Chapter 4 One-Step Vapor–Solid Reaction Growth of Sn@C Core–Shell
4.3.3 Electrochemical properties of Sn@C core–shell NWs
4.3.3 Electrochemical properties of Sn@C core–shell NWs
According to the literature, employing SnO2 and Sn for anode applications in LIBs presents several drawbacks. These include the irreversibility associated with the SnO2 reduction, the capacity losses due to the formation of the solid electrolyte interface (SEI) on the Sn surfaces, and large volume variations during the Li+ insertion and extraction reaction steps.38–40 For the Sn@C core–shell NWs prepared in this study, we performed the following electrochemical studies. Half-cells composed of a Li foil, as the negative electrode (anode), and Sn@C core–
shell NWs, as the positive electrode (cathode), were assembled into test cells. In Figure 4.9a, the first cyclic voltammetric (CV) profile of the NWs shows a reduction peak (cathodic scan) at 1.52 V. The peak is irreversible because it disappears in the later cycles. One possible origin
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of the peak was from the SEI on Sn.41 However, another signal peak at 1.05 V usually accompanies the one near 1.52 V. In addition, since most of the Sn containing nanostructures are wrapped inside the C shells, it is less likely that the SEI could form on the Sn surface. The peak at 1.52 V cannot be assigned to the SEI on the C shells either because the reported signals were at 0.80 and 1.25 V.41 Thus, we ascribe the peak at 1.52 V to the formation of LiOH due to the presence of traces of H2O in the electrolyte.42 The other reduction peaks are observed at 0.58 V, 0.46 V and 0.30 V during the discharging/alloying step while the oxidation peaks (anodic scan) are found at 0.54 V, 0.71 V, 0.80 V and 0.85 V during the charging/dealloying step. These are correlated with the interconversions among Sn and several LixSn (0 x 4.4) phases, including Li2Sn5, LiSn, Li7Sn3, Li5Sn2, Li13Sn5, Li7Sn2, and Li4.4Sn.10,17,43,44 Previous studies indicated that primary lithiation and delithiation processes of amorphous carbon and graphite materials usually occurred below 0.25 V while amorphous carbon nanotube (CNT) and single-wall CNT showed Li deintercalation at 0.9-1.2 V.45–50 The CV profiles in Figure 4.9a display gradual current decreases below 0.25 V and no clear deintercalation peaks at 0.9-1.2 V in the initial cycles. The observations suggest that Li+ ions probably accumulate irreversibly in the carbon shells.
Figure 4.9b displays the voltage profiles of the half-cell made of the Sn@C NWs. The data were taken at a cycling rate of 500 mA g-1 and between 0.005 and 2.00 V versus Li/Li+. The voltage profiles indicate that the NW electrode exhibits the characteristics of a Sn electrode.10 The first discharge and charge steps deliver specific capacities of 1760 and 1405 mA h g-1, respectively. They correspond to a coulombic efficiency of 80%. The large initial capacity loss can be attributed, as mentioned above, to the formation of the SEI layer on the electrode surface during the first discharge step and the storage of Li+ ions in EC/DMC-based electrolytes, which are difficult to be extracted.10,17,39 The presence of multiple plateaus in the initial discharge–
charge curves is assigned to the formation and decomposition of LixSn.43,44 In the following cycles, the plateaus become less and less discernible. For example, the plateaus at 0.58 V (the
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transition between Sn and LiSn phases) and 0.46 V (the transition between LiSn and Li7Sn3
phases) decrease gradually with increasing cycle numbers. These observations agree with the CV results. They suggest that Li+ ions stored in the related LixSn phases may not undergo reversible cycling. The discharge capacities are found to be 1384, 804, 582, 515 and 490 mA h g-1 at the second, tenth, twenty-fifth, fiftieth and one hundredth cycles, respectively.
Figure 4.9c depicts the specific capacity and the coulombic efficiency of the discharge–
charge process of the half-cell with a cycling rate at a current density of 500 mA g-1. Obviously, the capacity dropped swiftly for the first ten cycles. As discussed above, there are several possible pathways which may provide irreversible storage of Li+ ions. These include the formation of the SEI layer, the decomposition of electrolytes, and the accumulation of Li+ ions in carbon and tin materials. They offer an apparent capacity which exceeds the theoretical value.
Between the twenty-fifth and the fiftieth cycles, the half-cell remains stable during the cycling and exhibits a fade rate of 0.47% per cycle. The half-cell continues to be stable for the next fifty cycles with a fade rate of 0.09% per cycle. The specific capacity is 490 mA h g-1 at the end of the one hundredth cycle. The coulombic efficiency remains relatively stable at over 98% after the twenty-fifth cycle. In contrast, the cycling performance of a half-cell constructed from commercial Sn powders (particle size distribution 1–10 m)
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Figure 4.9 (a) Cyclic voltammograms of a Sn@C core–shell NW electrode (scan speed 0.5 mV s-1). Electrochemical performance of the electrode cycled between 0.005 V and 2.0 V vs. Li/Li+ (cycling rate 500 mA g-1); (b) voltage profiles of the electrode after 1, 2, 10, 25, 50, and 100 cycles and (c) capacity fading of the electrode. Coulombic efficiency and reversibility of each cycle of the electrode are presented in the secondary y-axis on the right of (c). The discharge capacity of an electrode fabricated from commercial Sn powder was cycled at 100 mA g-1.
0 500 1000 1500 2000
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Figure 4.10 shows the cycling rate capability of the half-cell constructed from the Sn@C core–shell NWs with a voltage window of 0.005–2.0 V. Figure 4.10 shows the discharge capacities of the device at current densities of 100, 500, 1000, and 3000 mA g-1. After one hundred cycles, the discharge (Li alloying) capacities at these rates are 525, 490, 486, and 286 mA h g-1, respectively. Both of the discharge capacities at 500 and 1000 mA g-1, compared to the one at 100 mA g-1, are about 93%. Even at 3000 mA g-1, 55% of the capacity at 100 mA g
-1 remained. In Figure 4.10, a capacity of 905 mA h g-1 is observed after the battery is cycled at 100 mA g-1 for five times. Then, after continued cycling at 500 mA g-1, 1000 mA g-1, and finally 3000 mA g-1 for five times each, the half-cell returns back to 626 mA h g-1 at 100 mA g-1. This demonstrates that even after fast discharge–charge cycles at 3000 mA g-1, the electrode was not severely degraded and the half-cell still exhibited excellent cycling properties. Consequently, we conclude that the improved electrochemical performance of the half-cell is rooted in the special morphology of the Sn@C core–shell NWs.
Figure 4.10 Electrochemical performance of a Sn@C core–shell NW electrode cycled between 0.005 V and 2.0 V vs. Li/Li+. (a) Curves of specific capacity versus cycle number of the electrode at cycling rates of 100 mA g-1, 500 mA g-1, 1000 mA g-1 and 3000mA g-1. (b) Discharge capacities of the electrode as a function of discharge rates (100–3000 mA g-1).
The NW structure may shorten the transport lengths for both electrons and Li+ ions. The C shell layer probably acts as a good electronic conductor and serves as a buffer for the volume change during the lithiation–delithiation process. Figure 4.11a and 4.11b shows the images of
0 10 20 30 40 50 60 70 80 90 100
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the as-prepared electrode material containing the Sn@C core–shell NWs. After one hundred discharge–charge cycles at 500 mA g-1 (cycled between 0.005 and 2.0 V), many Sn@C core–
shell NWs still maintain original diameters with clear core–shell structures in Figure 4.11c and 4.11d. Clearly, the volume expansion caused by the lithiation of the Sn cores was confined within the C shells. Two possible reasons are proposed to rationalize the observed confinements.
One of them is attributed to the voids observed in Figure 4.1 and 4.7, which may accommodate some of the expansions. The other one may originate from the flexible amorphous C shells which could be deformed extensively during the volume variations.51 In Figure 4.11c and 4.11d, only some NWs are found to be destructed into aggregated particles, probably due to non-uniform volume changes. On the other hand, as shown in Figure 4.12, the electrode fabricated from commercial Sn powders shows severe aggregations after fifty discharge–charge cycles.
Because the NWs retain most of the original morphology, they exhibit better battery performance than the Sn particles do. Comparison of our results to other literature reports is discussed below.
Figure 4.11 SEM images of the electrode prepared from Sn@C core-shell NWs. (a) Low and (b) high magnification views. (c) Low and (d) high magnification views of the electrode after 100 cycles of lithiation and delithiation. The electrode was fabricated with a mixture of carbon black and binder.
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Figure 4.12 SEM images of (a) the commercial Sn powder, (b) the electrode prepared from the powder, and (c) low and (d) high magnification views of the electrode after 50 cycles of lithiation and delithiation. The Sn particles shown in (b) and (c) were mixed with carbon black and binder used for the electrode fabrication.
Employing SnO2 as a potential anode material for LIBs has been attempted in many studies.21,22 Unfortunately, the oxide suffered from a high first-cycle irreversible capacity during the first charging because it has to be reduced to metallic Sn. To solve this problem, it is anticipated that metallic Sn could be applied directly. However, utilizing the metal presents another difficulty as discussed earlier. It undergoes enormous volume variation during the Li alloying and dealloying processes.17,20 One possible solution for this is to use nanostructured hollow and/or porous Sn containing materials. It is anticipated that the local empty spaces may accommodate the large volume change.52,53 Another possible answer is to restrain the active phases by controlling the size and morphology of the electrode materials. To achieve these, many nanostructured core–shell and composite materials have been developed.9–11,25 Literature reports related to solving the issues are summarized in Table 4.1. For example, the initial capacity of SnO2 nanotubes (NTs) was 940 mA h g-1. After eighty cycles at 0.05 mA cm-2, a capacity of 525 mA h g-1 was retained, corresponding to 55.8% of the initial value.21 In another example, SnO2–core/C–shell NTs were investigated.22 This nanocomposite could be repeatedly
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charged and discharged up to two hundred cycles with a capacity of 542 mA h g-1 at 0.3 mA cm-2. Some Sn metal containing cases presented in Table 4.1 are discussed below. For hydrobenzamide-capped Sn, a charge capacity of 500 mA h g-1 was observed after thirty cycles at 0.2C.23 In the second case, electrochemical tests of vertical arrays of 1D Sn NWs on Si substrates showed that a discharge capacity of 400 mA h g-1 could be maintained after fifteen cycles at a high discharge-charge rate of 4200 mA g-1.24 Cases of Sn@C@CNT nanostructures and rambutan-like Sn–C nanocomposites offered capacities of 490 mA h g-1 and 311 mA h g-1, respectively.9,10 Lastly, examples of nanocomposites of Sn and Group 11 metals are presented.
Sn@Cu core–shell NPs retained a capacity of 560 mA h g-1 after one hundred seventy cycles at a rate of 0.8C.25 The discharge capacity of Sn NPs supported on nanoporous Au was found to be 440 mA h g-1 after one hundred forty cycles at 0.1C.11 Performance of the devices fabricated from our core–shell NWs compares favorably against the results of the better ones shown in Table 4.1. The capacity of 525 mA h g-1 was retained after one hundred cycles at a rate of 100 mA g-1. Even at a high rate of 1000 mA g-1 for one hundred cycles, our cell maintained a capacity of 486 mA h g-1.
4.4 Conclusions
In summary, we have fabricated Sn@C core–shell NWs in high yields by a simple and straightforward process. By reacting vapor phase C2H2 with SnO2 solids, we grew the 1D NWs via the VSRG pathway. Due to the unique 1D core–shell morphology, which provides an adequate void volume inside the flexible C shell, the NWs can respond to the large volume change caused by the formation and decomposition of LixSn alloys in the charge–discharge steps in LIBs. The electrodes show high reversible capacities, low rates of capacity fading, and consistent cycling performance.
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Chapter 5 Conclusions
In this thesis, we demonstrated a facile simple, and low cost fabrication of tin-based nanomaterials, including as SnO2 nanorods, SnO2 hollow spheres, SnO2 nanosheets and Sn@C core-shell nanowires. Report a unique reaction employing vapor phase and solid particles as the reactants, one-dimensional SnO2 nanorods and Sn@C core-shell nanowires are formed via vapor-solid reaction growth pathway. SnO2 hollow spheres and nanosheets could be controlled by Sn+4/+2 precursors via a simple one-pot template-free hydrothermal method. These are summarized in Scheme 5.1. The SnO2 nanorods were length 1-2 m and diameter 10-20 nm. The shell thickness of the hollow spheres was around 200 nm with diameter 1-3 μm, while thickness of the nanosheets was 40 nm. The Sn@C core-shell nanowires was 100-350 nm with 30-70 nm C shell thickness and length was several micrometers.
The effect of morphological modification on the electrochemical properties of tin-based nanomaterials for LIBs were further evaluated. Based on the results, reducing the active particle size to the nanometer range and using the specific morphology of active materials can significantly improve the cycling performance. Also, propose a conception that due to the even distribution of the voids among the inactive matrix, the mechanical stress caused by the volume changes could be alleviated during cycles. At the end of the one hundredth cycle, the specific capacity and the columbic efficiency stay relatively stable and the columbic efficiency over 98% are observed. The specific capacity of SnO2 nanorods, SnO2 hollow spheres, SnO2
nanosheets and Sn@C core-shell nanowires are 435, 522, 490 and 525 mA h g−1, respectively.
In contrast, the cycling performances of commercial tin-based powders are poor. We anticipate that the tin-based nanomaterials can be utilized for LIBs upon further optimization.
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Scheme 5.1 Illustration of tin-based nanostructures as electrochemical test for LIBs.
Current rate: 100 mA g-1
Sn@C core/shell NWs SnO2 NRs
SnO2 HSs SnO2NSs
VSRG pathway
Hydrothermal method
0 20 40 60 80 100
0 500 1000 1500 2000
Sn@C NWs, 525 mA h g-1 SnO2 HSs, 522 mA h g-1 SnO2 NSs, 490 mA h g-1 SnO2 NRs, 435 mA h g-1
Columbic efficiency (%)
Specific capacity / mAh g-1
Cycle number
60 70 80 90 100