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Performance enhancement of Sn–Co alloys for lithium-ion battery by electrochemical dissolution treatment
Journal of Alloys and Compounds 574 (2013) 206–211
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Performance enhancement of Sn–Co alloys for lithium-ion battery by electrochemical dissolution treatment Chunhui Tan a, Gongwei Qi a, Yeping Li a, Jing Guo a, Xin Wang a, Delong Kong b, Hongjun Wang b, Shuyong Zhang a,⇑ a b
Key Laboratory of Colloid and Interface Chemistry, Ministry of Education, School of Chemistry and Chemical Engineering, Shandong University, Jinan 250100, China Shandong Sacred Sun Power Co., Ltd., Qufu 273100, China
a r t i c l e
i n f o
Article history: Received 25 November 2012 Received in revised form 18 March 2013 Accepted 25 March 2013 Available online 8 April 2013 Keywords: Lithium-ion battery Anode materials Sn–Co alloy Porous structure Electrochemical dissolution treatment
a b s t r a c t Three sets of Sn–Co alloy materials with improved porous structure are designed and facilely prepared via a two-step strategy comprising of electrodeposition followed by electrochemical dissolution treatment. The structure, composition and morphology of the Sn–Co alloy materials are investigated by scanning electron microscopy, X-ray diffraction and atomic absorption spectrophotometer. The electrochemical performance of the Sn–Co alloy materials as negative electrode materials for lithium-ion battery was tested by cyclic voltammetry, galvanostatic charge–discharge cycling and rate capability test. It is found that the resultant Sn–Co alloy materials after electrochemical dissolution treatment have porous structure and higher specific capacity than those of the as-deposited Sn–Co alloy materials. The initial capacity of a porous Sn–Co alloy sample attains 1109 mA h g1, and remains 580 mA h g1 after 70 cycles. This performance improvement can be mainly attributed to the formation of optimized porous structure during electrochemical dissolution treatment, which can buffer the great volume fluctuation of Sn phase during cycling and therefore reduce cracking and shelling of the materials. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Rapid development in portable electronic products and electric vehicles in recent years have brought about great progresses in lithium-ion battery (LIB). However, there are still many remained challenges. At present, most negative electrode materials for LIB are made of carbonaceous materials with low theoretical specific capacity, which limits the possibility for further improvement [1]. In order to enhance the safety, specific capacity and capacity retention of the negative electrode materials, researchers turned their attention to several materials with high-capacity, such as Sn, Si, and Sb [2,3]. The theoretical capacity of Sn is as higher as 990 mA h g1, about twice that of common carbonaceous materials [2]. However, a fatal defect of Sn-based materials is their large volume fluctuation during the intercalation–deintercalation of lithium ions, which results in a rapid destruction of the electrode structure and poor capacity retention [2,4–6]. Thus, improving the structural stability of the Sn-based materials has become one of the hottest topics in LIB research. Many approaches such as introduction of inert alloy component, reduction of particle size and creation of porous structure, have been tested. Introduction ⇑ Corresponding author. Tel.: +86 531 88361378; fax: +86 531 88364781. E-mail addresses: [email protected] (C. Tan), [email protected] (S. Zhang). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.291
of inert elements to form intermetallic compounds or composite alloys is a commonly adopted approach. Many kinds of alloy materials including Sn–Cu [7–11], Sn–Sb [12], Sn–Co [13–16], Sn–Ni [17–20], and Sn–Zn [21], have been studied. These alloys show improved cycling performance due to formation of an inert conductive network in their matrix. Reducing the particle size to submicron or even nanometer scales is the second way to improve the electrochemical performance of Sn-based material [7,17, 22–26]. However, small particles tend to aggregate, rapidly worsening their cycling performance [25]. Aside from large specific surface area, porous materials usually show better electric conductivity and extra room for volume change during charge–discharge cycling. Thus, porous electrode materials have received considerable attention [9,13,27–31]. Shin and Liu prepared a 3D porous Cu-Sn alloy material via electrochemical deposition using hydrogen bubbles as a template [9]. Ke et al. prepared a series of porous Sn alloys with a colloidal template [13,19,32]. Copper foam was also used as a framework for preparation of porous materials with improved performance by Huang et al. [27,33–35]. Yu et al. used 3D nanoporous Au-supported Sn as negative electrode materials for LIB and found that the capacity of this material remained 600 mA h g1 after 140 cycles [36]. In this paper, inert element Co was introduced and porous structure was built by electrochemical dissolution. Porous Sn–Co
C. Tan et al. / Journal of Alloys and Compounds 574 (2013) 206–211 Table 1 Bath composition for electrodeposition of different Sn–Co alloys.
207
3. Results and discussion
Baths
A/g L1
B/g L1
C/g L1
K4P2O73H2O Potassium sodium tartrate (C4H4KNaO64H2O) Glycine (C2H5NO2) SnCl22H2O CoCl26H2O Gelatin
192.2 10 10 33.8 11.9 1
192.2 10 10 28.2 17.8 1
192.2 10 10 22.6 23.8 1
alloy materials were fabricated by a two-step strategy: electrodeposition of Sn–Co alloy onto a piece of treated Cu foil substrate followed by electrochemical dissolution of part of active material. Using this strategy, a binder-free Sn-based alloy electrode with optimized porous structure is fabricated and improved electrochemical performance is observed.
2. Experimental 2.1. Preparations of porous Sn–Co alloy electrode Porous Sn–Co alloys loaded on a piece of commercial Cu foil substrate (10 lm in thickness) were prepared using a two-step strategy. Before electrodeposition, the Cu foil was treated in acetone and 0.1 mol L1 HCl to remove surface contamination and oxides. Sn–Co alloy was then deposited on one side of the treated Cu foil substrate in a home-made two-electrode plating cell without stirring. The Cu foil served as the cathode, a Ti foil as the anode and three different baths shown in Table 1 were used as electrolyte so as to obtain alloys with different Sn and Co contents. The area of the Cu cathode is 10 cm2. Electrodeposition was conducted galvanostatically at a current density of 10 mA cm2 for 10 min. After deposition, electrochemical dissolution was conducted in the same cell with the Sn–Co alloy loaded on the Cu foil serving as the anode and a Ni net serving as the cathode in a 0.1 mol L1 hydrochloric acid solution for 4 min at a current density of 10 mA cm2. Both electrodeposition and electrochemical dissolution were carried out on an HDV-7 potentiostat (Sanming, China). After electrodeposition and electrochemical dissolution, the Sn–Co alloy samples were washed with deionized water and absolute ethanol, and then air-dried. The Sn–Co alloys before and after electrochemical dissolution were marked with 1 and 2, respectively. For example, A1 and A2 correspond to the Sn–Co alloy electrodeposited from Bath A before and after electrochemical dissolution, respectively. The mass of the active materials was obtained from the weight change of the Cu foil substrate before and after electrodeposition and electrochemical dissolution. All the reagents used in the experiments were purchased from Sinopharm Chemical Reagent Co. Ltd. and used without further purification.
3.1. Morphology and structure of Sn–Co alloys The XRD patterns of the Sn–Co alloys before and after electrochemical dissolution treatment are shown in Fig. 1. In Fig. 1, presence of tetragonal Sn phase in the alloys (JCPDS Card No. 04-0673, space group: I41/amd, 141) is evidenced by diffraction peaks appearing at 2h = 30.6°, 32.0°, 43.9°, 44.9°, 55.3°, 62.5°, 63.8°, 64.6°, 72.4°, 73.2°, and 79.5°. Cubic Cu phase (JCPDS Card No. 04-0836, space group: Fm3m, 225) is evidenced by peaks appearing at 43.3°, 50.4° and 74.1°. Tetragonal SnCo intermetallic phase is evidenced by peaks appearing at 2h = 28.9°, 32.8°, 41.3°, 42.5°, 44.3°, 52.2°, 55.7°, 59.9°, 67.8° and 76.2°. The SnCo phase does not appear in the thermal phase diagram of the Sn–Co bimetallic system and the information of this phase is introduced in the TREOR indexation program [37]. Therefore, the XRD pattern demonstrates the presence of Sn, Cu and intermetallic SnCo phases in all these alloys. It can be seen from the XRD patterns of A1, B1 and C1 that the diffraction peaks belong to Sn and SnCo phases become weaker with the increasing Co content. It means that, the alloy becomes more amorphous in this case. The diffraction peaks belonging to Sn phase can be observed in A1, B1 and C1 but hardly be seen in A2, B2 and C2. It suggests that most of the Sn phase dissolves during electrochemical dissolution treatment. The diffraction peaks of the intermetallic SnCo phase can be observed both before and after electrochemical dissolution treatment, which shows that SnCo phase is more stable than Sn phase. The diffraction peaks of Cu phase become stronger after electrochemical dissolution treatment, suggesting reduced thickness of the porous alloy deposition layers. The composition of the alloy before and after electrochemical dissolution treatment is determined using AAS with the results listed in Table 2. It suggests that the Sn and Co content can be easily controlled by changing the main salt concentrations in the bath. After electrochemical dissolution treatment, there are more than 65% Sn and less than 52% Co has lost. So the Co content in the residual alloy increased. The morphologies of the alloy before and after electrochemical dissolution treatment are shown in Fig. 2. Fig. 2 demonstrates that the electrode surface is mainly composed of submicron particles. The size of the particle reduces with increasing Co content in the alloy as evidenced also by XRD results. After electrochemical dissolution treatment, some hollow spheres
2.2. Cell assembling Test CR2025 coin cells were assembled with the Sn–Co alloy loaded on the Cu foil substrate (ca. 0.5 cm2) serving as the working electrode, a lithium sheet as the counter electrode and reference electrode in an argon-filled glove box at room temperature. The electrolyte was 1 mol L1 LiPF6 in a mixed solvent of ethylene carbonate/dimethyl carbonate/ethyl methyl carbonate in 1:1:1 (vol.%), provided by Beijing Institute of Chemical Reagents. The cathode and anode were separated by a Celgard 2400 membrane.
2.3. Structure analysis and electrochemical tests The morphology of the Sn–Co alloy electrode before and after electrochemical dissolution was observed using a JSM-7600F Field emission scanning electron microscope (SEM, JEOL, Japan). The composition of the alloy was analyzed using a 3510 Atomic absorption spectrophotometer (AAS, HP, USA). The crystallographic information of the sample was recorded using a D8 Advance X-ray diffractometer (XRD, Bruker, Germany) with Cu Ka radiation at a scan rate of 0.5°s1. Cyclic voltammetry (CV) was performed using a CHI604A Electrochemical workstation (CH Instrument, Shanghai, China) at a scan rate of 0.5° mV s1 within the potential range of 0–2 V (vs. Li/Li+). The charge and discharge cycling tests were performed using a CT2001C-001 Land battery testing system (Jinnuo, Wuhan, China) with a current density of 198 mA g1 (0.2 C rate) at a cut-off voltage of 0.05– 1.5 V (vs. Li/Li+).
Fig. 1. XRD patterns of the Sn–Co alloys before and after electrochemical dissolution treatment.
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C. Tan et al. / Journal of Alloys and Compounds 574 (2013) 206–211
Table 2 Composition of the Sn–Co alloys before and after electrochemical dissolution. Alloy
A1 A2 B1 B2 C1 C2
Mass (mg)
Molar Ratio (%)
Sn
Co
Sn
Co
30.1 8.79 22.4 6.82 16.3 5.68
1.03 0.96 1.81 1.00 2.27 1.10
93.6 82.0 86.0 77.2 78.1 71.9
6.40 18.0 14.0 22.8 21.9 28.1
can be observed, which can be related to the Sn dissolution as we have learnt from Table 2. Formation of hollow sphere also suggests that some particles are of core–shell structure with the shell mainly composed of intermetallic SnCo. The structure of these Sn–Co alloys is much similar to that of the porous Ni–Cu alloy reported elsewhere [38]. Owing to the less particle size and hollow spherical structure, the specific area of the samples after electrochemical dissolution treatment is expected to be much higher. 3.2. Electrochemical performance of Sn–Co alloys According to Fig. 3, the initial specific capacity of Sn–Co alloys before electrochemical dissolution treatment is 905 mA h g1 for A1, 759 mA h g1 for B1, and 706 mA h g1 for C1, respectively. For these three alloys, the different initial capacities are related to the Sn content in the deposited alloys. After 70 cycles, the specific capacity of them decays to 127 mA h g1, 215 mA h g1 and 414 mA h g1 respectively. It is clear that, like other researchers reported, the capacity retention of the materials gets improved but capacities of the materials get reduced as the increasing of Co
content in the alloy. The initial specific capacity of Sn–Co alloys after electrochemical dissolution treatment is 967 mA h g1 for A2, 1072 mA h g1 for B2 and 1109 mA h g1 for C2. The initial capacity of them is much higher and coulombic efficiency is lower than A1, B1 and C1. The high initial capacity is due to their higher specific area and more oxides formed on their surface. After 70 cycles, the capacity of A2, B2, and C2 is 248 mA h g1, 389 mA h g1 and 578 mA h g1, respectively. Compared with A1, B1 and C1, their capacity at 70th cycle is increased by 95.3%, 80.9% and 39.6%, respectively. From Fig. 3b, the coulombic efficiency of A2, B2 and C2 is obviously higher than that of A1, B1 and C1 at most of the time and most of coulombic efficiency of B2 and C2 are more than 95%. Therefore, electrochemical dissolution treatment after deposition can greatly improve specific capacity, capacity retention and coulombic efficiency of the Sn–Co alloy. The first five CV curves of the C1and C2 samples are recorded and shown in Fig. 4. In Fig. 4, the reduction current appearing in the first negative scan near 0 V vs. Li+/Li is assigned to the lithiation of Sn–Co alloy and formation of LixSn alloy and isolated Co as shown by the following equation: þ
Coy Sn þ xLi þ xe ! Lix Sn þ yCo ð0 < x < 4:4Þ
ð1Þ
This reaction is irreversible. This means that Co separates from the Sn–Co alloy during the first lithiation. The reduction peak shifted positively to ac. 0.3 V from the second reverse scan which stands for the formation of LiySn (3.5 < y < 4.4). The oxidation peaks appearing at ac. 0.7 V in the positive scan is due to the dealloying reaction from LixSn to Li and Sn as shown by the following equation [2,6]: þ
Lix Sn þ yCo ! Sn þ xLi þ yCo þ xe
Fig. 2. SEM images of the as-deposited Sn–Co alloys and porous Sn–Co alloys.
ð2Þ
C. Tan et al. / Journal of Alloys and Compounds 574 (2013) 206–211
209
Fig. 3. Discharge capacity (a) and Coulombic efficiency (b) of Sn–Co alloys before and after electrochemical dissolution treatment.
Fig. 4. Cyclic voltammograms profiles of C1 (a) and C2 (b).
In which Co remains as an inert component. The increasing discrepancy between the forward scan and reverse scan curves within the potential range of 1.5–2.0 V corresponding to charge of doublelayer capacitance reflects the increased surface area of the Sn–Co alloys after electrochemical dissolution treatment. The peak current of C1 increased with cycling in the first few cycles can be due to the continuous formation of new surface during lithiation and delithiation. For C2, except for the first cycle, the curve of other cycles coincides well, demonstrating improved reversibility. Charge–discharge curves of as-deposited alloy and porous alloy are obviously different. Charge–discharge curves of C1 and C2 were taken for example shown in Fig. 5. The charge–discharge curves of the C1 are similar to that of Sn. The platforms appearing at 0.65 V, 0.51 V and 0.41 V in discharge curves stand for the lithiation of Sn phase and formation of Li0.4Sn, LiSn and Li2.33Sn. The corresponding platforms in charge curves appearing at 0.77 V, 0.71 V and 0.58 V relate to the delithiation and formation of Sn, Li0.4Sn and LiSn [2]. The platform of C1 in the first discharge curve is lower due to resistance polarization. Since significant loss of Sn occur during electrochemical dissolution treatment, C2 have different charge– discharge behaviors with only one obvious platform observed in each curve. The platform appears at ca. 0.3 V in discharge curves and ca. 0.6 V in charge curves, relating to Eqs. (1) and (2), respectively. According to literature, for first discharge, reaction taking place at potential higher than 0.6 V gives irreversible capacity [39,40]. This irreversible capacity is due to decomposition of
electrolyte, reduction of tin oxide and formation of solid electrolyte interface (SEI) on the surface of electrode. C2 have higher irreversible capacity in the first discharge, which can be due to their higher surface area that consumes more capacity for formation of SEI and reduction of tin oxide. Porous structure owns larger surface area and can shorten diffusion distance of Li which will reflect on rate capability. The improved rate capability of porous structure formed from electrodeposition and electrochemical dissolution treatment was tested taking C1 and C2 as an example and shown in Fig. 6a. The charge–discharge rate rose stepwise from 0.2C to 10C and reduced to 0.2C after that. It is very clear that C2 delivered remarkable rate ability with the stable performance and high capacity retention. The 10C capacity of C2 is ca. 400 mA h g1 which is 67% of the 0.2C capacity tested after 10C. The high coulombic efficiency of C2 can be also observed at every rate and never less than 95%. Fig. 6b displayed the charge–discharge curves at different rates which performed identical shape. The C2 delivered 782 mA h g1, 618 mA h g1, 573 mA h g1, 485 mA h g1 and 400 mA h g1 at the rate of 0.2C, 1C, 2C, 5C and 10C respectively. According the above XRD and SEM results, we ascribe the improved performance in specific capacity, capacity retention and rate capability to several changes: (1) loss of Sn phase during electrochemical treatment reduces the volume expansion of the material during lithiation; (2) dissolution of Sn phase leaves many pores behind which buffers the volume expansion during lithiation; and
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Fig. 5. The charge and discharge curves of C1 (a) and C2 (b).
Fig. 6. Rate capabilities of C1 and C2 (a) and charge and discharge curves of C2 in different rate (b).
(3) electrochemical dissolution reduces the size of the alloy particles. All these changes can improve the cycling performance of the materials. And the reduced particle size and formation of hollow sphere increase the area contacting with electrolyte, which promotes the capacity of the materials. 4. Conclusions A new method to improve electrochemical performance of electrochemically deposited Sn–Co alloy materials for lithium-ion battery by subsequent electrochemical dissolution is put forward. During electrochemical dissolution treatment, part of the Sn phase in the Sn–Co alloys was dissolved and the surface structure of the residual alloy was improved. The reduced particle size and formation of hollow sphere promote the specific capacity of the resultant materials. After electrochemical dissolution treatment, specific capacity, capacity retention and rate capability of a Sn–Co alloy are greatly increased which fully shows subsequent dissolution treatment is an effective way to improve the electrochemical performance of Sn–Co alloy materials for lithium-ion batteries. Acknowledgements This work was financially supported by Natural Science Foundation of Shandong Province (Project No. ZR2009BM012) and the 973 Project of China (No. 2011CB935901).
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Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Performance enhancement of Sn–Co alloys for lithium-ion battery by electrochemical dissolution treatment Chunhui Tan a, Gongwei Qi a, Yeping Li a, Jing Guo a, Xin Wang a, Delong Kong b, Hongjun Wang b, Shuyong Zhang a,⇑ a b
Key Laboratory of Colloid and Interface Chemistry, Ministry of Education, School of Chemistry and Chemical Engineering, Shandong University, Jinan 250100, China Shandong Sacred Sun Power Co., Ltd., Qufu 273100, China
a r t i c l e
i n f o
Article history: Received 25 November 2012 Received in revised form 18 March 2013 Accepted 25 March 2013 Available online 8 April 2013 Keywords: Lithium-ion battery Anode materials Sn–Co alloy Porous structure Electrochemical dissolution treatment
a b s t r a c t Three sets of Sn–Co alloy materials with improved porous structure are designed and facilely prepared via a two-step strategy comprising of electrodeposition followed by electrochemical dissolution treatment. The structure, composition and morphology of the Sn–Co alloy materials are investigated by scanning electron microscopy, X-ray diffraction and atomic absorption spectrophotometer. The electrochemical performance of the Sn–Co alloy materials as negative electrode materials for lithium-ion battery was tested by cyclic voltammetry, galvanostatic charge–discharge cycling and rate capability test. It is found that the resultant Sn–Co alloy materials after electrochemical dissolution treatment have porous structure and higher specific capacity than those of the as-deposited Sn–Co alloy materials. The initial capacity of a porous Sn–Co alloy sample attains 1109 mA h g1, and remains 580 mA h g1 after 70 cycles. This performance improvement can be mainly attributed to the formation of optimized porous structure during electrochemical dissolution treatment, which can buffer the great volume fluctuation of Sn phase during cycling and therefore reduce cracking and shelling of the materials. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Rapid development in portable electronic products and electric vehicles in recent years have brought about great progresses in lithium-ion battery (LIB). However, there are still many remained challenges. At present, most negative electrode materials for LIB are made of carbonaceous materials with low theoretical specific capacity, which limits the possibility for further improvement [1]. In order to enhance the safety, specific capacity and capacity retention of the negative electrode materials, researchers turned their attention to several materials with high-capacity, such as Sn, Si, and Sb [2,3]. The theoretical capacity of Sn is as higher as 990 mA h g1, about twice that of common carbonaceous materials [2]. However, a fatal defect of Sn-based materials is their large volume fluctuation during the intercalation–deintercalation of lithium ions, which results in a rapid destruction of the electrode structure and poor capacity retention [2,4–6]. Thus, improving the structural stability of the Sn-based materials has become one of the hottest topics in LIB research. Many approaches such as introduction of inert alloy component, reduction of particle size and creation of porous structure, have been tested. Introduction ⇑ Corresponding author. Tel.: +86 531 88361378; fax: +86 531 88364781. E-mail addresses: [email protected] (C. Tan), [email protected] (S. Zhang). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.291
of inert elements to form intermetallic compounds or composite alloys is a commonly adopted approach. Many kinds of alloy materials including Sn–Cu [7–11], Sn–Sb [12], Sn–Co [13–16], Sn–Ni [17–20], and Sn–Zn [21], have been studied. These alloys show improved cycling performance due to formation of an inert conductive network in their matrix. Reducing the particle size to submicron or even nanometer scales is the second way to improve the electrochemical performance of Sn-based material [7,17, 22–26]. However, small particles tend to aggregate, rapidly worsening their cycling performance [25]. Aside from large specific surface area, porous materials usually show better electric conductivity and extra room for volume change during charge–discharge cycling. Thus, porous electrode materials have received considerable attention [9,13,27–31]. Shin and Liu prepared a 3D porous Cu-Sn alloy material via electrochemical deposition using hydrogen bubbles as a template [9]. Ke et al. prepared a series of porous Sn alloys with a colloidal template [13,19,32]. Copper foam was also used as a framework for preparation of porous materials with improved performance by Huang et al. [27,33–35]. Yu et al. used 3D nanoporous Au-supported Sn as negative electrode materials for LIB and found that the capacity of this material remained 600 mA h g1 after 140 cycles [36]. In this paper, inert element Co was introduced and porous structure was built by electrochemical dissolution. Porous Sn–Co
C. Tan et al. / Journal of Alloys and Compounds 574 (2013) 206–211 Table 1 Bath composition for electrodeposition of different Sn–Co alloys.
207
3. Results and discussion
Baths
A/g L1
B/g L1
C/g L1
K4P2O73H2O Potassium sodium tartrate (C4H4KNaO64H2O) Glycine (C2H5NO2) SnCl22H2O CoCl26H2O Gelatin
192.2 10 10 33.8 11.9 1
192.2 10 10 28.2 17.8 1
192.2 10 10 22.6 23.8 1
alloy materials were fabricated by a two-step strategy: electrodeposition of Sn–Co alloy onto a piece of treated Cu foil substrate followed by electrochemical dissolution of part of active material. Using this strategy, a binder-free Sn-based alloy electrode with optimized porous structure is fabricated and improved electrochemical performance is observed.
2. Experimental 2.1. Preparations of porous Sn–Co alloy electrode Porous Sn–Co alloys loaded on a piece of commercial Cu foil substrate (10 lm in thickness) were prepared using a two-step strategy. Before electrodeposition, the Cu foil was treated in acetone and 0.1 mol L1 HCl to remove surface contamination and oxides. Sn–Co alloy was then deposited on one side of the treated Cu foil substrate in a home-made two-electrode plating cell without stirring. The Cu foil served as the cathode, a Ti foil as the anode and three different baths shown in Table 1 were used as electrolyte so as to obtain alloys with different Sn and Co contents. The area of the Cu cathode is 10 cm2. Electrodeposition was conducted galvanostatically at a current density of 10 mA cm2 for 10 min. After deposition, electrochemical dissolution was conducted in the same cell with the Sn–Co alloy loaded on the Cu foil serving as the anode and a Ni net serving as the cathode in a 0.1 mol L1 hydrochloric acid solution for 4 min at a current density of 10 mA cm2. Both electrodeposition and electrochemical dissolution were carried out on an HDV-7 potentiostat (Sanming, China). After electrodeposition and electrochemical dissolution, the Sn–Co alloy samples were washed with deionized water and absolute ethanol, and then air-dried. The Sn–Co alloys before and after electrochemical dissolution were marked with 1 and 2, respectively. For example, A1 and A2 correspond to the Sn–Co alloy electrodeposited from Bath A before and after electrochemical dissolution, respectively. The mass of the active materials was obtained from the weight change of the Cu foil substrate before and after electrodeposition and electrochemical dissolution. All the reagents used in the experiments were purchased from Sinopharm Chemical Reagent Co. Ltd. and used without further purification.
3.1. Morphology and structure of Sn–Co alloys The XRD patterns of the Sn–Co alloys before and after electrochemical dissolution treatment are shown in Fig. 1. In Fig. 1, presence of tetragonal Sn phase in the alloys (JCPDS Card No. 04-0673, space group: I41/amd, 141) is evidenced by diffraction peaks appearing at 2h = 30.6°, 32.0°, 43.9°, 44.9°, 55.3°, 62.5°, 63.8°, 64.6°, 72.4°, 73.2°, and 79.5°. Cubic Cu phase (JCPDS Card No. 04-0836, space group: Fm3m, 225) is evidenced by peaks appearing at 43.3°, 50.4° and 74.1°. Tetragonal SnCo intermetallic phase is evidenced by peaks appearing at 2h = 28.9°, 32.8°, 41.3°, 42.5°, 44.3°, 52.2°, 55.7°, 59.9°, 67.8° and 76.2°. The SnCo phase does not appear in the thermal phase diagram of the Sn–Co bimetallic system and the information of this phase is introduced in the TREOR indexation program [37]. Therefore, the XRD pattern demonstrates the presence of Sn, Cu and intermetallic SnCo phases in all these alloys. It can be seen from the XRD patterns of A1, B1 and C1 that the diffraction peaks belong to Sn and SnCo phases become weaker with the increasing Co content. It means that, the alloy becomes more amorphous in this case. The diffraction peaks belonging to Sn phase can be observed in A1, B1 and C1 but hardly be seen in A2, B2 and C2. It suggests that most of the Sn phase dissolves during electrochemical dissolution treatment. The diffraction peaks of the intermetallic SnCo phase can be observed both before and after electrochemical dissolution treatment, which shows that SnCo phase is more stable than Sn phase. The diffraction peaks of Cu phase become stronger after electrochemical dissolution treatment, suggesting reduced thickness of the porous alloy deposition layers. The composition of the alloy before and after electrochemical dissolution treatment is determined using AAS with the results listed in Table 2. It suggests that the Sn and Co content can be easily controlled by changing the main salt concentrations in the bath. After electrochemical dissolution treatment, there are more than 65% Sn and less than 52% Co has lost. So the Co content in the residual alloy increased. The morphologies of the alloy before and after electrochemical dissolution treatment are shown in Fig. 2. Fig. 2 demonstrates that the electrode surface is mainly composed of submicron particles. The size of the particle reduces with increasing Co content in the alloy as evidenced also by XRD results. After electrochemical dissolution treatment, some hollow spheres
2.2. Cell assembling Test CR2025 coin cells were assembled with the Sn–Co alloy loaded on the Cu foil substrate (ca. 0.5 cm2) serving as the working electrode, a lithium sheet as the counter electrode and reference electrode in an argon-filled glove box at room temperature. The electrolyte was 1 mol L1 LiPF6 in a mixed solvent of ethylene carbonate/dimethyl carbonate/ethyl methyl carbonate in 1:1:1 (vol.%), provided by Beijing Institute of Chemical Reagents. The cathode and anode were separated by a Celgard 2400 membrane.
2.3. Structure analysis and electrochemical tests The morphology of the Sn–Co alloy electrode before and after electrochemical dissolution was observed using a JSM-7600F Field emission scanning electron microscope (SEM, JEOL, Japan). The composition of the alloy was analyzed using a 3510 Atomic absorption spectrophotometer (AAS, HP, USA). The crystallographic information of the sample was recorded using a D8 Advance X-ray diffractometer (XRD, Bruker, Germany) with Cu Ka radiation at a scan rate of 0.5°s1. Cyclic voltammetry (CV) was performed using a CHI604A Electrochemical workstation (CH Instrument, Shanghai, China) at a scan rate of 0.5° mV s1 within the potential range of 0–2 V (vs. Li/Li+). The charge and discharge cycling tests were performed using a CT2001C-001 Land battery testing system (Jinnuo, Wuhan, China) with a current density of 198 mA g1 (0.2 C rate) at a cut-off voltage of 0.05– 1.5 V (vs. Li/Li+).
Fig. 1. XRD patterns of the Sn–Co alloys before and after electrochemical dissolution treatment.
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Table 2 Composition of the Sn–Co alloys before and after electrochemical dissolution. Alloy
A1 A2 B1 B2 C1 C2
Mass (mg)
Molar Ratio (%)
Sn
Co
Sn
Co
30.1 8.79 22.4 6.82 16.3 5.68
1.03 0.96 1.81 1.00 2.27 1.10
93.6 82.0 86.0 77.2 78.1 71.9
6.40 18.0 14.0 22.8 21.9 28.1
can be observed, which can be related to the Sn dissolution as we have learnt from Table 2. Formation of hollow sphere also suggests that some particles are of core–shell structure with the shell mainly composed of intermetallic SnCo. The structure of these Sn–Co alloys is much similar to that of the porous Ni–Cu alloy reported elsewhere [38]. Owing to the less particle size and hollow spherical structure, the specific area of the samples after electrochemical dissolution treatment is expected to be much higher. 3.2. Electrochemical performance of Sn–Co alloys According to Fig. 3, the initial specific capacity of Sn–Co alloys before electrochemical dissolution treatment is 905 mA h g1 for A1, 759 mA h g1 for B1, and 706 mA h g1 for C1, respectively. For these three alloys, the different initial capacities are related to the Sn content in the deposited alloys. After 70 cycles, the specific capacity of them decays to 127 mA h g1, 215 mA h g1 and 414 mA h g1 respectively. It is clear that, like other researchers reported, the capacity retention of the materials gets improved but capacities of the materials get reduced as the increasing of Co
content in the alloy. The initial specific capacity of Sn–Co alloys after electrochemical dissolution treatment is 967 mA h g1 for A2, 1072 mA h g1 for B2 and 1109 mA h g1 for C2. The initial capacity of them is much higher and coulombic efficiency is lower than A1, B1 and C1. The high initial capacity is due to their higher specific area and more oxides formed on their surface. After 70 cycles, the capacity of A2, B2, and C2 is 248 mA h g1, 389 mA h g1 and 578 mA h g1, respectively. Compared with A1, B1 and C1, their capacity at 70th cycle is increased by 95.3%, 80.9% and 39.6%, respectively. From Fig. 3b, the coulombic efficiency of A2, B2 and C2 is obviously higher than that of A1, B1 and C1 at most of the time and most of coulombic efficiency of B2 and C2 are more than 95%. Therefore, electrochemical dissolution treatment after deposition can greatly improve specific capacity, capacity retention and coulombic efficiency of the Sn–Co alloy. The first five CV curves of the C1and C2 samples are recorded and shown in Fig. 4. In Fig. 4, the reduction current appearing in the first negative scan near 0 V vs. Li+/Li is assigned to the lithiation of Sn–Co alloy and formation of LixSn alloy and isolated Co as shown by the following equation: þ
Coy Sn þ xLi þ xe ! Lix Sn þ yCo ð0 < x < 4:4Þ
ð1Þ
This reaction is irreversible. This means that Co separates from the Sn–Co alloy during the first lithiation. The reduction peak shifted positively to ac. 0.3 V from the second reverse scan which stands for the formation of LiySn (3.5 < y < 4.4). The oxidation peaks appearing at ac. 0.7 V in the positive scan is due to the dealloying reaction from LixSn to Li and Sn as shown by the following equation [2,6]: þ
Lix Sn þ yCo ! Sn þ xLi þ yCo þ xe
Fig. 2. SEM images of the as-deposited Sn–Co alloys and porous Sn–Co alloys.
ð2Þ
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Fig. 3. Discharge capacity (a) and Coulombic efficiency (b) of Sn–Co alloys before and after electrochemical dissolution treatment.
Fig. 4. Cyclic voltammograms profiles of C1 (a) and C2 (b).
In which Co remains as an inert component. The increasing discrepancy between the forward scan and reverse scan curves within the potential range of 1.5–2.0 V corresponding to charge of doublelayer capacitance reflects the increased surface area of the Sn–Co alloys after electrochemical dissolution treatment. The peak current of C1 increased with cycling in the first few cycles can be due to the continuous formation of new surface during lithiation and delithiation. For C2, except for the first cycle, the curve of other cycles coincides well, demonstrating improved reversibility. Charge–discharge curves of as-deposited alloy and porous alloy are obviously different. Charge–discharge curves of C1 and C2 were taken for example shown in Fig. 5. The charge–discharge curves of the C1 are similar to that of Sn. The platforms appearing at 0.65 V, 0.51 V and 0.41 V in discharge curves stand for the lithiation of Sn phase and formation of Li0.4Sn, LiSn and Li2.33Sn. The corresponding platforms in charge curves appearing at 0.77 V, 0.71 V and 0.58 V relate to the delithiation and formation of Sn, Li0.4Sn and LiSn [2]. The platform of C1 in the first discharge curve is lower due to resistance polarization. Since significant loss of Sn occur during electrochemical dissolution treatment, C2 have different charge– discharge behaviors with only one obvious platform observed in each curve. The platform appears at ca. 0.3 V in discharge curves and ca. 0.6 V in charge curves, relating to Eqs. (1) and (2), respectively. According to literature, for first discharge, reaction taking place at potential higher than 0.6 V gives irreversible capacity [39,40]. This irreversible capacity is due to decomposition of
electrolyte, reduction of tin oxide and formation of solid electrolyte interface (SEI) on the surface of electrode. C2 have higher irreversible capacity in the first discharge, which can be due to their higher surface area that consumes more capacity for formation of SEI and reduction of tin oxide. Porous structure owns larger surface area and can shorten diffusion distance of Li which will reflect on rate capability. The improved rate capability of porous structure formed from electrodeposition and electrochemical dissolution treatment was tested taking C1 and C2 as an example and shown in Fig. 6a. The charge–discharge rate rose stepwise from 0.2C to 10C and reduced to 0.2C after that. It is very clear that C2 delivered remarkable rate ability with the stable performance and high capacity retention. The 10C capacity of C2 is ca. 400 mA h g1 which is 67% of the 0.2C capacity tested after 10C. The high coulombic efficiency of C2 can be also observed at every rate and never less than 95%. Fig. 6b displayed the charge–discharge curves at different rates which performed identical shape. The C2 delivered 782 mA h g1, 618 mA h g1, 573 mA h g1, 485 mA h g1 and 400 mA h g1 at the rate of 0.2C, 1C, 2C, 5C and 10C respectively. According the above XRD and SEM results, we ascribe the improved performance in specific capacity, capacity retention and rate capability to several changes: (1) loss of Sn phase during electrochemical treatment reduces the volume expansion of the material during lithiation; (2) dissolution of Sn phase leaves many pores behind which buffers the volume expansion during lithiation; and
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Fig. 5. The charge and discharge curves of C1 (a) and C2 (b).
Fig. 6. Rate capabilities of C1 and C2 (a) and charge and discharge curves of C2 in different rate (b).
(3) electrochemical dissolution reduces the size of the alloy particles. All these changes can improve the cycling performance of the materials. And the reduced particle size and formation of hollow sphere increase the area contacting with electrolyte, which promotes the capacity of the materials. 4. Conclusions A new method to improve electrochemical performance of electrochemically deposited Sn–Co alloy materials for lithium-ion battery by subsequent electrochemical dissolution is put forward. During electrochemical dissolution treatment, part of the Sn phase in the Sn–Co alloys was dissolved and the surface structure of the residual alloy was improved. The reduced particle size and formation of hollow sphere promote the specific capacity of the resultant materials. After electrochemical dissolution treatment, specific capacity, capacity retention and rate capability of a Sn–Co alloy are greatly increased which fully shows subsequent dissolution treatment is an effective way to improve the electrochemical performance of Sn–Co alloy materials for lithium-ion batteries. Acknowledgements This work was financially supported by Natural Science Foundation of Shandong Province (Project No. ZR2009BM012) and the 973 Project of China (No. 2011CB935901).
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