Influence of copper addition for silicon–carbon composite as anode materials for lithium ion batteries

Abstract

A series of porous Si–C and Si–C/Cu composites have been successfully fabricated by a simple sol–gel and pyrolysis process. In the Si–C/Cu composites, nanoscale Si and Cu particles are homogeneously dispersed in the pyrolyzed carbon matrix. Furthermore, Cu₃Si phase has formed during the carbonization process confirmed by X-ray diffraction (XRD) and high-resolution transmission electron microscopy (HRTEM). As an anode material for lithium ion batteries, the Si–C/Cu2 composite exhibits a high initial discharge capacity of 2234 mA h g⁻¹ and a reversible discharge capacity of 947 mA h g⁻¹ after 100 cycles at the current density of 100 mA g⁻¹, respectively. With the current density gradually increasing to 1000 mA g⁻¹, the composite shows an average capacity of 848 mA h g⁻¹, exhibiting superior rate capability. The excellent cycling stability and rate discharge performance of the Si–C/Cu2 composite can be attributed to the improved conductivity owing to the addition of Cu, and the nanoporous structures as well as the formation of Cu₃Si, which both have good buffer effect to release volume expansion and maintain the integrity of the electrode during the charge–discharge cycles.

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1. Introduction

With the rapid development of electric vehicles and large-scale energy storage systems, rechargeable lithium ion batteries (LIBs) have attracted great attention owing to their high energy density and power density, long cycle life, low self-discharge rate, high operating voltage, wide operating temperature range, little memory effect and environmental friendliness compared to other rechargeable batteries in recent years.^1–5 However, the current commercial graphite anode has been greatly restricted due to its low theoretical capacity of 372 mA h g⁻¹ and safety problem caused by the close operation potential compared with lithium.^6,7 Therefore, many researchers make tremendous efforts to find new anode materials with higher specific capacity to meet the ever increasing energy requirements. Among them, silicon (Si) has been considered as one of the most promising anode candidates for the next-generation lithium ion batteries due to its natural abundance (the second most abundant element in the earth's crust), high theoretical capacity (∼3579 mA h g⁻¹ for the Li₁₅Si₄ phase) and low discharge potential (<0.5 V vs. Li⁺/Li).^4,8 Nevertheless, the practical application of silicon anode is currently hindered by its gigantic volume expansion (∼300%) during lithium insertion–extraction processes, relatively low electronic conductivity and instability of the solid electrolyte interface (SEI).^9–12 Therefore, much more research works are urgently needed to overcome the problems.

Various strategies have been developed to improve the electrochemical performances of Si-based anodes and significant progresses have been acquired in addressing the problems. These strategies mainly contains: (1) fabricating nanoscale Si with different structures including silicon nanowires and nanotubes,^9,13–16 silicon nanofilms,¹⁷ silicon nanoparticles and hollow nanospheres;^18,19 (2) preparing Si–C composites, such as carbon coated Si, graphene wrapped Si^20,21 and core–shell structures composites;²² (3) synthesizing Si–metal alloys and Si–metal–C composites, for instance, Si–Fe,^23,24 Si–Co,²⁵ Si–Ni,²⁶ Si–Cu,²⁷ Si–Al,²⁸ Si–Mg,²⁹ Si–Mn,³⁰ Si–Ti,³¹ Si–Zn,³² Si–Sn,³³ Si–Ca,³⁴ Si–Cr,³⁵ Si–Ni–C,³⁶ Si–Sn–C,³⁷ Si–Cu–C,^38,39 Si–Zn–C,⁴⁰ Si–Co–C,⁴¹ Si–Ti–C (ref. 1) have been explored in recent years. Among all the Si-based anode materials, Si–metal–C composites are the most possibly commercialized anode materials for lithium ion batteries due to their superior electrochemical performances, simple production technology, higher yields and lower production costs. As a representative example, Sukeun Yoon³⁹ et al. have prepared a carbon-coated Si–Cu₃Si composite material by means of mechanical milling and pyrolysis technique. The cycling performance of the composite delivers a stable capacity of 850 mA h g⁻¹ for 30 cycles. The improved cycling performance can be attributed to the copper silicide and pyrolyzed carbon, which both provide a better electrical contact with the current collector and a buffering effect for the volume expansion during cycling. By a simple high energy mechanical milling process, Jing Zhang⁴¹ et al. prepared a Si–Co–C composite with a initial discharge capacity of 1283.3 mA h g⁻¹ and a coulombic efficiency of 83.3%. In addition, the discharge capacity still remains above 610 mA h g⁻¹ after 50 cycles. The improved electrochemical performance can be contributed to the synergistic effect between C and Co, which improve the conductivity of the material and provide better electrical contact during cycling and suppress the volume expansion. All the mentioned Si–metal–C composites have shown good electrochemical performances in comparison with that of the corresponding binary systems as anode materials for lithium ion batteries.

Although great progresses have been obtained for the synthesis of Si–metal–C composites for lithium ion batteries, the synthetic route is almost by means of high energy mechanical milling and arc-melting. In this study, a series of Si–C and of Si–C/Cu composites have been successfully fabricated through a sol–gel and pyrolysis process. The characteristics of these composites as anode materials for lithium ion batteries are investigated using various analytical techniques.

2. Experimental section

2.1. Sample preparation

The Si–C/Cu composites were prepared by a simple sol–gel and along with a pyrolysis process. In brief, resorcinol (R, 18.2 mmol), formaldehyde (F, 36.4 mmol), hexadecyl trimethyl ammonium bromide (CTAB, 0.0364 mmol) and different quantity of cupric acetate (Cu(CH₃COO)₂·H₂O, 0.364 mmol, 0.607 mmol, 1.82 mmol) were dissolved in a certain amount of ethanol (342.5 mmol) by a magnetic stirrer at room temperature. After forming a homogeneous solution, nanocrystalline silicon powders (Si, 17.8 mmol, 99.99%, 80–100 nm, Tebo Technology Co., Ltd, China) were added, and dispersed through magnetic stirring, followed by ultrasonic treatment and finally magnetic stirring for 1 h, respectively. The wet samples were then sealed and cured at 70 °C for 7 days to obtain the corresponding xerogels. The xerogels were then dried in an oven at 70 °C under ambient pressure for over 12 h and finally carbonized at 800 °C for 2 h in Ar flow with a heating rate of 3 °C min⁻¹ and cooled down to room temperature naturally. Herein, the contents of the Cu(CH₃COO)₂·H₂O were set to 0.364, 0.607 and 1.82 mmol, and the corresponding composites were denoted as Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3, respectively.

As a comparison, another Si–C composite was prepared using the same procedure and prescription above, only without adding the cupric acetate. The as-prepared sample was denoted as Si–C.

2.2. Sample characterizations

X-ray diffraction (XRD) patterns of the as-prepared samples were collected by using a Bruker D8 Focus power X-ray diffractometer (40 kV and 40 mA) with CuKα radiation at a scan rate of 6° min⁻¹ in the range of 10–80° at room temperature. Raman spectra were taken on a JYLABRAM-HR Confocal Laser Micro-Raman spectrometer at room temperature. Scanning electron microscopy (SEM) images of the as-prepared products were obtained through a Hitachi S-4800 instrument at an acceleration voltage of 10 kV. Transmission electron microscopy (TEM) images and energy-dispersive X-ray spectroscopy (EDS) element mapping were performed on a FEI Tecnai G2 S-Twin instrument with a field emission gun operating at 200 kV. Thermo-gravimetric (TG) analysis were carried out from 25 to 800 °C with a heating rate of 10 °C min⁻¹ in air environment using an STA 449 Jupiter (NETZSCH) thermogravimetry analyzer to evaluate the carbon content. Inductively coupled plasma (ICP) optical emission spectrometer (Perkin-Elmer Instruments) was employed to measure the mass ratio of Si and Cu.

An amount of the as-prepared samples (150–200 mg) were heated to 200 °C under vacuum (10⁻⁵ Torr) for 2 h to remove all the adsorbed species. Nitrogen adsorption–desorption isotherms were then obtained at liquid nitrogen temperature (−196 °C) with a Micromeritics ASAP 2010 surface area and porosity analyzer. The total surface area was calculated according to the Brunauer–Emmett–Teller (BET) method. The pore volume and size distribution were analyzed by t-plot theory, Barrett–Joyner–Halenda (BJH) theory and density functional theory (DFT).

2.3. Electrochemical measurements

The electrochemical properties of the as-prepared samples were investigated by CR2025 coin-type cells with two electrodes assembled in an argon-filled glove box. The working electrode was prepared by mixing the active material (Si–C/Cu1, Si–C/Cu2, Si–C/Cu3 and Si–C), acetylene black, and polyvinylidene fluoride (PVDF) with a weight ratio of 70 : 20 : 10 in N-methyl-2-pyrrolidone (NMP) to form a homogeneous slurry, which was then coated onto a copper foil and dried in a vacuum oven at 60 °C for 12 h. After that, the electrodes were pressed using a small tablet press under a pressure of 10 MPa to enhance the contact between the material and the copper foil. Then, the CR2025 coin-type cells were assembled by using pure lithium foil as counter and reference electrode, a Celgard 2400 membrane as separator, and 1 M LiPF₆ dissolved in ethylene carbonate and diethylene carbonate with a volume ratio of 1 : 1 as electrolyte. It should be noted that this process was performed in a glove box filled with highly pure argon gas. The charge–discharge performance was tested galvanostatically in the voltage range from 0.005 to 1.5 V at various current densities using a programmable battery testing system (LAND CT2001A) at room temperature. Material loading was 1.15–1.2 mg cm⁻² (including the weight of acetylene black and the binder). The cyclic voltammogram (CV) and electrochemical impedance spectroscopy (EIS) measurements were performed on a Bio-Logic VMP3 Electrochemical Workstation. CV test was conducted with a scanning rate of 0.1 mV s⁻¹ and the potential vs. Li⁺/Li ranging from 0.005 to 1.5 V. EIS measurement was carried out with an oscillation amplitude of 5 mV in the frequency range from 700 kHz to 100 mHz.

3. Results and discussion

3.1. Characterizations of the as-prepared samples

The phase and purity of the as-prepared samples are investigated by XRD. As shown in Fig. 1(a), all the diffraction peaks can be well indexed to Si, Cu, Cu₃Si and C. The relatively broader peak of carbon suggests that the carbon generated during the carbonization process is amorphous carbon.^42,43 Meanwhile, metal Cu which is reduced can react with Si forming Cu₃Si during the pyrolysis process.⁴⁴ Notedly, the diffraction intensity of Cu₃Si is influenced by the content of Cu in the composites. The XRD intensity of Cu₃Si increases with the increase of Cu content. Fig. 1(b) shows the Raman spectra of the Si–C/Cu2 composite, as can be seen, there are two peaks located at about 1340 and 1598 cm⁻¹, which are corresponding to the characteristic D-band (the disordered band) and G-band (the graphene band) of graphite carbon. While the intensity ratio of the two peaks (I_D/I_G) is calculated to be 0.83, indicating the carbon in the composite is fairly amorphous.⁴⁵

Fig. 1 (a) XRD patterns of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3. (b) Raman spectrum of Si–C/Cu2.

The SEM images of the as-prepared samples are shown in Fig. 2. Seen from Fig. 2(a)–(d), the particles with diameter of several tens of nanometers are interconnected into grape-like aggregates, which are connected into a bulk network. It can be vividly seen that there are many nanopores among the particles and aggregates. This kind of unique highly porous network structure is very helpful for the transmission of lithium ion.⁴⁶

Fig. 2 SEM images of (a) Si–C, (b) Si–C/Cu1, (c) Si–C/Cu2 and (d) Si–C/Cu3.

The TEM images and EDS analysis of the Si–C/Cu2 composite are shown in Fig. 3. As depicted in Fig. 3(a), the dark domains are uniformly dispersed in the light area, indicating that all the Si, Cu and Cu₃Si are well distributed in the carbon matrix. To further confirm the morphology feature of the Si–C/Cu2 composite, high-resolution TEM (HRTEM) is carried out. As shown in Fig. 3(b)–(d), the lattice spacing of 0.181, 0.314 and 0.245 nm can well correspond to the (200) plane of Cu phase, (111) plane of Si phase and (200) plane of Cu₃Si phase, respectively. The highly distributed characteristic of the elements in the Si–C/Cu2 composite is also affirmed by EDS analysis. It can be seen from Fig. 3(e) that all the elements of the composite are uniformly dispersed and coexisted.

Fig. 3 (a) TEM image, (b)–(d) high-resolution TEM images and (e) EDS mapping of Si–C/Cu2.

The thermal behavior of the as-prepared samples are analyzed by thermo-gravimetric (TG) measurement. As shown in Fig. 4, the slight weight change below 300 °C is mainly associated with the loss of adsorbed water and ethanol on the surface of the prepared samples. The main weight loss of the obtained samples is after 350 °C, which is attributed to the oxidation of carbon in air environment to generate some gaseous products, such as CO and CO₂. According to the TG curves, the final weigh loss of the prepared samples is 62.3% for Si–C, 56% for Si–C/Cu1, 51.4% for Si–C/Cu2 and 46.1% for Si–C/Cu3, suggesting the carbon contents in the as-prepared samples decrease with the increase of Cu content. Besides, ICP analysis demonstrates the mass ratio of Si and Cu is about 31.6: 1 for Si–C/Cu1, 21.4: 1 for Si–C/Cu2 and 5.05: 1 for Si–C/Cu3, respectively. According to the TG results and ICP analysis and without considering the oxidation of Cu, the element analysis result is displayed in Table S1.†

Fig. 4 TG curves of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3.

Nitrogen adsorption–desorption isotherms are employed to further confirm the developed pore structures. As shown in Fig. 5(a), all the as-prepared samples have been demonstrated a typical IV isotherm with a distinct hysteresis loop according to the IUPAC classification, indicating the existence of mesoporous nanostructures. It can be noticed that the adsorption amount reduces with the increase of Cu content, confirming the degree of mesoporosity decreases. The nitrogen adsorption–desorption isotherms show uptakes at the relative low pressure suggest that all the as-prepared samples contain a certain amount of micropores.

Fig. 5 (a) Nitrogen adsorption–desorption isotherms and (b) corresponding pore size distribution curves of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3.

According to the nitrogen adsorption–desorption isotherms, surface area, total pore volume, average pore size and pore size distribution curves of the as-prepared samples are calculated and the results are displayed in Table 1 and Fig. 5(b). As can be seen, the surface area (S_BET) and total pore volume (V_total) decrease as the increase of Cu content, while the average pore size (D) increases. That is partly because cupric acetate itself is a catalyst and can facilitate the rapid formation of the gel network, resulting in the collapse of some pores. The pore size analysis further confirm the mesoporous nanostructures of the samples on the basis of the average pore size with the range of 4.03–5.26 nm.

Table 1 Surface area and porosity analysis of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3 composites

Sample	SBET (m^2 g^−1)	Smic (m^2 g^−1)	Sext (m^2 g^−1)	Vtotal (cm^3 g^−1)	D (nm)
Si–C	128.29	75.25	53.04	0.13	4.03
Si–C/Cu1	113.07	59.17	53.90	0.13	4.68
Si–C/Cu2	88.61	41.75	46.85	0.12	5.26
Si–C/Cu3	71.45	35.81	35.65	0.09	5.26

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3.2. Electrochemical performance of the as-prepared samples

To evaluate the electrochemical properties of the obtained samples, various electrochemical measurements have been carried out. Fig. 6(a)–(d) displays the initial two discharge–charge curves of all the samples measured at the current density of 100 mA g⁻¹ between 0.005 and 1.5 V. Note that the current density is based on the weight of the active material (Si + in situ carbon + Cu₃Si). The initial discharge and charge capacities are 1058 and 416 mA h g⁻¹ for Si–C, 1347 and 615 mA h g⁻¹ for Si–C/Cu1, 2234 and 1166 mA h g⁻¹ for Si–C/Cu2, 1844 and 1006 mA h g⁻¹ for Si–C/Cu3, respectively. Accordingly, the initial coulombic efficiencies of the samples are 39%, 46%, 52% and 55%. The irreversible capacity loss in the first cycle can be mainly attributed to the formation of solid electrolyte interphase (SEI) films on the electrode surface owing to the decomposition of the electrolyte,^47,48 and also to the reaction of Li with active sites in the electrode.⁴⁶ The following discharge–charge process reveals superior reversible capacities, there are 481 and 409 mA h g⁻¹ of the discharge and charge capacity for Si–C, 798 and 670 mA h g⁻¹ of the discharge and charge capacity for Si–C/Cu1, 1086 and 996 mA h g⁻¹ of the discharge and charge capacity for Si–C/Cu2, 975 and 889 mA h g⁻¹ of the discharge and charge capacity for Si–C/Cu3, with coulombic efficiency of 85%, 84%, 92% and 91%, respectively.

Fig. 6 Galvanostatic discharge–charge curves of the 1st and 2nd cycles for (a) Si–C, (b) Si–C/Cu1, (c) Si–C/Cu2 and (d) Si–C/Cu3.

In order to further assess the electrochemical performances of the samples, the cyclic voltammogram (CV) measurement has been conducted. Fig. 7(a)–(d) shows the CV profiles of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3 composites in a potential window of 0.005 to 1.5 V (vs. Li⁺/Li) at a scanning rate of 0.1 mV s⁻¹ for the initial five cycles. As can be seen, the CV curves of each sample are similar in shape. In the first cathodic scanning process, there are two reduction peaks located at approximately 1.43 and 0.7 V (vs. Li⁺/Li), which disappears at the subsequent cycles, can be contributed to the irreversible reactions between the electrode and electrolyte and the formation of a SEI film on the surface of active particles.⁴⁹ Both would cause the large irreversible capacity loss of the electrode in the first cycle. The sharp reduction peak between 0.2 and 0.01 V corresponds to the alloying process, specifically refers to the formation from crystallite Si to amorphous Li_xSi alloy, and finally transformation to crystalline Li₁₅Si₄ phase. Correspondingly, the anodic peaks at approximately 0.35 and 0.52 V (vs. Li⁺/Li) are related to the decomposition of crystalline Li₁₅Si₄ phase to amorphous Li_xSi phase, and finally to amorphous Si phase.^3,50 In the following cycles, a much more distinct lithiation peak appeared at about 0.18 V owing to the formation of amorphous Li_xSi phase.³ The increase in intensity of the current peaks could be attributed to the repeated activation (lithiation and delithiation) during the cycling process.⁵¹ The CV curves from the second cycle are increasingly overlapped, which suggests the high reversibility of the electrode in the lithiation–delithiation reaction during the electrochemical activation process.

Fig. 7 Cyclic voltammogram (CV) curves of (a) Si–C, (b) Si–C/Cu1, (c) Si–C/Cu2 and (d) Si–C/Cu3 at the scanning rate of 0.1 mV s⁻¹ in the range of 0.005 to 1.5 V (vs. Li⁺/Li).

The cycling behaviors and rate performances of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3 composites are tested and shown in Fig. 8(a) and (b). As can be seen from Fig. 8(a), Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3 deliver a discharge capacity of 1058, 1347, 2234 and 1844 mA h g⁻¹ in the first cycle at the current density of 100 mA g⁻¹, respectively. While the first charge capacities are 416, 615, 1166 and 1006 mA h g⁻¹ for each sample, corresponding to the initial coulombic efficiency of 39%, 46%, 52% and 55%, respectively. After the first few cycles, the coulombic efficiency of all the prepared composites steadily increases to around 99% and then remains stable. It is obvious that each composite displays a relative stable cycling performance even after 100 cycles, which could be contributed to the promising nanoscale carbon matrix and porous structures. The developed carbon matrix and continuous nanoporous structures could efficiently buffer the mechanical stress generated by the volume change of active material and maintain the structural integrity of anode material during the lithiation–delithiation process. Among all the samples, Si–C/Cu2 shows the most outstanding performance with a reversible discharge capacity of 947 mA h g⁻¹ and capacity retention of 88% (compared with the discharge capacity of the second cycle) after 100 cycles, respectively. Furthermore, Si–C/Cu2 also displays an initial discharge capacity of 2.569 mA h cm⁻² and a reversible discharge capacity of 1.089 mA h cm⁻² after 100 cycles under the current density of 0.115 mA cm⁻² (Table S2†). It is reasonable to assume that the high discharge capacity of the Si–C/Cu2 composite is probably owing to the relatively high content of Si in the composite and the addition of Cu can improve the conductivity of the composite and the utilization of active materials to some extent. While the formation of Cu₃Si can also have the buffer effect to release the stress from volume expansion and raise the cycle stability of the materials.³⁹ In addition, RF-derived carbon is also synthesized and its cycle performance is tested in order to elucidate the contribution of carbon in the as-prepared composites. The carbon only shows a reversible average discharge capacity of 190 mA h g⁻¹ at the current density of 100 mA g⁻¹ (Fig. S1†), thus the high discharge capacity of Si–C/Cu2 is mainly contributed by Si.

Fig. 8 (a) Cycling performance at the current density of 100 mA g⁻¹ and (b) the rate capabilities of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3.

The high rate discharge capacities of the samples are conducted at various current densities ranging from 100 to 1000 mA g⁻¹, and then back to 100 mA g⁻¹. It can be seen that Si–C/Cu2 exhibit large reversible average discharge capacity of 1049, 990, 945 and 848 mA h g⁻¹ at the current density of 100, 200, 500 and 1000 mA g⁻¹, respectively. When the current density returns from 1000 to 100 mA g⁻¹, the reversible discharge capacity recovers to 933 mA h g⁻¹ and maintains at the value of 899 mA h g⁻¹ after 100 cycles. Meanwhile, the ultra-high rate test of the Si–C/Cu2 composite is also carried out, it shows an average discharge capacity of 455 and 299 mA h g⁻¹ at the current density of 5 and 10 A g⁻¹ (Fig. S2†), respectively. The excellent performance of Si–C/Cu2 composite can be ascribed to the following factors: (1) the continuous nanoporous structures are helpful for the diffusion and transportion of the electrolyte and lithium ions; (2) the continuous nanoporous structures can also relax the stress and alleviate the structure decomposition induced by lithium ions insertion–extraction. (3) The addition of a mount of Cu can enhance the conductivity of the composite, and thus improving the utilization of active material and charge–discharge efficiency of the anode electrode; (4) the formation of Cu₃Si also has the buffer effect to accommodate volume expansion and maintain the integrity of the electrode during the charge–discharge process.

To further clarify the resistance against charge and ion transfer during cycles, the EIS measurement of the as-prepared composites has been performed before discharge–charge process (Fig. 9). As can be seen, the Nyquist plot consists a depressed semicircle in the high and medium frequency region corresponding to the charge-transfer resistance (R_ct), and a sloped straight line in the low frequency region corresponding to the Warburg diffusion resistance (W).^52,53 Obviously, Si–C/Cu2 displays the lowest R_ct value among all the prepared composites. This result indicates that moderate addition of Cu can provide better electrical contact and improve the conductivity of the composite, and thus boost the electron transport and charge transfer on the electrode/electrolyte interface.

Fig. 9 Electrochemical impedance spectrum (EIS) of Si–C, Si–C/Cu1, Si–C/Cu2 and Si–C/Cu3 before cycling.

4. Conclusions

In summary, a series of porous Si–C and Si–C/Cu composites have been successfully fabricated by a simple sol–gel and pyrolysis process. In the Si–C/Cu composites, Si and Cu can be homogeneously dispersed in the carbon matrix. The addition of Cu can improve the conductivity of the composites and enlighten the utilization of active materials, thus the Si–C/Cu composites have higher reversible discharge capacities and rate performances compared with the Si–C composite. Among all the Si–C/Cu composites, Si–C/Cu2 exhibits the best cycling stability and rate discharge ability. In view of the simple synthetic method and excellent electrochemical performances, the porous Si–C/Cu composites are considered to have wide application prospect for lithium ion batteries in the future.

Acknowledgements

This work is supported by the Creative Research Groups of the National Natural Science Foundation of China (Grant No. 21221061).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra12332e

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