SnO2 Model Electrode Cycled in Li-Ion Battery Reveals the Formation

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SnO2 model electrode cycled in Li-ion battery reveals the formation of Li2SnO3 and Li8SnO6 phases through conversion reactions Giulio Ferraresi, Claire Villevieille, Izabela Czekaj, Michael Horisberger, Petr Novák, and Mario El Kazzi ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b19481 • Publication Date (Web): 14 Feb 2018 Downloaded from http://pubs.acs.org on February 15, 2018

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ACS Applied Materials & Interfaces

SnO2 model electrode cycled in Li-ion battery reveals the formation of Li2SnO3 and Li8SnO6 phases through conversion reactions

Giulio Ferraresi†, Claire Villevieille†, Izabela Czekaj§, Michael Horisberger‡, Petr Novák†, Mario El Kazzi†, * †

Paul Scherrer Institute, Electrochemical Energy Storage Section, CH-5232 Villigen PSI, Switzerland

§

Krakow University of Technology, Faculty of Chemical Engineering and Technology, Krakow, Poland

‡

Paul Scherrer Institute, Laboratory for Scientific Developments and Novel Materials, CH-5232 Villigen PSI, Switzerland

Keywords: SnO2; Thin film; Anode; Electrochemistry; XPS; Li-ion battery; Li2SnO3; Li8SnO6; DFT; SEM

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ABSTRACT

SnO2 is an attractive negative electrode for Li-ion battery owing to its high specific charge compared to commercial graphite. However, the various intermediate conversion and alloy reactions taking place during lithiation/delithiation, as well as the electrolyte stability have not been fully elucidated, and many ambiguities remain. An amorphous SnO2 thin film was investigated for use as a model electrode by a combination of post-mortem X-ray photoelectron spectroscopy (XPS) supported by density functional theory (DFT) calculations and scanning electron microscopy (SEM) to shed light on these different processes. The early stages of lithiation reveal the presence of multiple overlapping reactions leading to the formation of Li2SnO3 and Sn0 phases between 2 V and 0.8 V vs. Li+/Li. Between 0.45 V and 5 mV vs. Li+/Li Li8SnO6, Li2O and LixSn phases are formed. Electrolyte reduction occurs simultaneously in two steps, at 1.4 V and 1 V vs. Li+/Li, corresponding to the decomposition of the LiPF6 salt and EC/DMC solvents, respectively. Most of the reactions during delithiation are reversible up to 1.5 V vs. Li+/Li, with the re-appearance of Sn0 accompanied by the decomposition of Li2O. Above 1.5 V vs. Li+/Li, Sn0 is partially re-oxidized to SnOx. This process tends to limit the conversion reactions in favor of the alloy reaction, as also confirmed by the long-term cycling samples.

INTRODUCTION The replacement of current Li-ion battery electrodes with a new family of materials is crucial to improve their electrochemical performance and deliver higher energy density1. On the anode side, the specific energy of the battery can be improved by replacing commercial graphite with


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materials able to provide a higher specific charge in a similar potential window. Conversion materials like SnO2 have drawn attention because of their high theoretical specific charge (1491 mAh/g)

2-3

. The peculiarity of SnO2 is its ability to react with lithium in two distinct

processes4-8: conversion and alloy reactions. The conversion reaction (I) is predicted to occur at a potential above 1.2 V vs. Li+/Li, followed by the alloy reaction (II) at a potential below 0.5 V vs. Li+/Li as demonstrated by several experimental techniques including transmission electron microscopy (TEM), X-ray diffraction (XRD) and X-ray absorption spectroscopy (XAS) 4, 9-10. Reaction (I)

SnO2 + 4Li+ + 4e- ↔ Sn + 2Li2O (711 mAh/g)

Reaction (II) Sn + xLi+ + xe- ↔ LixSn (x ≤ 4.4) (783 mAh/g) The conversion reaction (I) was believed to be irreversible, causing the initial high irreversibility. The alloy reaction (II) is considered to be the main reversible reaction, delivering a theoretical specific charge of 783 mAh/g.2 To date, most of the research on these materials has focused on the improvement of the electrochemical performance by investigating SnO2 nanoparticles.9, 11-12 Simultaneously various experimental reports5,

13-14

have also attempted to look deeper into the conversion and alloy

reaction pathways that occur upon lithiation/delithiation to elucidate the origin of the extra specific charge achieved compared to the theoretical maximum.15 This extra specific charge observed during delithiation is thought to be related to the reversibility of Reaction (I) through the decomposition of Li2O, leading to the oxidation of Sn0 into a SnOx phase. Soft and hard Xray photoelectron spectroscopy (XPS) investigations16 revealed the overlap of the conversion, alloy, and electrolyte decomposition reactions in the 1.2–0.4 V vs. Li+/Li region. These overlapping reactions are suggested to lead to the formation of a passivation layer on the surface of Sn0 that prevents the further conversion of SnO2 particles, thereby limiting the full discharge


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of the battery. Such a conclusion contradicts those of former studies that state that the extra specific charge detected on nanoparticles implies that SnO2 is fully taking part in the conversion/alloy reactions. Furthermore, theoretical calculation17-18 reveals the presence of the intermediate Reactions (III), (IV), (V) and (VI) occurring between the Reaction (I) and (II) during SnO2 lithiation. Reaction (III)

SnO2 + Li → 0.5 SnO + 0.5 Li2SnO3

Reaction (IV) 2 SnO + 1.33 Li → 0.67 Li2SnO3 + 1.33 Sn Reaction (V)

0.67 Li2SnO3 + 1.33 Li → 0.33 Li8SnO6 + 0.33 Sn

Reaction (VI) 0.33 Li8SnO6 + 1.33 Li → 2 Li2O + 0.33 Sn Similar reactions have been reported for SiO219-20 but have never been experimentally reported in the literature for SnO2. The difficulties encountered in identifying such intermediate reactions using X-ray diffraction, absorption and photoemission techniques is believed to be related to the amorphous nature of the new formed phases and to the little changes in absorption and photoelectron signals between SnO2, SnO and Li2SnO3.21-22 This challenge is overcome here by investigating model electrodes. Thin films electrodes are known for their outstanding electrochemical performance thanks to their fast Li-ion diffusion and the improvement of the mechanical stability, enabling better resolution of the redox reaction peaks in the cyclic voltammetry and allowing overlapped reactions to be de-convoluted. Thin film electrodes composed only of active material simplify the investigation by XPS considerably, thanks to the flat, dense, and smooth surface, together with the suppression of interference signals from carbon and binder (see also Note 1). Additionally, with such a flat surface any volume change related to the SnO2 or Sn lithiation can be easily observed with the SEM, since a strong surface morphological change is expected for the structures like Li2SnO3 and Li8SnO6 (see Table S1


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Supporting Information). Finally, the SEI is expected to be thin on model film compared to that with conventional electrodes, making the monitoring of the oxidation of the active materials possible even at low potentials without the need for a hard X-ray source. We previously investigated Si thin films,19 showing the benefits of using thin films as model electrodes coupled to surface investigations to simplify the understanding of the XPS spectra and aid the identification of electrochemical reactions. In this context and in order to shed light on these controversial reactions, we revisited the topic using a 20-nm film of amorphous SnO2 film as a model electrode. The surface and interface between the thin film and carbonate-based electrolyte were investigated using post-mortem X-ray photoemission spectroscopy (XPS) supported by the density functional theory (DFT) to estimate the binding energy shift of the Sn3d core level of the different phases and scanning electron microscopy (SEM) to monitor any volume changes of the SnO2 film during lithiation.

2. EXPERIMENTAL SECTION 2.1 Thin film deposition 20-nm thick amorphous SnO2 films were deposited on a copper foil at room temperature using a DC magnetron sputtering apparatus (TIPSI) (the thickness is confirmed by the SEM image displayed in Figure 1b). A Sn target was mounted on a Pk75 3″ water cooled cathode. The preliminary vacuum was 6.7 × 10-7 mbar. The distance between the target and the foil was about 80 mm. During the deposition, the partial pressures of Ar and O2 were 1.4 × 10-3 mbar and 1.9 × 10-3 mbar, respectively. The gas plasma was power regulated at a constant power of P = 80 W, and the deposition rate was about 0.4 nm/s. 2.2 Cells and Electrochemistry


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The test electrodes were obtained by cutting the deposited film in 13 mm diameter samples, which were stored in an Ar-filled glovebox after being dried at 80 °C overnight under dynamic vacuum. The mass of the active material was indirectly calculated by considering a fully dense film (6.85 g/cm3). Half cells were assembled using Celgard 2400 soaked with LP30 electrolyte (1 M LiPF6 salt in ethylene carbonate: dimethylcarbonate 1:1 wt.% - BASF SE) as a separator and lithium metal as the counter electrode. Galvanostatic cycling (GC) and cycle voltammetry (CV) measurements were measured using a Bio-Logic VMP3 instrument at room temperature (25 °C). Galvanostatic cycling was performed at a 1C rate (one hour lithiation, one hour delithiaiton) in a potential range of 5 mV – 2.0 V vs. Li+/Li. All the potentials reported in the study are given vs. Li+/Li redox potential. Cyclic voltammetry was performed at a scan rate of 50 µV/s in the same potential range. For post-mortem investigation, each cell was stopped at specific potential values, then disassembled inside an argon-filled glovebox and rinsed with a dimethyl carbonate solution (BASF SE) before being dried and transferred to the XPS or SEM using an argon-filled or a vacuum transfer chamber, respectively. The sample measured at the open-circuit potential (OCP), was placed inside the cell without applying any current and rest for 48 h before being treated as described above. 2.3 X-ray photoemission spectroscopy (XPS) The XPS measurements were conducted with a VG ESCALAB 220iXL spectrometer (Thermo Fisher Scientific) using focused monochromatized Al Kα radiation (1486.6 eV) with a beam size of ∼500 µm2 (power of 150 W). The pressure in the analysis chamber was approximately 2 × 10−9 mbar. The spectrometer was calibrated on the clean silver surface by measuring the Ag3d5/2 peak at a binding energy (BE) of 368.25 eV with a full width at half maximum (FWHM)


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of 0.78 eV. All the spectra were recorded under the conditions of 30 eV pass energy and 50 eV for the surveys in steps of 50 meV and dwell time of 50 ms. The peak deconvolutions were realized by applying the sum of 70% Gaussian – 30% Lorentzian line shapes after Shirley-type background subtraction. Control spectra were recorded at the beginning and at the end of each experiment as well as on different spots to ensure the reproducibility of the measurement and the homogeneity of the surface and to certify that the sample did not evolve with time or suffer beam damage. No calibration of the binding energy peak positions is applied in this study since all the core levels evolve during lithiation/delithiation and have the tendency to shift, even the hydrocarbon (CH2) peaks in the C1s spectra, which are commonly selected as a reference in the literature.16, 23-24 As we previously reported for an amorphous Si thin film19, the Si2p signal is not affected by the charging effect making the assignment of the lithiation/delithiation phases of Si and SiO2 very accurate. The same methodology was applied to the SnO2 film, since the Sn3d5/2 is also not expected to experience any charging effects. 2.4 Scanning Electron Microscopy (SEM) The cycled samples analyzed by XPS were also analyzed by SEM and compared to pristine and OCP samples. SEM measurements were performed with a Carl Zeiss Ultra55 scanning electron microscope, using the secondary electron detector to emphasize the surface morphology. The images were taken at an acceleration voltage of 3 kV and a working distance of ~ 3 mm to ensure limited damage of the species formed upon cycling, especially polymeric species coming from the electrolyte decomposition. During the lithiation of the amorphous SnO2 film, an important surface modification is expected based on the large volume change of the different phases as reported in Table S1 (Supporting Information).


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3. RESULTS AND DISCUSSION 3.1 Pristine electrode The amorphous SnO2 thin film was first characterized to assess its surface morphology and composition. Figure 1 shows the SEM images and the XPS analysis of the pristine (as-deposited) sample. From a morphological point of view, the film homogeneously covered the Cu substrate as a uniform layer (Figure 1a). As shown in Figure 1b, the SnO2 film, with an estimated thickness of 20 nm ± 3 nm, keeps its integrity even if it is scratched from the substrate. The XPS survey (Figure 1c) of the pristine sample gives information about the chemical composition of the film at a depth of ~ 8 nm. As expected, we detected the Sn3d, O1s peaks associated with the deposited film together with the C1s peak, related to adventitious carbon, always observed on the surface of samples exposed to air. Since no peaks related to copper were observed, we assumed that the SnO2 film was homogeneously deposited on the substrate and that no relevant alloy reaction with the Cu took place. Sn was acquired in the +4 oxidation state based on the BE position and the FWHM of the Sn3d5/2 core level measured at 487.15 eV and 1.5 eV, respectively (Figure 1d). However, it is well known that it is very difficult to discriminate between the Sn +4 (SnO2) and +2 (SnO) oxidation states by XPS, as the BE difference is just 0.6 eV.25-26 The Sn3d5/2 core level was also measured on a native Sn oxide film that was grown on top of a thin Sn film, where a mixture of both SnO2 and SnO is expected to be formed (Figure S1, Supporting Information), to support our assumption. Indeed, a much larger FWHM of 1.9 eV with a BE of 487.7 eV was measured, confirming the absence of SnO in the deposited SnO2 film.


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Figure 1. (a) Top-view SEM image of the SnO2 film; (b) Detail on peeled-off SnO2 film; (c) XPS survey, (d) Sn3d5/2, (e) C1s and (f) O1s core levels spectra of 20 nm SnO2 pristine film.

3.2 Electrochemical investigation The electrochemical investigations of the 20 nm SnO2 electrode films in half-cell configuration are reported in Figure 2. The cyclic voltammetry (CV) measurement displayed in Figure 2a shows well-defined peaks contrary to the results found in the literature of conventional electrodes with nm-/µm-sized spherical SnO2 particles.16,

22

Five distinct peaks were

distinguishable during lithiation and three peaks during delithiation. A high background current contribution was observed between 2.0 and 0.7 V vs. Li+/Li during the first lithiation, which is believed to be related to side reactions happening on the copper substrate (i.e., copper native oxide reduction and electrolyte reduction), processes already described in our previous study on


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Si thin film19. Such reactions can be much more enhanced in the case SnO2 thin film because thicker copper oxide layer is expected since the film is deposited under oxidative atmosphere. Moreover, in the same potential range we cannot exclude as well the contribution of the electrolyte reduction Regardless of this process, during lithiation a first broad peak was detected at ~1.7–1.5 V followed by a very narrow peak at 1.2 V. From experimental reports

6-7

, SnO2 is

expected to react directly with lithium to form Sn0 and Li2O above 1.0 V according to Reaction (I); however, theoretical calculations17-18 suggest that SnO and LixSnOy phases would be formed before the conversion reaction was complete. To the best of our knowledge, the intermediates in Reactions (III), (IV) and (V) have never been experimentally observed. Pursuing the lithiation, the peak observed at 0.95 V can be attributed to electrolyte decomposition as the conversion reaction is expected to be finished above 1.0 V. The two main peaks at 0.6 V and 0.25 V are generally attributed to the alloy reaction between Li and Sn, leading to the formation of two LixSn alloys with different lithium contents.4 Three main peaks were recorded during the first delithiation: a narrow peak at 0.5 V; a larger peak between 0.9–1.2 V, both associated with the reversible delithiation of LixSn phases; and a broad shoulder between 1.3–2.0 V attributed to the reversibility of Reaction (I) with oxidation of Sn0 to SnOx. During the second lithiation, a net difference was observed compared to the first lithiation: the high background current originating from the current collector/electrolyte decomposition was no longer observed and only a broad shoulder around 1.2 V was visible at potentials above 1 V. The peak at 0.6 V showed the same onset as in the first cycle, while the peak at 0.25 V could be deconvoluted into two peaks at 0.3 V and 0.22 V. During the second delithiation, an extra peak appears at 0.65 V, while the peak at ~1.1 V showed a higher intensity than in the first lithiation.


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In the following 3rd, 5th and 10th cycle, the main differences come from the splitting of the peak at 0.6 V during lithiation and the continuous fading of the peak at 0.25 V. This peak evolution suggests that the reaction mechanisms and Li pathways are different from one cycle to another and that new phases develop during cycling. During further lithiation, the main changes appear at 0.6 V, where the peak intensity increases followed by the formation of a new peak at 0.7 V and the decrease in intensity of the process localized between 1.3 and 2.0 V. Comparing the cycling of the SnO2 thin film with that of a metallic Sn thin film (Figure S2, Supporting Information) showed that the CV of SnO2 clearly became more and more similar to that of Sn as the cycle number increased. Such behavior suggests that the reaction mechanism occurs stepwise, evolving from an initial conversion/alloy reaction to predominantly an alloy reaction, and that SnO2 is less and less recombined. Figure 2b shows the first galvanostatic cycle profile, which is in agreement with the data obtained using cyclic voltammetry. Several key potentials (labeled from A to I in Figure 2b) have been selected to be characterized by postmortem XPS and SEM to investigate properly the complex and multiple electrochemical reactions taking place during the first cycle.


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Figure 2. (a) Cyclic voltammogram of a 20 nm SnO2 film on Cu cycled in LP30 electrolyte vs. Li metal at a rate of 50 µV/s; (b) first galvanostatic cycle measured at 1.6 C-rate. Key potentials examined by post-mortem XPS and SEM are marked in red during lithiation and blue during delithiation.

3.3 Post-mortem SEM Post-mortem SEM images displayed in Figure 3 were collected for an electrode kept at OCP and compared to those of electrodes extracted after cycling. For the sample measured at OCP (~ 2.9 V), the film appeared to be very similar to the pristine one (Figure 1a). After the cell was cycled and then stopped at 2.0 V, small morphological changes, such as the initially rough surface, were observed. However, most important modifications were detected at 1.4 V, with wrinkle-like defects appearing along the stripes of the Cu substrate coupled to several cracks. At 0.81 V, the film appears to be completely wrinkled and partially lifted from the substrate as consequence of volume change due to the electrochemical reactions. Indeed, between 2 V and 0.8 V, Li2SnO3 and Li8SnO6 phases are supposed to be formed via Reaction (III) and (IV) with


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an average volume of 3.3 and 1.8 times bigger, respectively, than the volume of SnO2 (see Table S1, Supporting Information). Similar features with more pronounced textures and denser wrinkles are observed at 5 mV, and relate to the LixSn alloy reactions according to Reaction (II). During delithiation, at 2.0 V, the surface morphology does not recover to the initial flat situation. Although we expected electrolyte decomposition products to cover the surface of the film, no clear features were detected by SEM.

Figure 3. Post-mortem SEM images of 20 nm SnO2 thin films at different key potentials: during open-circuit potential (OCP), 1st lithiation (2.0 V, 1.4 V, 0.81 V, full lithiation) and full delithiation.

3.4 Post-mortem XPS during the first lithiation XPS analyses of the Sn3d5/2, C1s, and O1s core levels that were acquired during the first lithiation are presented in Figure 4, and the F1s and Li1s core levels are shown in Figure S3 (Supporting Information). The values of BE and FWHM for the Sn3d5/2, C1s, and O1s core


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levels are reported in Table S2 (Supporting information). The absolute intensities related to the area under the peaks of the acquired core levels are also presented in Figure S4 (Supporting Information). After resting at OCP for 48 h, no obvious change was observed in the Sn3d5/2 core level compared to the pristine except a slight decrease in the intensity followed by a slight intensity increase in the C1s and O1s peaks. The presence of fluorine and phosphorus traces was also detected at this stage. The latter originates from the adsorbed DMC and EC species as well as from the remaining LiPF6 salt and LiF formation. At 2.0 V, no major changes were detected, except the continuous increase of the LiF component in the F1s spectrum. The first noticeable changes were observed at 1.4 V, which coincides with the initial electrochemical activity detected by CV measurement and SEM image (Figure 3, cracks and beginning of wrinkles-like formation). We notice the appearance of a second component in the Sn3d5/2 at 485.4 eV, which corresponded to Sn0, followed by a net intensity decrease of the peak. Simultaneously, no significant variation was found for the C1s and O1s peaks. This behavior is believed to be related to the initial lithiation of the SnO2 film through Reaction (III), leading to the formation of a Li2SnO3 phase. It is challenging to distinguish the SnO2, SnO, and Li2SnO3 phases in Sn3d5/2 core level XPS as the BE difference between the three components is very small. Thus, we used DFT calculation to clarify this point (see part 3.7). Based on experimental and DFT results, we confirmed that the conversion from SnO2 to Sn0 starts with Reaction (III), followed by Reaction (IV). Additionally, we noticed that, at 1.4 V, only LiPF6 was reduced to form LiF (intensity doubled) (Figure S3 and S4 in Supporting Information), whereas the DMC and EC solvents remain relatively stable as confirmed by the C1s core level (Figure 4).


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From 1 V to 0.8 V, drastic changes occur in the C1s peak, with the appearance of Li2CO3 species at 291.4 eV and three additional components located at 286.5 eV, 288 eV, and 289.9 eV, which are associated with C-H, C-O, and C=O, respectively. These changes in the C1s peak originate from the decomposition of EC and DMC solvents. The intensity of the Sn3d5/2 and F1s peaks dropped significantly as a consequence of the formation of organic/inorganic species, supporting their homogenous spread over the SnO2 film and LiF. Additionally, the conversion Reaction (IV) evolves progressively and reaches its maximum at 0.8 V, where the intensity of the Sn0 component becomes comparable to that of Li2SnO3. Further changes occurred at 0.45 V and 5 mV, where the alloy reaction was expected to begin. In the Sn3d5/2 core level, three distinct components were detected. The first one at 487.6 eV is believed to be related to SnFx species because it is shifted to a higher BE position than the SnO2 component (see the Figure S5 in Supporting Information). The component at 485.5 eV is associated with the Li8SnO6 phase, while the third ones at 485.4 eV (for 0.45 V) and at 485 eV (for 5 mV) is affiliated to LixSn phases. It is also worth mentioning that the Sn3d5/2 peak intensity dropped by roughly half at 0.45 V and by a factor of 10 at 5 mV. Surprisingly, this intensity drop was not accompanied by an increase of the C1s, O1s or F1s, indicating the absence of any evolution of the organic/inorganic species originates from the electrolyte reduction. Simultaneously, a Li2O component is detected in the tail of the O1s peak at low BE at 530.8 eV (indicated by the arrow on the right side of Figure 4), and it constantly evolves to reach a maximum at 5 mV. Thus, the intensity attenuation of the Sn3d5/2 is basically related to the Li2O phase generated by Reaction (VI). The Li2O mainly grew around the Sn0 as confirmed by the strong attenuation of the Sn0 component in respect to the Li8SnO6 at 5 mV. At this stage, three reactions overlap between 0.45 V and 5 mV: Reaction (V) linked to the conversion of Li2SnO3


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into Li8SnO6 and Sn0, Reaction (VI) linked to the conversion of Li8SnO6 into Li2O and Sn0, and finally, Reaction (II) linked to the alloy of Sn into LixSn. The fact that at 5 mV the Li8SnO6 remains visible confirms that Reaction (VI) did not reach completion. 3.5 Post-mortem XPS during the first delithiation During the delithiation process at 0.9 V, the three components of SnFx, Li8SnO6, and LixSn within the Sn3d5/2 core level remained visible. The shift back to the initial position of the LixSn component at 484.6 eV was in agreement with the de-alloying reaction observed in the CV curve at 0.47 V. At this potential; the Sn3d5/2 peak intensity starts to increase again, implying the dissolution of the organic/inorganic species from the surface. In this situation, we clearly observed the partial decomposition of Li2CO3 in the C1s peak as well as that of Li2O in the O1s spectra. No significant changes were observed at 1.5 V except a slight decrease in the CH2 group intensity. The Li2O component is believed to be detected in the noise level. However, at the end of delithiation at 2 V, we notice the back-conversion of Li8SnO6 and Sn0 into SnOx. The Sn3d5/2 spectrum is almost recovered with just two components at 485.5 eV and 487.1 eV associated to SnOx and Sn0 respectively, indicating that the reaction mechanism is mostly reversible in agreement with the CV curves. At the same time, the intensity of the CH2 group keeps decreasing while Li2O component is hardly detected in the tail of the O1s peak.


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Figure 4. XPS spectra for Sn3d5/2, C1s and O1s core levels acquired on 20 nm SnO2 thin film at different potential during first lithiation and delithiation.

3.6 Long-term cycling and angle-resolved XPS analysis XPS analyses after long-term cycling were also performed in order to reach a better understanding about the reversibility of the conversion reaction between SnO2 and Sn0. Two samples were investigated after being stopped at 0.8 V during the second lithiation and stopped at 2 V after five complete cycles, respectively. The corresponding Sn3d5/2 core levels are presented in Figure 5 and can be directly compared to the samples stopped at 0.8 V during the


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first lithiation and to the sample stopped at 2 V at the end of the first delithiation. During the second lithiation at 0.8 V (Fig.5b), more Sn0 compared to that in the first cycle (Fig.5a) is detected by directly comparing the intensity ratio Li2SnO3/Sn0. At the end of the delithiation after five cycles (Fig.5d), less Sn0 was reversibly converted to SnO2 compared to that in the first cycle (Fig.5c) Such a result further confirms the partial reversibility of Reaction (I) from Sn0 to SnO2, as already observed after the first cycle and the domination of the alloy Reaction (II) along cycling. The electrochemical activity in Figure 2 and in Figure S2 (Supporting Information) are in line with the XPS showing fewer features related to the conversion process in the subsequent cycles.

Figure 5. XPS spectra comparison of Sn 3d5/2 core level acquired at 0.8 V during (a, b) first and second lithiation at 0.8 V and (c, d) first and fifth cycle after full delithiation at 2 V.

In parallel to the long-term surface analyses, XPS measurements were acquired in angle-resolved


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mode on the sample stopped at 0.8 V during the first cycle. The aim of this experiment was to perform non-destructive depth profile analysis within the first 8 nm from the surface by changing the take-off angle between the analyzer and the sample holder. Such a measurement on Sn3d5/2 allows the distribution of Li2SnO3 and Sn phases to be determined vertically with respect to the surface. The Sn3d5/2 peaks were collected with four different angles (0º, 45º, 60º and 75º) as presented in Figure 6. The ratio between the SnO2 and Sn0 components remained constant (Table S3, Supporting Information) independent of the take-off angle. This behavior confirms that the two phases are homogenously distributed vertically down to 8 nm (corresponding to almost 40% of the film thickness).

0°

45°

60°

75°

Intensity [a.u.]

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


492 490 488 486 484 482492 490 488 486 484 482 492 490 488 486 484 482

492 490 488 486 484 482

Binding energy [eV]

Figure 6. Angle-resolved XPS analysis of Sn3d5/2 at 0.8 V during the first lithiation.

3.8 Density functional theory calculation As mentioned above, experimentally the BE shift of the Sn3d5/2 core level between the two Sn oxides (SnO2 and SnO) is reported to be 0.6 eV with FWHM around 1.5 eV. Such a small chemical shift and large FWHM make the discrimination between the two oxides with XPS very challenging 25-26. The same applies to the lithiated phases of SnO2 such as Li2SnO3 and Li8SnO6 or lithiated Sn phases like LiSn and Li4.4Sn. Obtaining experimentally the exact BE of Sn3d5/2


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from reference materials is far from trivial because they are not offered commercially and most of them are metastable phases. In this context, DFT calculation can help to better assess the core level chemical shift and the assignment of the various phases. Seven optimized clusters representing the structure of Sn, SnO2, SnO, and their lithiated phases Li2SnO3, Li8SnO6, LiSn, and Li17Sn4, presented in Figure S5 (Supporting information), were investigated. The binding energy of Sn3d5/2 of the different systems was calculated using DFT within the local density approximation with Vosko, Wilk, and Nusair functional (VWN)27 implemented into the cluster code StoBe (LCAO, Gaussian basis sets) (described in Note 2 in Supporting info). The models used include Sn atom as a central atom surrounded by two shells of neighbors. The theoretical BEs of Sn3d5/2 electrons are presented in Table 1 referenced to vacuum and calculated as the total energy difference between the ground and the ionized core state. The theoretically obtained values of BE were adjusted by subtracting the sample work-function and by adding energy correction of 3 eV for Sn, which includes a relative energy correction and orbital relaxation correction28. The effective core potentials were used for all other Sn atoms in the certain model to localize the core hole on the particular Sn center. The absolute theoretical BEs are difficult to compare directly with the XPS experimental measurement because the DFT models do not capture all the physics. Namely, they include only two shells of neighbors and no infinite longrange order. The trend of the core level shift or the binding energy difference (BE) are more significant than the absolute BE values. Additionally, results obtained by DFT describe BEs of single Sn centre, while in experimental XPS analysis populations of centres with the same binding energy are detected. For the lithiated SnO2 the theoretical BE separation of the different phases (Table 1) are: ∆(SnO2-SnO) = 0.45 eV, ∆(SnO2-Li2SnO3) = 0.8 eV, and ∆(SnO2-Li8SnO6) = 2.18 eV


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The BE separation between the Sn0 and lithiated Sn are: ∆(Sn-LiSn) = 0.4 eV and ∆(Sn-Li17Sn4) = 2.05 eV Although the calculated shift between SnO2 and SnO is slightly underestimated (0.45 eV theoretically vs. 0.6 eV experimentally), the trend of Sn3d5/2 ∆BE follows the experimental observations. In the case of the Li2SnO3 and LiSn phases, the Sn 3d5/2 BE shift with respect to SnO2 (∆BE = 0.8 eV) and Sn0 (∆BE = 0.4 eV) is rather small and can hardly be experimentally discriminated. This difficulty is exacerbated if the two phases are sub-stoichiometric in Li, which can make the chemical shift even smaller29-30. However, for Li8SnO6 and Li17Sn4 phases, the BE shift is over 2 eV with respect to SnO2 (∆BE = 2.18 eV) and Sn (∆BE = 2.05 eV) respectively. Such shifts are significant and can be detected easily within the Sn3d5/2 peak even if the lithiated phases are sub-stoichiometric in Li. The BE shift in lithiated oxide systems has the tendency to decrease when the lithium content increases. It is the consequence of the Li being a donor of electrons to Sn, making the core electrons easier to be removed from the more lithiated system. The DFT calculation supports the XPS experimental observation, showing that the Sn3d5/2 core levels acquired on SnO2, SnO and Li2SnO3 phases overlap and are very difficult to differentiate. Similar observations were made for Sn0 and LixSn phases.

Table 1. Calculated Sn3d5/2 binding energy for Li-Sn-O systems Structure SnO2 SnO Li2SnO3 Li8SnO6 Sn LiSn Li17Sn4

BE [eV]LDA_VWN 486.83 eV 486.38 eV 486.06 eV 484.65 eV 484.64 eV 484.22 eV 482.59 eV

∆BE ∆(SnO2-SnO) = 0.45 eV ∆(SnO2-Li2SnO3) = 0.8 eV ∆(SnO2-Li8SnO6) = 2.18 eV ∆(Sn-LiSn) = 0.4 eV ∆(Sn-Li17Sn4) = 2.05 eV


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4. DISCUSSION To date, the electrochemical reaction mechanisms of SnO2 during lithiation have been assumed to consist of a conversion reaction leading to Sn and Li2O formation at ~1.2 V (Reaction I) followed by an alloy reaction of LixSn below 0.5 V (Reaction II). During delithiation, both Reaction (I) and (II) were reported to be reversible through the decomposition of Li2O and the reformation of SnOx phase as a result of the reaction of Sn0 16, 22. In the present study, we report new insights on the presence of intermediate reactions taking place during lithiation, changing the reaction mechanisms drastically. Indeed, Reactions (III), (IV), (V) and (VI), which were already predicted by theoretical calculations, have to be considered as well as their possible kinetic/thermodynamic limitations in bulk electrode. Thanks to the DFT calculations, we anticipated the behavior of the XPS Sn3d5/2 core level upon SnO2 and Sn lithiation, showing the difficulty in discriminating between SnO2, SnO and Li2SnO3 as well as Sn and LiSn. Thus, postmortem SEM pointed out a large volume change at early lithiation, giving an indication of the Li2SnO3 phase formation resulting from Reaction (III) between 2 V and 0.8 V. The combination of XPS supported by DFT calculations and SEM coupled with the cyclic voltammetry allowed the voltage dependency of the reaction mechanisms during the first cycle to be illustrated (Figure 7). From OCP until 0.8 V the electrochemical activity begins with LiPF6 reduction at 1.4 V leading to a strong increase of the LiF species. Simultaneously, the conversion Reaction (III) is observed below 2 V with the appearance of Li2SnO3. Afterwards, between 1.4 V and 0.8 V, Reaction (III) overlaps with Reaction (IV) forming both Li2SnO3 and Sn0. The detection of Reaction (III) and (IV) during the early lithiation are in agreement with the predicted theoretical calculations17 expected to occur at around 2 V. However, the reduction of DMC and EC solvents starts at 1 V covering the surface of the electrodes with Li2CO3 and oligomers-like species.


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Three more additional reactions were detected between 0.45 V and 5 mV: Reaction (V) (the conversion of Li2SnO3 into Li8SnO6 and Sn0), Reaction (VI) (the conversion of Li8SnO6 into Li2O and Sn0), and Reaction (II) (the alloying of Sn into LixSn). Those reactions overlap in these potential ranges, as both Li8SnO6 and Li2O were detected at 0.45 V as well as the alloying of Sn0. The formation of Li2SnO3 and Li8SnO6 phases was expected in this study because the same behavior was already observed for Si particles and thin films19, 23, where the conversion of the native SiO2 to lithium silicate (Li4SiO4) was reported to occur at 0.2 V and overlaps with the formation of the Si alloy with lithium. An operando X-ray diffraction measurement was performed on 100 nm SnO2 film (Figure S7, Supporting Information) and confirms that all the phases formed during the lithiation are amorphous as expected from the amorphous nature of the pristine film. The XPS depth profile analysis confirmed the absence of any preferential surface passivation since both Sn and LixSnOy phases were observed to be homogeneously distributed in the film. Otherwise, in agreement with previous TEM studies9, the Li2O generated from Reaction (VI) was found to cover mainly the Sn0 particles, and its thickness reached its maximum at 5 mV. Concerning electrolyte decomposition, no clear further surface evolution was observed below 1 V. During delithiation, the reactions are mostly reversible. At 0.9 V, the partial decomposition of Li2O and Li2CO3 species was detected, leading to their dissolution. However, no major changes occurred in the Li8SnO6 phase while the LixSn phases start to be delithiated. Similar behavior was observed at 1.5 V, where no oxidation of Sn0 started yet despite the Li2O component decreasing in intensity such that it is only detected in the noise level. The oxidation of Li8SnO6 and partial oxidation of Sn to SnOx happened at 2 V above the potential of the Li2O decomposition. This behavior led us to believe that the mechanism and the origin of the Sn


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oxidation might not be only related to Li2O decomposition. The contribution of the electrolyte as a source of oxygen to oxidize the Sn0 cannot be confirmed nor excluded at this stage. The decomposition of Li2CO3 can also be exempt since the component in the C1s did not evolve between 1.5 V and 2 V. Finally, after further cycles, less Sn0 is reversibly converted to SnOx, and the alloy reaction dominates the electrochemical mechanism leading to a steady fade in the specific charge.

Figure 7. Illustration of the proposed reaction mechanism during first lithiation and delithiation.

CONCLUSION


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We have demonstrated the benefit of investigating amorphous thin film model electrodes to dissect the electrochemical reaction mechanisms of conversion-based materials for Li-ion batteries. Amorphous 20 nm SnO2 thin films analyzed by coupling electrochemistry, postmortem XPS and SEM supported by DFT calculation have revealed the potential dependency of the various reaction occurring during the early stages of cycling as well as after long-term cycling. We demonstrated that the electrolyte reduction occurs in two steps, at 1.4 V and 1 V with the reduction of the LiPF6 salt and EC/DMC, respectively. Regarding the SnO2 film, we showed the presence of partially reversible intermediate reactions during the lithiation through the formation of Li2SnO3 and Li8SnO6 phases. The presence of Li2O is also evidenced below 0.5 V and found to selectively cover the Sn particles. During the delithiation process, the decomposition of the Li2O followed by its dissolution took place at 0.9 V in parallel with the LixSn de-alloying reaction. The onset of SnOx phase re-appears only above 1.5 V making the mechanism of the Sn oxidation potentially not only related to Li2O decomposition. Finally, the long-term cycling investigation demonstrates that the alloy reaction becomes dominant leading to specific charge fading. These findings advance the understanding of conversion-based materials and provide an important comprehension and perspective of such complicated reaction mechanism during battery operation. Furthermore, they are of valuable support to design the cycling protocols of SnO2 electrodes or graphite electrodes with SnO2 additives in order to optimize their electrochemical performances. By tuning the cell cut-off voltages the SnO2 can cycle reversibly through conversion, alloy or both conversion and alloy reactions providing certain tolerance in minimizing the impact of the volume expansion. In the same manner, fundamental knowledge is gained in this study relative to the analysis of alloy and conversion materials by using the XPS. Such knowledge can be applied later on much more complex


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commercial-like electrodes to better assess the various reactions taking place at the interface electrode/electrolyte.

ASSOCIATED CONTENT Supporting Information. Volume of the different Li-Sn-O compounds; XPS of Sn 3d5/2 on Sn thin film; CV of SnO2 vs. Sn thin films; XPS of Li 1s and F 1s core levels; Binding energy and Full-Width-Height-Maximum values of Sn 3d5/2, C 1s and O 1s core levels; XPS of Sn 3d5/2 comparison between LP30 and LC30 electrolyte at 5mV during 1st cycle. Operando XRD experiment. AUTHOR INFORMATION Corresponding Author Mario El Kazzi: [email protected] Funding Sources Centre of Competence for Energy and Mobility and Swiss Electric Research (project nr. 911) are gratefully acknowledged for financial support. REFERENCES [1] Nitta, N.; Wu, F.; Lee, J. T.; Yushin, G., Li-ion battery materials: present and future. Mater. Today 2015, 18 (5), 252-264. [2] Winter, M.; Besenhard, J. O., Electrochemical lithiation of tin and tin-based intermetallics and composites. Electrochim. Acta 1999, 45 (1–2), 31-50. [3] Besenhard, J. O.; Yang, J.; Winter, M., Will advanced lithium-alloy anodes have a chance in lithium-ion batteries? J. Power Sources 1997, 68 (1), 87-90. [4] Oehl, N.; Schmuelling, G.; Knipper, M.; Kloepsch, R.; Placke, T.; Kolny-Olesiak, J.; Plaggenborg, T.; Winter, M.; Parisi, J., In situ X-ray diffraction study on the formation of [small


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alpha]-Sn in nanocrystalline Sn-based electrodes for lithium-ion batteries. CrystEngComm 2015, 17 (44), 8500-8504. [5] Sivashanmugam, A.; Kumar, T. P.; Renganathan, N. G.; Gopukumar, S.; WohlfahrtMehrens, M.; Garche, J., Electrochemical behavior of Sn/SnO2 mixtures for use as anode in lithium rechargeable batteries. J. Power Sources 2005, 144 (1), 197-203. [6] Mohamedi, M.; Lee, S.-J.; Takahashi, D.; Nishizawa, M.; Itoh, T.; Uchida, I., Amorphous tin oxide films: preparation and characterization as an anode active material for lithium ion batteries. Electrochim. Acta 2001, 46 (8), 1161-1168. [7] Brousse, T.; Retoux, R.; Herterich, U.; Schleich, D. M., Thin‐Film Crystalline SnO2‐ Lithium Electrodes. J. Electrochem. Soc. 1998, 145 (1), 1-4. [8] Reddy, M. V.; Linh, T. T.; Hien, D. T.; Chowdari, B. V. R., SnO2 Based Materials and Their Energy Storage Studies. ACS Sustainable Chemistry & Engineering 2016, 4 (12), 62686276. [9] Lee, S.-Y.; Park, K.-Y.; Kim, W.-S.; Yoon, S.; Hong, S.-H.; Kang, K.; Kim, M., Unveiling origin of additional capacity of SnO2 anode in lithium-ion batteries by realistic ex situ TEM analysis. Nano Energy 2016, 19, 234-245. [10] Rhodes, K. J.; Meisner, R.; Kirkham, M.; Dudney, N.; Daniel, C., In Situ XRD of Thin Film Tin Electrodes for Lithium Ion Batteries. J. Electrochem. Soc. 2012, 159 (3), A294-A299. [11] Zhou, X. S.; Dai, Z. H.; Liu, S. H.; Bao, J. C.; Guo, Y. G., Ultra-Uniform SnOx/Carbon Nanohybrids toward Advanced Lithium-Ion Battery Anodes. Adv. Mater. 2014, 26 (23), 39433949. [12] Yin, Y. X.; Xin, S.; Wan, L. J.; Li, C. J.; Guo, Y. G., SnO2 hollow spheres: Polymer bead-templated hydrothermal synthesis and their electrochemical properties for lithium storage. Sci. China-Chem. 2012, 55 (7), 1314-1318. [13] Böhme, S.; Edström, K.; Nyholm, L., On the electrochemistry of tin oxide coated tin electrodes in lithium-ion batteries. Electrochim. Acta 2015, 179, 482-494. [14] Kilibarda, G.; Szabó, D. V.; Schlabach, S.; Winkler, V.; Bruns, M.; Hanemann, T., Investigation of the degradation of SnO2 electrodes for use in Li-ion cells. J. Power Sources 2013, 233, 139-147. [15] Hu, R.; Chen, D.; Waller, G.; Ouyang, Y.; Chen, Y.; Zhao, B.; Rainwater, B.; Yang, C.; Zhu, M.; Liu, M., Dramatically enhanced reversibility of Li2O in SnO2-based electrodes: the effect of nanostructure on high initial reversible capacity. Energy & Environmental Science 2016, 9 (2), 595-603. [16] Böhme, S.; Philippe, B.; Edström, K.; Nyholm, L., Photoelectron Spectroscopic Evidence for Overlapping Redox Reactions for SnO2 Electrodes in Lithium-Ion Batteries. The Journal of Physical Chemistry C 2017, 121 (9), 4924-4936. [17] Cheng, Y.; Nie, A.; Gan, L.-Y.; Zhang, Q.; Schwingenschlogl, U., A global view of the phase transitions of SnO2 in rechargeable batteries based on results of high throughput calculations. Journal of Materials Chemistry A 2015, 3 (38), 19483-19489. [18] Jain, A.; Ong, S. P.; Hautier, G.; Chen, W.; Richards, W. D.; Dacek, S.; Cholia, S.; Gunter, D.; Skinner, D.; Ceder, G.; Persson, K. A., Commentary: The Materials Project: A materials genome approach to accelerating materials innovation. APL Materials 2013, 1 (1), 011002. [19] Ferraresi, G.; Czornomaz, L.; Villevieille, C.; Novák, P.; El Kazzi, M., Elucidating the Surface Reactions of an Amorphous Si Thin Film as a Model Electrode for Li-Ion Batteries. ACS Applied Materials & Interfaces 2016, 8 (43), 29791-29798.


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[20] Philippe, B.; Dedryvère, R.; Gorgoi, M.; Rensmo, H.; Gonbeau, D.; Edström, K., Improved Performances of Nanosilicon Electrodes Using the Salt LiFSI: A Photoelectron Spectroscopy Study. J. Am. Chem. Soc. 2013, 135 (26), 9829-9842. [21] Courtney, I. A.; Dahn, J. R., Electrochemical and In Situ X‐Ray Diffraction Studies of the Reaction of Lithium with Tin Oxide Composites. J. Electrochem. Soc. 1997, 144 (6), 20452052. [22] Böhme, S.; Edström, K.; Nyholm, L., Overlapping and rate controlling electrochemical reactions for tin(IV) oxide electrodes in lithium-ion batteries. J. Electroanal. Chem. 2017, 797, 47-60. [23] Philippe, B.; Dedryvère, R.; Allouche, J.; Lindgren, F.; Gorgoi, M.; Rensmo, H.; Gonbeau, D.; Edström, K., Nanosilicon Electrodes for Lithium-Ion Batteries: Interfacial Mechanisms Studied by Hard and Soft X-ray Photoelectron Spectroscopy. In Chem. Mater., American Chemical Society: 2012; Vol. 24, pp 1107-1115. [24] Pereira-Nabais, C.; Światowska, J.; Chagnes, A.; Ozanam, F.; Gohier, A.; Tran-Van, P.; Cojocaru, C.-S.; Cassir, M.; Marcus, P., Interphase chemistry of Si electrodes used as anodes in Li-ion batteries. Appl. Surf. Sci. 2013, 266, 5-16. [25] Stranick, M. A.; Moskwa, A., SnO2 by XPS. Surf. Sci. Spectra 1993, 2 (1), 50-54. [26] Stranick, M. A.; Moskwa, A., SnO by XPS. Surf. Sci. Spectra 1993, 2 (1), 45-49. [27] Vosko, S. H.; Wilk, L.; Nusair, M., Accurate spin-dependent electron liquid correlation energies for local spin density calculations: a critical analysis. Can. J. Phys. 1980, 58 (8), 12001211. [28] Takahashi, O.; Pettersson, L. G. M., Functional dependence of core-excitation energies. The Journal of Chemical Physics 2004, 121 (21), 10339-10345. [29] Guittet, M. J.; Crocombette, J. P.; Gautier-Soyer, M., Bonding and XPS chemical shifts in ZrSiO4 versus SiO2 and ZrO2:Charge transfer and electrostatic effects. Physical Review B 2001, 63 (12), 125117. [30] Giustino, F.; Bongiorno, A.; Pasquarello, A., Modeling of Si 2p core-level shifts at Si– (ZrO2)x(SiO2)1−x interfaces. Appl. Phys. Lett. 2002, 81 (22), 4233-4235.


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TOC 155x95mm (120 x 120 DPI)


SnO2 Model Electrode Cycled in Li-Ion Battery Reveals the Formation

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