Interpreting PES Peak Positions for Cycled Li-ion Battery Electrodes

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Breaking Down a Complex System: Interpreting PES Peak Positions for Cycled Li-Ion Battery Electrodes Fredrik Lindgren, David Rehnlund, Ida Källquist, Leif Nyholm, Kristina Edstrom, Maria Hahlin, and Julia Maibach J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b08923 • Publication Date (Web): 16 Nov 2017 Downloaded from http://pubs.acs.org on November 19, 2017

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The Journal of Physical Chemistry

Breaking Down a Complex System: Interpreting PES Peak Positions for Cycled Li-ion Battery Electrodes Fredrik Lindgren†‡,, David Rehnlund†, Ida Källquist‡, Leif Nyholm†, Kristina Edström†, Maria Hahlin‡*, Julia Maibach† †

Department of Chemistry - Ångström Laboratory, Uppsala University, Box 538, 75121 Uppsala,

Sweden ‡

Department of Physics and Astronomy, Uppsala University, Box 516, 75120 Uppsala, Sweden

ABSTRACT

Photoelectron spectroscopy (PES) is an important technique for tracing and understanding the side reactions responsible for decreasing performance of Li-ion batteries. Interpretation of different spectral components is dependent on correct binding energy referencing and for battery electrodes this is highly complex. In this work, we investigate the effect on binding energy reference points in PES in correlation to solid electrolyte interphase (SEI) formation, changing electrode potentials and state of charge variations in Li-ion battery electrodes. The results show that components in the SEI have a significantly different binding energy reference point relative to the bulk electrode material (i.e. up to 2 eV). It is also shown that electrode components with electronically insulating/semi-conducting nature are shifted as a function of electrode potential relative to highly conducting materials. Further, spectral changes due to 1 ACS Paragon Plus Environment


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lithiation are highly depending on the nature of the active material and its lithiation mechanism. Finally, a strategy for planning and evaluating PES experiments on battery electrodes is proposed where some materials require careful choice of one or more internal reference points while others may be treated essentially without internal calibration.

INTRODUCTION In Li-ion batteries, important and critical reactions take place on the electrode surfaces and at interfaces to the electrolyte. The operation voltage of these electrodes is typically close to or even outside the stability window of the electrolyte, which leads to electrolyte decomposition. However, on the negative electrode, this process is self-limiting as reduction of the electrolyte solvents and salts forms a passivating layer. This is the so called solid electrolyte interphase (SEI) layer and the passivating properties of this interphase layer are essential to prevent extensive and continuous electrolyte decomposition.1 On the positive electrode side, oxidation of solvents, salt decomposition and dissolution of active materials and binders may decrease the battery performance.2 Altogether, the properties of interfaces and interphase layers in the Li-ion battery influence to a large extent parameters such as safety, capacity loss, rate capability and cycle life.3,4 One widely adopted technique to study surfaces and interfaces in electrode materials is photoelectron spectroscopy (PES) because of its sensitivity to chemical environments/chemical shifts and as the redox state can be followed. PES can also provide information on elemental compositions, relative amounts, and the depth distribution of the species analyzed.5 The chemical shift is used when assigning certain atoms to their nearest neighbors and relative peak positions are therefore essential.6 However, also other mechanisms than a change in chemical shifts can affect the spectra and the measured binding energies which makes data interpretation complex.7 Thus, as a basis for conclusive data interpretation, accurate binding energy calibration of spectra is required. 2 ACS Paragon Plus Environment

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Although PES is a well-established surface analysis technique, new material classes and emerging complex structures might call for a re-evaluation of the binding energy calibration practices to achieve consistent data interpretation. The composite electrodes used in modern Liion batteries are good examples of this: they pose many challenges with respect to calibration as they comprise materials with different properties, a three-dimensional structure, lithiation and de-lithiation reactions and also formation of new interphase layers during cycling. To demonstrate this complexity, a SEM/EDX image of a Li-ion battery electrode comprising different electrode materials, a polymeric binder and carbon black (CB) is shown in Figure 1(a) and a schematic cross-section view of such an electrode is presented in Figure 1(b). This electrode, although not of commercial interest, is used as a model system in this work where different battery materials are investigated within the same electrode.

(a)

(b)

Binder

Active materials

Carbon black

Current collector

Figure 1. (a) SEM top view of a composite electrode for Li-ion batteries with different active materials highlighted in different colors (Li4Ti5O12 (LTO) – yellow, Li2FeSiO4 (LFS) – blue,

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LiFePO4 (LFP) – green and LiMn2O4 (LMO) – pink, carbon black (CB) and particles yielding insufficient elemental contrast particles - grey). (b) shows a simplified cross-section view of a composite electrode containing different active materials, a binder, CB and a current collector. The challenge faced in the energy calibration of these measurements has attracted some attention considering the electrochemical potential in intercalation battery materials8 and the energy scale referencing for Li metal used as anode material.9 However, more work is required to decode all the information in PES spectra from Li-ion battery electrodes to understand the whole system in more detail. During our work with PES analyses of various electrode materials, we have encountered a number of different factors influencing the measured binding energies.10-12 These include surface layer formation, changed electrode electrochemical potential as well as the expected changes in chemical shift due to different lithiation and delithiation mechanisms. Since more than one of these effects may occur in samples at different stages during the electrochemical cycling, it is challenging to establish the reason for an observed peak shift. Consequently, the determination of peak positions on the binding energy scale, both on a relative and on a calibrated scale, can be challenging. In a previous article,13 we proposed a model based on a potential gradient at the buried interface (i.e. between the SEI layer and the active electrode material) in order to explain relative shifts in measured binding energies. In this model, the active material is in good contact with the ground of the spectrometer, thus ensuring that the bulk sample and the spectrometer are at the same potential. SEI components and binders are essentially electronically insulating and the local potential of these compounds may be different, resulting in relative peak shifts of the corresponding spectra. However, the exact chemical/physical origin of this potential gradient is not yet fully understood.


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In this work, composite electrodes and model systems are examined to estimate the influence of the spectral changes due to: SEI formation, changing electrode electrochemical potential and state of charge variations. Based on the results, we propose guidelines on how to interpret PES results for cycled battery electrodes.

EXPERIMENTAL SEI formation on gold. Half-cells with Au-disc working electrodes (Goodfellow, 99.9%), Li-foil counter electrodes and two layers of Solupor® separators (thickness 10 µm) soaked in 1 M LiPF6 in ethylene carbonate (EC) (BASF):diethyl carbonate (DEC) (BASF) (2:1 by volume) were assembled in so-called pouch-cells inside an argon filled glovebox. SEI formation was achieved by cyclic voltammetry (CV) at 0.1 mV s-1 from the initial open circuit voltage (OCV) (~3 V vs. Li+/Li) to 0.4 V vs. Li+/Li and then by different number of cycles between 0.4 and 1.0 V vs. Li+/Li, also at 0.1 mV s-1. After the desired number of cycles, all samples were stopped at 1.0 V vs. Li+/Li and electrodes were taken out for PES analysis. One sample was also swept from the initial OCV (~3 V vs. Li+/Li) to 1.0 V vs. Li+/Li where it was stopped and analyzed, this sample is referred to as “Start of SEI formation”. The CV was performed on a BioLogic MPG 2 potentiostat. Mixed materials electrode. Mixed materials electrodes were prepared by ball milling (Retsch PM4) a mixture of Li4Ti5O12 (LTO), carbon coated Li2FeSiO4 (LFS) (Höganäs), carbon coated LiFePO4 (LFP) (Phostech P2), LiMn2O4 (LMO) (SP30, Merck), CB (Erachem Comilog Super P) and PVdF-HFP binder (Kynarflex 2801) dissolved in N-Methyl-2-pyrrolidone (NMP)(VWR). The mixture was bar coated on a carbon coated Al-foil and dried in air. The dry electrode coating comprised (by weight) 20% LTO, 20% LFS, 20% LFP, 20% LMO, 10% CB and 10% PVdF-HFP. Electrodes with a diameter of 20 mm were punched out using a Hohsen precision punching tool, brought into an argon filled glovebox and dried at 120 °C for six hours.



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Half-cells were prepared with Li-foil counter electrodes and two layers of Solupor® separators (thickness 10 µm) soaked in 1 M LiPF6 in EC:DEC (BASF) (2:1 by volume) sealed in pouchcells. The average cell capacity was estimated to 150 mAh g-1 (approximate practical capacity of LFP: 165 mAh g-1, LMO: 120 mAh g-1, LTO: 175 mAh g-1, LFS: 120 mAh g-1) and the mass of the electrode coating was about 1.8 mg cm-2. Electrochemical de-lithiation and lithiation was performed on a Digatron BTS 600 galvanostat with a constant current corresponding to a rate of C/10. The upper and lower cut-off values during the cycling were set to 1.3 and 4.5 V vs. Li+/Li respectively. Before PES analysis, the half-cells were stopped in the middle of the characteristic voltage plateaus for the individual active materials and were allowed to relax to achieve a steady OCV. After the PES measurements the electrodes were re-assembled in identical half-cells. The surface morphology and elemental composition of a pristine composite electrode was analyzed using a Merlin high resolution scanning electron microscope (Zeiss) equipped with an energy dispersive X-ray (EDS) detector (Oxford Instruments). The EDS spectra were collected at an acceleration voltage of 15 kV. LiMn2O4 (LMO) electrodes. LMO (SP30, Merck), CB (Super P, Erachem Comilog) and PVdF-HFP (Kynarflex 2801) were dissolved in NMP (VWR) and mixed for one hour (Retsch PM4). The mixture was bar coated onto carbon coated Al-foil and dried in air. The dried electrode contained (by weight) 80% LMO, 10% CB and 10% PVDF-HFP and had a mass loading of 2.4 mg cm-2. 20 mm diameter electrodes were punched out and dried in the glovebox as described previously. Half-cells were assembled with a Li-foil counter electrode and two layers of a Solupor® separator (10 µm thick) soaked in 1 M LiPF6 in EC:DEC (BASF)(2:1 by volume) in pouch cells. The half-cells were galvanostatically cycled with a current corresponding to a rate of C/10 between 4.5 and 3.0 V vs. Li+/Li to achieve the desired degree of lithiation using a Digatron BTS 600 galvanostat.


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Nano Si composite electrodes. A slurry containing powders of Si nanoparticles (CEA, 100200 nm), CB (Erachem Comilog Super P) and Na-carboxymethyl cellulose (CMC) (MW=700 000 g mol-1, D.S = 0.90, Sigma-Aldrich) were mixed together using a buffered water solution (pH 2.8) containing citric acid (Sigma-Aldrich, 99,5%) and KOH (Eka Nobel) according to a procedure described by Mazouzi et al.14 The slurry was prepared by ball milling in a planetary ball mill (Retsch PM4) for one hour and the slurry was thereafter bar coated onto Cu-foil. Coated foils were dried in air at 60°C and the dry weight composition of the composite film was 67% Si, 10% CB, 7% CMC, 14% citric acid and 2% KOH. The average mass loading of Si was 1.0 mg cm-2. Electrodes and half-cells were prepared as described above although in this case the electrolyte consisted of 0.6 M lithium 4,5-dicyano-2-(trifluoromethyl)imidazolide (LiTDI) in dimethyl carbonate (DMC)(BASF):EC(BASF):fluoroethylene carbonate (SigmaAldrich):vinylene

carbonate

(BASF)

(2:1:0.1:0.02

by

volume).

Half-cells

were

galvanostatically lithiated with at a C-rate of C/75 (i.e. -48 mA per g of Si assuming a theoretical capacity of 3579 mAh g-1) until the desired charge capacity or the lower cut-off voltage (0.01V vs. Li+/Li) was reached, using a Digatron BTS 600 galvanostat. Reference samples for the Si electrodes were prepared using undoped (100) Si wafers which were allowed to react with Li. The polished (100) side of a Si wafer was put in contact with a Li-foil, sealed in a pouch cell under argon, and stored for one week at 60°C. A sample where a Li-foil was put in contact with the unpolished side (backside) of an identical (100) Si wafer was also prepared the same way. Third, a pristine (100) Si wafer sample was prepared analogously, but without any Li-foil, and this sample was likewise stored at 60°C for one week.

PES analysis. PES sample preparation of the cycled half-cells was performed in an argon filled glovebox (H2O ~ 1 ppm, O2 ~ 1 ppm). The working electrodes were removed from the pouch, rinsed in DMC (BASF) and mounted on the sample holder using conductive Cu-tape.



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One gold electrode with an SEI (stopped after one CV cycle) was also mounted on nonconductive tape and thereby analyzed without being connected to the ground of the spectrometer (this will be referred to as a floating sample). In this case a low energy electron flood gun was used for charge compensation. The flood gun parameters were adjusted to achieve as steady and narrow peaks as possible and this was achieved at a current of 23 mA with the electron energy set to 20%. Pouches with Si-wafers were opened inside the glovebox and each wafer was mounted on sample holders using Cu-tape and a screw to ensure good electrical contact; the Si wafers were hence not rinsed in DMC. All samples were transferred to the analysis chamber with a special transfer unit to avoid exposure to air and moisture.15 Samples that were re-assembled in pouch cells were also brought back to the glovebox using the same transfer unit without exposure to the atmosphere. All measurements, except the analyses of the Si electrodes, were performed with a Perkin Elmer PHI 5500 using monochromatic Al Kα radiation with an excitation energy of 1486.7 eV. Silicon samples at different stages during the first lithiation were analyzed by HAXPES with an excitation energy of 6015 eV. The HAXPES measurements were performed at the HIKE end station at the KMC-1 beamline16,17 at the synchrotron facility BESSY II operated by the Helmholtz-Zentrum Berlin. Data analysis and curve fitting was performed using the Igor Pro 4.07 software. For the LMO electrodes and the mixed electrode material no binding energy calibration was employed, i.e. the corresponding spectra are presented as measured. The peak parameters established by Nesbitt and Banerjee18 (based on multiplet splitting for Mn2+, Mn3+, and Mn4+) were used for the Mn2p3/2 peak fitting for the pure LMO electrode. For the SEI on gold samples, the spectra are presented both as-measured and calibrated using the surface hydrocarbon peak. The floating sample was energy calibrated in such a way that the Au4d5/2 peak is located at 335.1 eV, consistent with the grounded samples of the SEI on gold electrodes. The Si electrodes at various


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degrees of lithiation were referenced in binding energy either to the SiO2 peak or to the lithium silicate peak of the reference Si wafers; more details are provided in the results and discussion section.

RESULTS AND DISCUSSION The following sections present different examples of processes in Li-ion batteries electrodes that affects peak positions and the overall shape of the PES spectra. First, we demonstrate a model system where a buried interface dipole layer is built up during SEI formation. Secondly, we investigate how a change in the electrode electrochemical potential affects the relative peak positions. Thirdly, we show spectral changes as a function of lithiation (state of charge) for a well-known positive electrode material (LMO). Lastly, we present PES results for silicon, a possible future negative electrode material, obtained after lithiation and we discuss and apply an internal binding energy calibration to this system.

Binding energies affected by SEI formation. In our previous work we proposed that a dipole layer at the buried interface shifts the observed binding energies of the components in the SEI layer relative to the underlying electrode material.13 Similar effects have also been observed in other experimental and theoretical studies.19-21 On samples with thick SEI layers and in highly surface sensitive measurements essentially only products within this SEI layer are detected, and PE binding energy shift relative to the underlying materials can hence not be observed. With higher PE excitation energies or thinner SEI layers, the spectra will contain signals from both the SEI layer and the electrode material underneath, and the binding energy calibration becomes complicated due to the peak shifts induced by the dipole layer at the buried interface. In real battery electrode systems, it can be challenging to separate the effects on the binding energy originating from the dipole layer at the buried interface since during the SEI



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formation also chemical shifts from the lithiation reactions of the underlying electrode material may occur simultaneously. In order to only study the influence of the SEI formation effect on the binding energies relative to the underlying electrode, gold disc electrodes with differently thick SEI layers were analyzed by in-house PES. Gold has no native surface oxide22 that may undergo conversion reactions with lithium but as a gold-lithium alloy is formed at potentials below ~0.2 V vs. Li+/Li 23

the electrodes were only cycled between 0.4 and 1.0 V vs. Li+/Li using CV. To evaluate the

effect of the CV scans on the gold electrodes, peak fitting of the Au 4f7/2 spectra were examined in detail (see Figure S1 and Table S1 in the supporting information). The Au 4f7/2 peak position shifts slightly from 83.9 to 84.1 eV and the full width at half maximum increases from 0.80 to 0.87 eV over the six samples. This could indicate a slight change in the gold electrode. However, in comparison to the large shifts of the surface layer components, these changes are relatively small and it does not inflict the dipole at the buried interface. Hence, with this cycling procedure, an SEI layer could be formed (as the reduction of EC usually takes place at around 1 V vs. Li+/Li) without affecting the underlying electrode material by extensive lithiation. This means that only the effects of the dipole layer of the SEI should be influencing the spectra. It should be pointed out that all the cycled samples (“Start of SEI formation” to “150 cycles”) were stopped at the same potential (i.e. at 1.0 V vs. Li+/Li) so that any influence of electrochemical potential or accumulated charges at the interfaces should be the same for all cycled samples. The Au 4d5/2 – C 1s spectra for the gold samples with increasing SEI layer thicknesses are presented in Figure 2(a) and (b). These spectra are intensity normalized with respect to the most intense C 1s peak. The binding energy scales of the spectra in Figure 2(a) and (c) are presented as measured and the position of the Au 4d5/2 peak appears at 335.1 eV 24 for all samples. The relative intensity of the Au 4d5/2 compared to the C 1s peak decreases upon the CV cycling,


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indicating an increasing SEI layer thickness. In the C 1s spectra of the “start of SEI formation” sample, a new set of peaks appears at approximately 2 and 5 eV higher than the hydrocarbon peak. These are attributed to the formation of Li-alkyl carbonates and similar compounds, which are characteristic for SEI layers formed by carbonate based electrolytes.25 After the SEI formation, the C 1s spectra have very similar appearances but are shifted to higher binding energies as the SEI layer is increasing in thickness, see the detailed C 1s spectra in Figure 2(c). A total peak shift of about 2 eV is observed between the hydrocarbon peak in the pristine sample and the hydrocarbon peaks from the SEI after 50 and 150 cycles.



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Figure 2. Au 4d5/2 – C 1s spectra, (a) not calibrated and (b) calibrated vs. the hydrocarbon signal. Detailed C 1s spectra for the two calibration routines are presented in (c) and (d), respectively. These results are consistent with our proposed model where a dipole layer at the buried interface gives rise to a potential gradient between the electrode and the SEI layer, which results in a shift in the binding energy. Further, this potential gradient should extend only over the immediate interface region (i.e. the double layer region) and not over the whole SEI thickness, as in the latter case all SEI peaks would be asymmetric.26 The exact composition of this dipole layer is yet undetermined but it is reasonable to assume that species at the buried interface, such 12 ACS Paragon Plus Environment

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as LiF and Li alkyl carbonates with a strong dipole moment, could be responsible for the potential gradient and hence the observed peak shifts. The nature of the buried dipole layer was investigated further by analyzing a cycled gold sample with PES using both ground and flood gun for charge compensation. The resulting spectra are presented in Figure 3, where the floating sample was energy calibrated using the Au 4d5/2 peak. It is seen that the binding energy difference between the bulk Au 4d5/2 and the surface C 1s peaks from the SEI components are the same in both measurements. This clearly means that the potential gradient at the buried interface is the same irrespective of the charge compensation method used. This support the conclusion that the potential difference causing the relative peak shift between bulk and surface is located at the buried interface and that this dipole is present in the sample at all times once the SEI layer has been formed.

Figure 3. Au 4d5/2 – C 1s spectra obtained after the first CV scan. The black spectrum is measured from a sample connected to ground while the dotted red line denotes the corresponding spectrum obtained for a floating sample using flood gun charge neutralization.



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If the species in the SEI layer are to be identified further, some kind of binding energy calibration of these spectra is required. This calibration can, depending on the origin of the hydrocarbon species, be done by vacuum level alignment by positioning adventitious hydrocarbons at 284.8 eV27 or polymeric hydrocarbons at 285.0 eV.28 In this work 285.0 eV is chosen for the C 1s spectra presented in Figure 2(b) and (d). Note that the Au 4d5/2 peaks are now shifting in binding energy implying that this kind of calibration prevents a meaningful interpretation of the data of the bulk material. However, for the surface confined species in the SEI, it should be possible to follow the evolution of specific compounds. These include the lithium alkyl carbonate peak, now positioned at ~290 eV and a pronounced carbon oxygen peak at 287 eV. Based on these results it is obvious that, depending on the point of interest, different reference points should be used when analyzing different parts of the electrode. A surface hydrocarbon/polymer hydrocarbon energy calibration is well suited for the SEI components while the bulk material would benefit from the use of a bulk specific reference. In this example, where the sample is fairly simple (a gold substrate covered by an SEI) it is straightforward to decide whether a compound originates from the bulk or the SEI and thereby to apply a proper energy calibration. In more complex battery samples, the situation is, however, more challenging as it might not be possible to determine the spatial/depth position of a compound within the sample. Reliable information on the spatial distributions of battery electrode species can preferably be obtained by using different photoelectron excitation energies, as high energy argon ion sputter depth profiling is known to alter the SEI composition.5 Also, by stepwise SEI formation and frequent analysis of samples during the buildup of SEI layers, any gradual transformations may be followed. The gold-SEI model system clearly shows the effect of SEI formation on the relative binding energies. Since the differences in the binding energy for the non-calibrated and the surface


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calibrated spectra are as large as 2 eV, an incorrect choice of calibration could clearly lead to incorrect conclusions regarding the surface chemical composition of the sample.

Electrode potential effects on the binding energy. During charge and discharge, the lithium content in the active materials increases or decreases as a result of the changes in the electrochemical potential of the material. Along with this, the spectra are expected to change since the varying lithium content alters the chemical environment in the active material. However, also materials not directly involved in electrochemical lithiation/de-lithiation, such as the binder and SEI components, seem to experience relative peak shifts as a function of the electrochemical potential.11, 29 The aim of this part of the study is to investigate the relationship between an electrode’s electrochemical potential and the binding energy of components in a composite electrode in the presence of a minimal SEI contributions. To assess the effect of a varied electrode electrochemical potential, four different Li-ion battery electrode materials (LTO, LFS, LFP and LMO) were incorporated in a mixed materials electrode. A SEM image of this electrode with the different active materials highlighted in different colors is shown in Figure 1(a) (see also Figure S2 for SEM-EDS images). This mix of materials was chosen because all the components can be cycled with very little SEI formation (i.e. within the electrochemical stability window of the electrolyte), they all have pronounced lithiation plateaus that ensure stable electrochemical potentials, and they cover a wide potential range (from ~1.6 to 4.1 V vs. Li+/Li) which should facilitate studies of the impact of the electrochemical potential on the binding energies. The cycling curve for the mixed electrode and the open circuit voltage at which the electrodes are analyzed with PES are indicated in Figure 4(a). The cycling is stopped in the middle of each material’s plateau to ensure that the potential of the sample is stable and that the voltages that are given in Figure 4(a) are the OCV obtained after the relaxation period. The



OCV of re-assembled cells after the PES measurement differed less than 0.1 V from those prior to the PES measurements, indicating that the electrochemical potentials of the samples were not significantly affected by the measurement or their re-assembly into new cells. Sample A is analyzed in the potential region where LTO is active, sample B in the LFS potential region and sample C and D in the LFP and LMO potential regions, respectively.

(a) (b)

Cell voltage vs. Li+/Li [V]

5

(c)

(d)

(e)

D - 4.13V (LMO)

4

C - 3.42 V (LFP)

3

B - 2.95 V (LFS)

2 A - 1.58 V (LTO)

1 0

0

4

8 12 16 Cycling time [h]

20

(f) Relative peak shift [∆ eV]

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0.0 -0.4 CB -CF2- in PVdF -F in PVdF LiF Ti 2p3/2 Mn 2p3/2

-0.8 -1.2 -1.6 0

1

2

3

4

5

Electrode open circuit voltage vs. Li+/Li [V]

Figure 4. (a) Cell voltage vs. cycling time for a mixed electrode containing LTO, LFS, LFP and LMO as the active materials. (b) F 1s, c) Mn 2p3/2, (d) Ti 2p3/2 and (e) C 1s spectra of the mixed electrode at different electrode OCVs. (f) Relative peak positions compared to sample A as a function of electrode potential for the mixed electrode.


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Figure 4(b) to (e) presents the PES results for the mixed materials electrodes as obtained for the pristine electrode and at the four different potentials, respectively. All spectra are normalized to the highest intensity but no binding energy calibration is performed. In Figure 4(b) the F1s spectra shows a C-F peak from the PVdF binder (green) and a small feature ascribed to LiF (light grey). Figure 4(c) and (d) present the Ti 2p3/2 and Mn 2p3/2 signals from LTO and LMO respectively. Noticeable features in the C 1s spectra in Figure 4(e) include the -CF2- peak from the PVdF binder (highlighted in red) and the CB peak at ~284 eV (dark grey). Note that the contribution from the carbon coatings on LFP and LFS appear at the same binding energy as the CB conductive additive. Overall, the C 1s spectra remains fairly unaffected by the cycling which confirms that there is no extensive SEI formation during this relatively short period of cycling. In these spectra (b-e) there is a clear shift of the peaks assigned to PVdF, LiF and LTO towards lower binding energies when the electrode potential is increased. At the same time, the peak positions for CB and LMO are almost constant in binding energy when the sample potential is increased. It should be noted that the LMO material partly may undergo an irreversible phase transition below 3 V vs. Li+/Li.30 However, this does not seem to have a major impact on the cycling curve as the LMO characteristic plateaus between 4.0 and 4.2 V vs. Li+/Li are present also during the second charge (after the discharge below 3 V). In addition, also the peak shape of the Mn2p3/2 peak remain essentially unaffected during this relatively short cycling period, indicating that the LMO material is not affected to a major extent. It should also be mentioned that the LTO peak in the pristine sample and in sample B to D represents a de-lithiated material and in sample A the LTO material is halfway lithiated. The Li-content in LTO is in this case very important since a sharp increase in electronic conductivity of LTO is observed at a lithiation degree of more than 4%.31 Therefore the LTO material in sample A (lithiation degree ~50%) should be considered a good electron conductor while in



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sample B to D (lithiation degree well below 4%, see calculations in supporting information) LTO should be considered a poor electron conductor. This change in conductivity may be the reason for the peak shift observed between the pristine sample and sample A and also between sample A and B. Since the LTO material in sample B to D already is fully de-lithiated, the shift observed between these samples, however, require another explanation. Unfortunately, signals from LFS and LFP are not possible to detect in this electrode due to the carbon coating that decreases the intensity for these material specific emissions (Si 2p and Fe 2p). In Figure 4(f) the relative peak shifts obtained for the individual emissions with respect to sample A are presented as a function of the electrode potential. There is a clear trend that the binding energy for some compounds (the PVdF binder, LTO and LiF) in the electrode decreases with increasing electrode potential while other materials (CB and LMO) remain essentially unaffected. The peaks ascribed to PVdF, LTO and LiF are shifting between 1.1 and 1.6 eV over a potential difference of about 2.6 V. A common denominator for the materials that show a significant peak shift is that they have poor electronic conductivity while LMO and CB are materials with comparatively high electronic conductivity. Due to the different electronic conductivity, we propose that a potential gradient is formed also at the interface between the conductive and non-conductive parts in the composite electrodes. The results show that the influence of this gradient depends on the open circuit potential of the sample and on whether or not the phases in contact with each other can adopt to this potential via electrochemical reactions (i.e. de-/lithiation in the electrochemically active materials) or by filling of electronic states (as e.g. in carbon black). Components like the binder or already fully de-lithiated LTO cannot adjust to the electrode potential and thus a dipole layer at the interface to the conductive/electrochemically active parts is formed, in analogy with the buried interface effects seen during SEI formation discussed above.


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The phenomenological awareness of this effect is most valuable since many electrode materials for Li-ion batteries have poor electronic conductivities. In turn, this means that the peak shift between the lithiated and de-lithiated states also could be an effect of a change in the electrode’s electrochemical potential in addition to the chemical shift due to the change in the lithium content. Generally, this model electrode system clearly shows that the effect is large enough to considerably affect the data interpretation.

Degree of lithiation Li intercalation in LMO. To investigate what happens to active materials in detail, an LMO electrode is chosen as a model system to investigate the type of spectral response that can be expected due to changes in its lithium content (or state of charge). The LMO active material exhibits well defined redox chemistry and can be cycled with minor surface layer formation.30 In this experiment, LMO electrodes were cycled to different degrees of lithiation prior to analysis. In Figure 5(a), the spectra showing the Mn 2p3/2 photoemission lines are presented whereas the C 1s photoemissions are presented in Figure 5(b). In the case of the de-lithiated sample (top in Figure 5(a)) Mn should theoretically only be present in the +IV oxidation state if all Li is extracted. However, in the spectrum the presence of a shoulder at low binding energy indicates the presence of a considerable amount of Mn in the +III and +II oxidation states (as is confirmed by peak fitting), probably due to incomplete de-lithiation and possibly also due to a particle surface with different composition. With increased degree of lithiation, the spectral response follows the expected behavior where the Mn-ions are reduced and relatively more intensity originates from the +III and +II oxidation states. The spectra in Figure 5(a) and (b) are presented as measured (i.e. no internal energy calibration has been performed). For the LMO material, our results show that essentially new



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peaks are formed due to lithiation but also that the overall peak shifts nevertheless are quite small. This is in good agreement with the results in the previous example (see Figure 4) where no obvious peak shifts are observed for LMO in the mixed materials electrode despite the extended potential range. The C 1s peak position from the CB material (often used as reference) is rather constant; its binding energy differs only by 0.1 eV between the de-lithiated to the fully lithiated sample. Since CB is electrochemically inactive at the high potentials of LMO (around 3.5 to 4.5 V vs. Li+/Li) no change in the binding energy due to reactions with Li is expected. Thus, the binding energy referencing for both LMO and CB is rather straightforward in this example. A slightly more pronounced peak shift is, on the other hand, observed for the PVdF components (dotted line indicated -CF2- in Figure 5(b)) as these binding energies increase with increasing lithium content (i.e. decreasing electrode potential) in accordance with the behavior seen for the mixed materials electrode in the previous example.


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Figure 5. (a) Mn 2p3/2 and (b) C 1s spectra for LMO electrodes at three different degrees of lithiation. The binding energies are presented as measured.

Lithiation of silicon. In many composite battery electrodes, an internal binding energy scale calibration is required to provide meaningful studies of the changes in the binding energies of bulk material components. Based on our experience, graphite, NiTiOPO4 and Si are a few examples of such materials.10-12 It is, however, not fully understood why these systems need



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this data treatment but volume expansion effects and extensive SEI formation may be expected to influence the measurements. Silicon is one of the most challenging battery electrode materials due to the nature of its alloying reaction with lithium, the large volume variations during the cycling and the presence of pronounced SEI formation. The lithiation of crystalline silicon occurs within a rather limited potential window (i.e. 0.15 - 0.01 V vs. Li+/Li, see Figure S3) and therefore only minor electrode potential effects on the spectra can be expected. However, at these potentials SEI formation is ongoing and a surface hydrocarbon calibration method will therefore not provide a meaningful calibration approach for the bulk material. In the following example we instead present spectra obtained after the first lithiation and SEI formation for a composite nano-Si electrode where an internal binding energy scale calibration is applied based on different reference samples. To calibrate this series of measurements a reasonable starting point is an undoped Si wafer (spectrum in Figure 6(a)). This spectrum is presented as measured (i.e. no binding energy calibration) and contains a SiO2 signal from the native surface oxide as well as the bulk silicon signal. The Si 1s spectrum for the pristine nano-Si electrode in Figure 6(d) has a similar appearance and it therefore has a comparable build-up. This spectrum is energy calibrated in order to align the SiO2 peak from the nano-Si electrode with SiO2 peak from the Si wafer. The binding energy of the bulk silicon peak is sensitive towards doping32 (or impurities) and is therefore considered less suitable for energy calibration. Any eventual doping (or impurities) in the pristine nano-Si particles could therefore explain why the bulk peaks from the nano-Si particles has a slightly higher binding than the undoped Si-wafer. The final reference sample for this series is the lithiated Si-wafer (spectra in Figure 6(c)). Also these spectra are presented as measured and no binding energy calibration has been performed. The Si 1s spectrum contains a more intense oxide peak compared to the pristine Si-


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wafer, most likely due to the contribution from the native oxide on the lithium foil from the preparation of this sample. By comparing the spectrum from the fully lithiated nano-Si electrode (Fig. 6 (i) – calibrated vs. the oxide peak on the lithiated Si-wafer) to the lithiated Si-wafer it is noteworthy that they are very similar in both appearance of the peaks and their binding energies. This indicates that these two ways of synthesizing the Li-Si material yields comparable compounds. Therefore, the chemically lithiated Si wafer (Fig, 6c) is used for binding energy calibration of the fully electrochemically lithiated nano-Si electrode (Fig 6.i). In the bulk signal from the silicon materials it is possible to fit two peaks (dark grey and dark yellow in Figure 6(c) and (i)), which indicate that at least two different silicon surroundings exist. This is in good agreement with theoretical33 and experimental work,34-36 where even at high degrees of lithiation both Li-Si and Si-Si bonds are reported. For the cycled electrodes the most stable material (or most stable binding energy reference) is considered to be the lithiated silicon oxide. This is because the transformation of the SiO2 into the lithiated oxide begins already at potentials below 1.8 V vs. Li+/Li.37 This means that when the lithiation of bulk Si is about to start at 0.15 V vs. Li+/Li (see Figure S3) the SiO2 should essentially already have been transformed into lithiated silicon oxide. The spectra for the cycled Si electrodes with different degrees of lithiation (Figure 6(e) to (i)) are therefore calibrated by positioning the lithiated oxide on the nanoparticles at the same binding energy as for the lithiated oxide on the fully lithiated electrode, thus essentially following the lithiated Si wafer reference (where the front side had been in contact with a lithium foil, see Figure 6(c)). Further, it is worth noting that with this calibration the CB peak (see C 1s spectra in Figure 6(e) to (i)) is shifted to higher binding energies and this could be explained by a lithiation reaction of this material that is also expected at these potentials. This peak shift is similar to lithiated graphite where peak shifts of about 1 eV to higher binding energies have been reported due to lithiation.38-40



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With the calibration performed it is now possible to interpret the spectra from the lithiation of the nano-Si electrodes. The pristine electrode, as stated earlier, consists of a native SiO2 layer and a bulk Si peak. After SEI formation, the native oxide has been transformed into a lithiated oxide and the corresponding peak is thereby shifted to a lower binding energy. The bulk Si peak is shifted to a higher binding energy, probably due to a doping effect caused by small amounts of lithium atoms inserted into the Si structure. This is in good agreement with both experimental n-doping of silicon32, first principles studies of the lithium silicon material33 and also the spectra in Figure 6(b) where a Si wafer was doped with Li and a similar peak shift to higher binding energies was observed. After a lithiation to 700 mAh g-1 (see Si 1s spectrum in Figure 6(f)) a new peak, highlighted with yellow, is emerging at a lower binding energy than the bulk Si peak (grey). This yellow peak continues to increase in intensity when more lithium is inserted and therefore this peak is assigned to a more lithiated Si environment. Also along with increased lithiation, both the grey and the yellow peak are successively shifted to lower binding energies. It is likely that the overall shift indicates that both Si environments are influenced by the overall increase in the Li content. In the Si 1s spectrum in Figure 6, the overall increase in Li content is illustrated with successively darker grey and yellow colors for the bulk Si peaks. The wide variety of Li-Si and Si-Si constellations possible for this material and the less defined redox chemistry makes a more comprehensive peak assignment challenging. Nonetheless, the reference samples synthesized in different ways than through electrochemical reactions in batteries provide valuable references for the binding energy calibration of this system. To extend the understanding of the lithiation mechanism of silicon further, it would be interesting to make comparisons to PES analysis of well-defined crystalline Li-Si materials. However, these spectra are collected from realistic Li-ion battery electrodes and confirm that


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at least two different Si-environments co-exist during the entire lithiation and thereby supports both theoretical33 and experimental work.34-36

Figure 6. Si 1s and C 1s spectra recorded at 6015 eV excitation energy. (a) to (c) Differently treated Si wafers and (d) to (i), crystalline nano-Si electrodes at different stages of electrochemical lithiation.



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Towards a PES data evaluation strategy To summarize, three distinct phenomena affecting the shapes of the spectra and the PE peak positions in cycled Li-ion battery electrodes are presented. These are: i) the formation of a SEI layer, ii) changes in the electrode potential and iii) lithiation and delithiation of the material.

During the SEI formation peaks from compounds within the SEI are successively shifted relative to the peaks from the bulk electrode materials as a consequence of the dipole layer formation at the buried interface. Due to the high electronic conductivity for the LMO and gold electrodes, these materials can easily adapt to the spectrometer grounding potential and the bulk signals require no internal binding energy calibration to present a meaningful picture. It is also clear that peaks originating from species in the SEI should be internally calibrated e.g. with respect to the surface hydrocarbon peak or similar prominent peaks. The choice of binding energy calibration approach should be based on if the species of interest are located in the bulk of the electrode or the SEI layer. Peak shifts are also observed for insulating and semiconducting components in the mixed materials electrode as a function of the electrode potential. This effect needs to be considered for components with poor electronic conductivity and when comparing spectra obtained at different electrochemical potentials, especially if there is a large potential difference. Generally, this effect poses a greater challenge regarding energy calibration as the peak shift does not seem to be identical for all peaks. Frequent measurements at different points of the cycling curve are hence needed to assure that the binding energy shifts can be followed properly. The degree of lithiation of the LMO and silicon material evidently influences the relative amounts of the different oxidation states and the relative intensities of these peaks will consequently be altered due to the lithiation and de-lithiation reactions. Each material and its lithiation mechanism will determine the changes in the spectra and the use of reference samples 26 ACS Paragon Plus Environment

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is therefore recommended. In the LMO electrode it is possible to follow the oxidation states whereas in the silicon electrode the lithiation mechanism is less defined and more complex. In these examples it is also important to collect samples frequently to be able to follow the gradual changes properly.

CONCLUSIONS This work, which emphasizes the challenges associated with the interpretation of PES results obtained for composite battery electrode materials, contributes to the ongoing discussion on energy calibration of PES data for complex materials and improved PES data interpretation in general. Specifically, it was shown that relative binding energy shifts occur within composite electrodes depending on their state of charge, electrode potential and in the case of SEI formation. Altogether, these shifts may alter the peak position up to several eVs and if not accounted for, they could lead to misinterpretations and incorrect peak assignments. It is, on the other hand, also possible that these effects could be used to obtain information regarding e.g. the electronic properties of composite electrodes and their individual components. It can be concluded that the positions of SEI components preferably should be calibrated based on the binding energies for surface compounds such as hydrocarbons or other prominent peaks from compounds usually found in the SEI. However, for referencing bulk materials, a surface based calibration approach is clearly not ideal, especially not in the presence of SEI layers as potential gradients at the buried interface introduce additional binding energy shifts. A different, bulk specific internal reference is to be preferred unless the composite has a good electronic conductivity, in which case an internal calibration may not be required at all.



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It is tempting to try to establish general guidelines on how to calibrate battery related PES measurements, but as the energy referencing procedure is highly sample dependent every system most likely requires its own calibration and reference measurements.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge Au 4f7/2 spectra and peak parameter table for samples in Figure 2. SEM micrograph and EDS elemental mapping of the mixed materials electrode. Voltage vs. capacity curve for the lithiation reaction of nano-Si electrodes.

AUTHOR INFORMATION Corresponding Author *[email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ACKNOWLEDGMENT The work presented in this paper is undertaken within a joint development project at the HVV (www.highvoltagevalley.se) consortium and financed by the Swedish Governmental Agency for Innovation Systems (Vinnova) and with the EU FP7 projects Eurolion and Hi-C. Further, we thank HZB for the allocation of synchrotron radiation beam time. The research leading to


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these results has received funding from the European Community's Seventh Framework 19 Programme (FP7/2007-2013) under grant agreement no. 312284. The authors thank for generous sponsoring of this work from the Swedish Research Council (2016-03545) and Swedish Energy Agency (40495-1). Many thanks to Reza Younesi for valuable discussions and to Andreas Blidberg for coming up with the mixed materials electrode idea.

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TOC Graphic


Interpreting PES Peak Positions for Cycled Li-ion Battery Electrodes

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