Sn-Co nanoalloys by solvothermal route for LIBs.13 It will be interesting to explore and systematically investigate amorphous and crystalline bimetallic Sn-Co nanoalloys for SIBs. Here, we provide the direct experimental evidences to show that amorphous bimetallic nanoalloys are superior over their crystalline counterparts as negative electrodes in SIBs. The feasibility was demonstrated by using a model system of amorphous nanoalloys of Sn-Co 4 ACS Paragon Plus Environment

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prepared by controlled co-reduction as the starting materials. Theoretical gravimetric capacity of Co3Sn2 is 487 mAh/g, which is higher than that of commercial graphite anodes in lithium-ion batteries using graphite with a practical value of 350 mAh/g. More importantly, the theoretical volumetric capacity of Co3Sn2 is 3794 mAh/cm3 based on theoretical density of 7.79 g/cm3. Therefore, the active-inactive and amorphous nanoparticle approach could be further optimized; the Co-Sn couples could be promising anode mateirals for SIBs. The amorphous Sn-Co nanoalloys were then facilely converted to crystalline Sn-Co with different degree of crystallinity by calcination in argon at different temperature. The cycling and rate performances of the amorphous, less-crystalline and crystalline Sn-Co nanoalloys were directly compared in terms of electrochemical sodium ion storage. Amorphous nanoalloys of Sn-Co demonstrated significantly better performances as compared to those crystalline ones as negative electrode materials for SIBs. Experimental Section Materials preparation. (1) Preparation of amorphous Co3Sn2 was based on a modified approach reported.13 In a typical procedure, 1.2 mmol of CoCl2·6H2O and 0.8 mmol of SnCl2·2H2O were fully dissolved in 30 ml of pure ethanol and then transferred to a 50 ml autoclave. 12 mmol of NaBH4 (excessive amount) was slowly added into the autoclave under stirring. The autoclave was sealed in a Teflon-lined autoclave and heated to 180 oC maintained for 24 h and then cooled down naturally. The black precipitate was collected, washed repeatedly with water and ethanol and dried in a vacuum oven at 50 oC overnight. (2) Crystalline Co3Sn2 with different degree of crystallinities were prepared by calcinating the as-obtained amorphous Co3Sn2 from (1) under the protection of argon for 2 h at 500 oC or 600 oC, with ramping rate of 10 oC/min.



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Materials characterization. Powder X-ray diffraction (XRD) was carried out with a Rigaku Smartlab X-ray diffractometer using Cu Kα radiation. Energy-dispersive X-ray spectroscopy (EDS) and morphology characterization were carried out on a field emission scanning electron microscopy (FESEM, JSM-7600, equipped with Pegasus Apex 2 integrated EDS) with accelerating voltage of 15 kV. TEM images were taken by transmission electron microscopy (TEM, JEOL 2010) with operating voltage of 200 kV. Electrochemical measurement. The as-prepared Co3Sn2 samples were mixed with carbon conductivity enhancer (Super-P carbon black, Timcal), and polyvinylidene fluoride (PVDF) binder in weigh ration of 70:20:10 in N-methylpyrrolidone (NMP) solvent to form a slurry. The slurry was pasted on copper discs and dried in vacuum oven at 100 oC overnight. The asprepared material-coated copper discs were assembled into 2032-type coin cells as the working electrodes, metallic Na discs as the counter electrode, 1 M solution of NaClO4 in a 50:50 v/v mixture of ethylene carbonate (EC) and diethyl carbonate (DEC) as the electrolyte, and Celgard 3501 membrane as separator. The as-assembled coin cell was tested at room temperature in a voltage window of 0.02 - 1.2 V with an initial current of 10 mA/g, on a MTI BST8-WA battery tester. To demonstrate the rate performance, the testing current was set to 10 to 20, 50, 100, 150 mA/g and came back to 10 mA/g for the interval of 10 cycles each. Results and Discussion A series of electrode materials with the same chemical composition of Co3Sn2 nanoalloys but different degree of crystallinity were synthesized and systematic investigated to reveal the relationship between their degrees of crystallinity and the corresponding electrochemical performances in reversible sodium ion storage. The different degrees of crystallinity were characterized by XRD (Figure 2). The Co3Sn2 nanoalloys prepared without any calcination are


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amorphous and no detectable XRD peaks could be observed (Figure 2a). Both metallic Sn and Co obtained under the similar conditions had strong and sharp XRD diffraction peaks (Figure S1 in SI), which rules out the formation of metallic Sn or Co. XRD analysis also ruled out the presence of other alloys, including CoSn, CoSn2, CoSn3, which all have totally different XRD patterns from that of Co3Sn2.18 EDS analysis revealed the atomic ratio between Co:Sn is 3:2 (Figure S4 in SI), which agree with the concentration of precursors added. The application of calcination temperature to control the crystallite size is well established and governed by the Scott equation. The temperature induced crystallite growth could be based on means of interfacial interaction/reaction during calcination.19-21 After treatment at 500oC under the protection of argon, the amorphous Co3Sn2 nanoalloys were converted to poorly crystalline nanoalloys with two broad diffraction peaks (Figure 2b). After treatment at 600oC under argon, the amorphous precursor was converted to crystalline Co3Sn2 nanoalloys with distinguishable and sharp diffraction peaks, identified as Co3 Sn2 (Figure 2c).



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Figure 2. XRD patterns for amorphous Sn-Co nanoalloys and the corresponding calcinated samples with different degrees of crystallinity: (a) amorphous Co3Sn2 nanoalloys prepared by coreduction in anhydrous ethanol without calcination; (b) low-degree crystalline Co3Sn2 nanoalloys obtained by calcinating the amorphous precursor (a) at 500oC for 2h; and (c) high-degree crystalline Co3Sn2 nanoalloys obtained by calcinating amorphous precursor (a) at 600oC for 2h. The amorphous Co3Sn2 nanoalloys and the corresponding calcinated samples with different degrees of crystallinity with the same chemical composition were further confirmed by TGADSC (Figure 3). The mass loss observed up to 200oC can be attributed to the removal of chemical and physically attached water, ethanol, and HO- and C2H5O- groups (Figure 3a). Without any heat treatment, the amorphous Sn-Co nanoalloys show a broad exothermal peak between 250-450oC (Figure 3a), which could be attributed to the short-range ordering of the Co


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and Sn atoms in the initially amorphous Co-Sn grains.22 The intensity of the exothermal peak dramatically reduced for the poorly crystalline Sn-Co nanoalloys (Figure 3b) and almost disappeared in crystalline Sn-Co nanoalloys (Figure 3c), as expected. The mass increase from ~400 to 750oC was mainly due to the oxygen adsorption and oxidation of metallic alloy which is associated with the overall heat generation profiles. The exothermal peak at ~650oC could be attributed to the oxidation of the bimetallic nanoalloys into SnO2 and Co3O4 nanoparticles. Further heat treatment could decompose Co3O4 into CoO by thermal reduction of Co3+ into Co2+ (Co3O4 3CoO + 0.5O2) as evidenced in all of the three samples where mass losses were observed at about 750oC. It is also interesting to highlight that the conversion from Co3O4 into CoO occurred at much lower temperature than typically reported at above 915oC.23 The lower temperature of thermal reduction of Co3O4 into CoO suggests that the incorporation of SnO2 could dramatically reduce the activation energy of the thermal reduction reaction by catalytic effects. This interesting observation is worthy for further investigation. The TGA analysis, compliment with the XRD results, further confirms that the three samples have the same chemical composition of Co3Sn2 with different degree of crystallinity.



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Figure 3. TGA-DSC analysis of amorphous Co-Sn nanoalloys and the corresponding calcinated samples with different degrees of crystallinity: (a) amorphous Co3Sn2 nanoalloys


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prepared by co-reduction in anhydrous ethanol without calcination with typical exothermal peaks associated with crystallization; (b) poorly crystalline Co3Sn2 nanoalloys obtained by calcinating the amorphous precursor (a) at 500oC for 2h ; and (c) high-degree crystalline Co3Sn2 nanoalloys obtained by calcinating amorphous precursor (a) at 600oC for 2h. The three samples of bimetallic Co-Sn nanoalloys of from amorphous, poorly crystalline to crystalline ones were also characterized by TEM and SAED (Figure 4). TEM images clearly shows that all the bimetallic samples are nanoscale particles. The amorphous Sn-Co nanoparticles are in the size of about ~10 nm (Figure 4a&b), and the poorly crystalline Sn-Co particles are about ~15 nm (Figure 4c&d), and the crystalline nanoparticles are about ~40 nm in size ((Figure 4e&f). The growth of particle size is associated with the improvement in crystallinity, which could be attributed to annealing induced interfacial aggregation during the crystallization. The different degrees of crystallinity among the three samples were directly evidenced by the different selected area electron diffraction (SAED) patterns. The amorphous Sn-Co nanoalloys didn’t reveal any measurable diffraction rings or spots, confirming the amorphous nature of the as prepared nanoalloys (inset of Figure 4a). For the poorly crystalline Sn-Co nanoalloys obtained by calcinating the amorphous precursor at 500oC, SAED pattern with diffraction rings was observed which clearly indicates the less-crystalline nature (inset of Figure 4c). The crystallinity of the Sn-Co nanoalloys obtained by calcinating the amorphous precursor at 600oC was dramatically improved. The corresponding SAED pattern reveals both diffraction rings and spots indicating crystalline state (inset of Figure 4e). The TEM and SAED results agree with the XRD and TGA analysis aforementioned. Therefore, the three samples with different degrees of crystallinity provide the ideal models to directly test the hypothesis that amorphous



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bimetallic nanoalloys could perform better than their counterparts of crystalline nanoalloys in reversible electrochemical sodium storage.

Figure 4. TEM and SAED characterization of the amorphous bimetallic Co-Sn nanoalloys and the corresponding calcinated samples with different degree of crystallinity: (a, b) low and highmagnification TEM images of the amorphous Co3Sn2 nanoalloys; the inset of (a) is the corresponding SAED pattern without any diffraction observed indicating amorphous nature; (c, d) low and high-magnification TEM images of the poorly crystalline Co3Sn2 nanoalloys obtained by calcinating the amorphous precursor (a) at 500oC; the inset of (c) is the corresponding SAED pattern with diffraction rings indicating less-crystalline nature; (e, f) low and high-magnification TEM images of the crystalline Co3Sn2 nanoalloys obtained by calcinating the amorphous


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precursor at 600oC; the inset of (e) is the corresponding SAED pattern with both diffraction rings and spots indicating crystalline state.

Their electrochemical performances in reversible sodium-ion storage were investigated and compared (Figure 5). The typical first two-cycle charge-discharge profiles for the amorphous, less-crystalline and crystalline Co3Sn2 nanoalloys are directly compared (Figure 5a). The first cycle charge and discharge capacities for the amorphous, less-crystalline and crystalline Co3Sn2 nanoalloys are 1100 and 149 mAh/g, 500 and 98 mAh/g, and 230 and 68 mAh/g, respectively. The huge difference between the amorphous nanoalloys and those crystalline nanoaloys shows the significant effect of degree of crystallinity on sodium storage. Amorphous Co3Sn2 nanoalloys demonstrated much higher first cycle capacities than those of crystalline ones. For the bimetallic nanoalloys, with the increase of degree of crystallinity, the sodium-ion storage capacity decreases dramatically, as evidenced experimentally. This phenomena has been observed in LIBs before.24 The better performance of amorphous Co3Sn2 nanoalloys, as compared to its corresponding calcinated crystalline ones, could be attributed to (1) better accessibility for sodium ions, (2) larger strain accommodation without pulverization, (3) more defect sites that may host sodium ions.25 It is estimated that the first cycle irreversible capacity losses are 86% (with amorphous Sn-Co nanoalloys), 80% (with less-crystalline Sn-Co nanoalloys), and 70% (with crystalline Sn-Co nanoalloys). The formation of solid electrolyte interphase (SEI) could have contributed to the huge irreversible capacity losses observed. Additionally, sodium ions could be trapped at dangling bonds, vacancies and voids, and parasitic reactions involving trace amount of water adsorbed. It is also interesting to highlight that the constant-voltage plateau at near ~1 V observed in discharge (sodiation), which can be associated with the alloying between



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Sn and Na, is much larger at ~ 800 mAh/g for amorphous Co3Sn2 nanoalloys than that of lesscrystalline Sn-Co nanoalloys at 100 mAh/g or that of crystalline Sn-Co nanoalloys with bare plateau observed. In other words, amorphous bimetallic nanoalloys, not crystalline ones, should be considered in the future for SIBs.

Figure 5. Preliminary electrochemical results of the amorphous bimetallic Co3Sn2 nanoalloys and the corresponding calcinated samples with different degree of crystallinity tested for reversible sodium storage: (a) comparison of the first two cycles of charge-discharge profiles for the amorphous (No Cal), less-crystalline (500oC) and crystalline (600oC) Co3Sn2 nanoalloys, from the top to the bottom, respectively; (b) comparison of the rate and cycling performances of the amorphous, less-crystalline, and crystalline Co3Sn2 nanoalloys in reversible sodium storage. Similarly, the cycling and rate performances of the amorphous, less-crystalline and crystalline Co3Sn2 nanoalloys were directly compared (Figure 5b). As expected, the amorphous Co3Sn2 nanoalloys demonstrated significantly better cycling and rate performances at different testing conditions than that of corresponding calcinated crystalline counterparts. As revealed from charge-discharge profiles, with the increase in degree of crystallinity controlled by calcination temperature, the range of plateau associated with forming NaxSn phase becomes shorter leading 14 ACS Paragon Plus Environment

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to decrease in sodium ion storage capacity. The dominant slopes observed in charge-discharge profiles in crystalline Co3Sn2 nanoalloys obtained by calcinations at 600oC indicate that the pseudocapacity is the main capacity contributing factor. The degree of capacitive effect can be determined through the relationship between measured current (i) and scan rate (ν) from the CV curves by measuring b value in the equation of i=υb, where i is the measured current and υ is the scan rate.26 The CV analysis of the crystalline Co3Sn2 at different sweep rates is shown in Figure S4 in SI. It is known that if the b value is close to 0.5, the charge storage mainly comes from intercalation; if the b value is close to 1, the contribution from capacitive effect is dominant.26-27 The b value measured in our case is 0.82. Therefore, the CV analysis also supports the contribution of capacitive effect. Given that the pseudocapcity occurs mainly on and near the surface, the crystalline Co3Sn2 nanoalloys demonstrated very low capacities, as expected. The amorphous Co3Sn2 nanoalloys demonstrated a very good cycling stability at C/5, C/2, 1C, and 1.5C rates, assuming fully charged in 1 h at rate of 100 mA/g. The good rate performance could be attributed to its unique amorphous bimetallic structure with Co as the facilitating conductive matrix to support the sodium active Sn. The bimetallic nanoalloys allow fast diffusion and transport of sodium ions in electrodes. When the current rate was reassumed to C/10, the capacity also reassumed back to about ~100 mAh/g, or 95% as compared to the 10th cycle. The preliminary results clearly indicate that amorphous bimetallic nanoalloys with disordered structures could be potential good electrodes for SIBs, which is worth for further investigation. Another important issue that require further study is the associated huge first cycle irreversible capacity loss, which could consume too much expensive coupled cathode materials in full cells. In fact, the first cycle irreversible capacity loss is a common issue observed in many anode materials for sodium-ion batteries, including carbon blace.10 Future studies are needed to



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mitigate the first cycle irreversible capacity loss problem. The possible strategies proposed include (1) pre-sodiation, (2) coating the active materials with protecting layers with good stability, (3) selection of suitable electrolytes with less SEI formation, and (4) use of electrolyte stabilizers. We are working in those directions in our next stage of investigation. Conclusions In summary, our experimental results clearly show that bimetallic active-inactive nanoalloys, using Co-Sn as a model, in the form of amorphous phase is superior over their crystalline counterparts in SIBs. The improved electrochemical performances of amorphous Co-Sn nanoalloys could be attributed to the easy accessibility for sodium ions, strain accommodation and the presence of defect sites that may host sodium ions. We anticipate that the Sony’s formula of Sn-Co/C anode in NexelionTM LIBs could be good anode for SIBs as well.

ASSOCIATED CONTENT Supporting Information. XRD, TEM and battery performances of similarly obtained Sn and Co particles and CV plots. This information is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author * Da Deng Address: 5050 Anthony Wayne Dr., Room 3164, Detroit, MI 48202; Telephone: 313-577-5940; Email: [email protected] Author Contributions 16 ACS Paragon Plus Environment

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18. He, J.; Zhao, H.; Wang, J.; Wang, J.; Chen, J. Hydrothermal Synthesis and Electrochemical Properties of Nano-Sized Co–Sn Alloy Anodes for Lithium Ion Batteries. J. Alloys Compd. 2010, 508, 629-635. 19. Yang, H. M.; Zhang, X. C.; Ao, W. Q.; Qiu, G. Z. Formation of Nife2o4 Nanoparticles by Mechanochemical Reaction. Mater. Res. Bull. 2004, 39, 833-837. 20. Ao, W. Q.; Li, J. Q.; Yang, H. M.; Zeng, X. R.; Ma, X. C. Mechanochemical Synthesis of Zinc Oxide Nanocrystalline. Powder Technol. 2006, 168, 148-151. 21. Hayat, K.; Gondal, M. A.; Khaled, M. M.; Ahmed, S.; Shemsi, A. M. Nano Zno Synthesis by Modified Sol Gel Method and Its Application in Heterogeneous Photocatalytic Removal of Phenol from Water. Appl. Catal., A 2011, 393, 122-129. 22. Ferguson, P. P.; Rajora, M.; Dunlap, R. A.; Dahn, J. R. ( Sn0.5co0.5 ) 1 − Y C Y Alloy Negative Electrode Materials Prepared by Mechanical Attriting. J. Electrochem. Soc. 2009, 156, A204-A208. 23. Liu, H.-C.; Yen, S.-K. Characterization of Electrolytic Co3o4 Thin Films as Anodes for Lithium-Ion Batteries. J. Power Sources 2007, 166, 478-484. 24. Deng, D.; Kim, M. G.; Lee, J. Y.; Cho, J. Green Energy Storage Materials: Nanostructured Tio2 and Sn-Based Anodes for Lithium-Ion Batteries. Energy Environ. Sci. 2009, 2, 818-837. 25. Liu, B.; Deng, D.; Lee, J. Y.; Aydil, E. S. Oriented Single-Crystalline Tio2 Nanowires on Titanium Foil for Lithium Ion Batteries. J. Mater. Res. 2010, 25, 1588-1594. 26. Wang, J.; Polleux, J.; Lim, J.; Dunn, B. Pseudocapacitive Contributions to Electrochemical Energy Storage in Tio2 (Anatase) Nanoparticles. J. Phys. Chem. C 2007, 111, 14925-14931. 27. Yu, P.; Li, C.; Guo, X. Sodium Storage and Pseudocapacitive Charge in Textured Li4ti5o12 Thin Films. J. Phys. Chem. C 2014, 118, 10616-10624.


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Table of Contents Graphic


Amorphous Bimetallic Co3Sn2 Nanoalloys Are Better Than Crystalline

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