On the Lithiation Mechanism of Amorphous Silicon ... - ACS Publications

Subscriber access provided by Nottingham Trent University

C: Energy Conversion and Storage; Energy and Charge Transport

On the Lithiation Mechanism of Amorphous Silicon Electrodes in Li-Ion Batteries Daniel Uxa, Bujar Jerliu, Erwin Hüger, Lars Dörrer, Michael Horisberger, Jochen Stahn, and Harald Schmidt J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b06011 • Publication Date (Web): 16 Aug 2019 Downloaded from pubs.acs.org on August 16, 2019

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 39 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

The Journal of Physical Chemistry

On the Lithiation Mechanism of Amorphous Silicon Electrodes in Li-Ion Batteries

Daniel Uxa1, Bujar Jerliu1, Erwin Hüger1, Lars Dörrer1, Michael Horisberger2, Jochen Stahn3, Harald Schmidt1,4,*, 1Technische

Universität Clausthal, Institut für Metallurgie, AG Mikrokinetik, Clausthal-Zellerfeld, Germany.

2Laboratory

for Scientific Developments and Novel Materials, Paul-Scherrer-Institut, Villigen PSI, Switzerland

3Laboratory

for Neutron Scattering and Imaging, Paul Scherrer Institut, Villigen PSI , Switzerland

4Clausthaler

Zentrum für Materialtechnik (CZM), Clausthal-Zellerfeld, Germany.

*Corresponding author; [email protected] Abstract Amorphous silicon is a high-capacity negative electrode material for use in advanced lithium-ion batteries. We investigated the mechanism of Li incorporation into and removal from this material during electrochemical lithiation and delithiation using a combination of in-operando Neutron Reflectometry and ex-situ Secondary Ion Mass Spectrometry. The results indicate that a heterogeneous lithiation mechanism is present for the first cycle and also for subsequent cycles during lithiation and delithiation, where a highly lithiated phase is penetrating the silicon electrode. During the first lithiation half-cycle a two-step process is acting, which is not present for delithiation and higher cycles. There, in a first step a Li-poor phase penetrates the silicon electrode leading to about 10 % of maximum capacity. Afterwards, during a second step a Li-rich phase moves into the electrode leading to complete lithiation in a slower process. The different phases are separated by a relatively sharp interface of several nanometers only. The Li-poor phase 1 ACS Paragon Plus Environment

The Journal of Physical Chemistry 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

Page 2 of 39

extended over the whole electrode is still present after delithiation in form of irreversibly trapped Li.

Introduction Rechargeable lithium-ion batteries (LIBs)

1–7

are superior to other types of batteries due to their

high energy density, high cell voltage and low self-discharge. In science and technology this type of battery is widely developed 8 for use in portable electronic devices (smart phones, laptops), in the automotive transportation sector (electric vehicles), as Megawatt power stations and also as nano-sized energy sources (smart cards, sensors, on chip applications). For future applications, battery weight, cycle life, energy density, power density and also costs have to be improved. In this context, an important point is the development of new negative electrode materials (termed anodes) with a higher specific capacity than traditional materials. Amorphous silicon is an anode material with a very high theoretical capacity between 3580 and 4200 mAh/g currently under discussion for applications 9,10. During electrochemical lithiation, lithium and silicon react and an amorphous LixSi alloy is formed. This process is associated with drastic structural changes

11,

enormous volume expansion (up to 400 %) 12,13 and mechanical stress 14. Consequently, a decay of the specific capacity during subsequent cycling takes place that is associated with mechanical fracture and irreversible side reactions that are triggered by the volume changes 9,12 In this context the exact mechanism of lithiation and delithiation is of considerable interest. Especially the discrimination between a homogeneous and a heterogeneous lithiation mechanism is of importance. For the first of these two cases the electrode is homogeneously filled with lithium, while (at least approximately) a constant Li concentration across the whole electrode is maintained during lithiation. A slight gradient in Li concentration may occur at the surface in order to allow diffusional flux of Li into the electrode. In the second case a highly lithiated phase with a sharp

2 ACS Paragon Plus Environment

Page 3 of 39 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


phase boundary penetrates successively the silicon electrode and the situation of two (or more) coexisting phases occurs. The large change of lithium concentration across the phase boundary of small width (some nanometers) may lead to stress development and possibly to electrode damaging and performance loss. In the present paper, in-operando and real time Neutron Reflectometry (NR) measurements on amorphous silicon electrodes during charging and discharging cycles are presented. These measurements are combined with ex-situ Secondary Ion Mass Spectrometry (SIMS) measurements with the aim to identify the lithiation mechanism.

Experimental details The electrochemical experiments were performed with a self-constructed three-electrode electrochemical cell sealed against air as described in ref.

15.

It can be used for both: the in-

operando experiments at neutron facilities and the ex-situ experiments with SIMS. Thin film electrodes prepared by sputtering were applied for the experiments. The working electrode for NR consists of a 1 cm thick quartz support coated with an about 400 nm thick copper layer as a back contact and current collector. The amorphous silicon film as active material is coated on top of the copper. The counter and reference electrode are made of metallic lithium (1.5 mm foil, 99.9 %, Alfa Aesar). A microporous polyethylene separator (Brückner Maschinenbau, Germany) was introduced between the two electrodes. Propylene carbonate (Sigma Aldrich, anhydrous, 99.7 %) or the deuterated form 1,2-Propylene-d6 carbonate (Sigma Aldrich, 98 % D) with 1M LiClO4 (Sigma Aldrich, battery grade) was used as an electrolyte. The electrochemical cell was assembled within an argon filled glove-box (water content < 0.1 ppm, oxygen content < 0.1 ppm). The active layer of amorphous silicon in the working electrode for the NR experiments was deposited by dc magnetron sputtering at PSI Villigen directly on the copper coated quartz substrate using argon sputter gas (base pressure: 2.7 x 10-4 Pa, operating pressure: 0.36 Pa, sputtering power: 3 ACS Paragon Plus Environment


150 W) at room temperature. The diameter of the circular electrode was about 4 cm. Alternatively, the silicon electrodes with a lower diameter of about 2 cm for the SIMS experiments were produced by ion-beam sputtering (IBC 681, Gatan) directly on Cu supports. The base pressure of the vacuum chamber was 5 x 10-7 mbar. Sputtering was also done with Ar+ ions at a working pressure of 5 x 10-5 mbar. The ion beam acceleration voltage was 5 kV and the ion beam current was 200 µA. The silicon layer of both types of electrodes was amorphous after deposition. An investigation by Grazing incidence X-ray diffraction (Bruker D5000, Co Ka1, 40 kV) showed only characteristic sharp Bragg peaks of copper. This confirms the amorphous nature of the electrode. Additional crystallographic phases were not detected. Galvanostatic cycling experiments were carried out with a computer controlled potentiostat (BioLogic, model SP-50) using electrical current densities between 1.6 A/cm2 and 30 A/cm2 resulting in lithiation potentials between 2.5 and 0.0 V vs. Li. The electrochemical cycling curves show no significant differences for the two types of electrodes. In-situ NR was performed at the Time-of-Flight reflectometer Amor (Paul-Scherrer-Institute Villigen) and at the / reflectometer V6 (Helmholtz Center Berlin). During NR measurements the quartz substrate is the incoming medium and a neutron beam hits the surface at a small angle, . The neutrons are partly reflected at and partly transmitted through the SiO2/Cu/Si/electrolyte interfaces, exit the quartz on the opposite side. Afterwards, they are detected. The intensity of the reflected beam as a function of scattering vector qz = 4π sin θ / λ is recorded ( wavelength). The reflectivity is the ratio between the reflected and the incoming beam and can be seen as an interference pattern originating from reflections of all interfaces present16. The contrast of the neutron index of refraction determines the intensity of the reflected beam perpendicular to the reflecting interface. Data analysis was performed using Parratt’s recursion formula within the Motofit software package 17 by fitting the measured NR patterns.


Page 4 of 39

Page 5 of 39 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


The measurements at V6 were carried with a monochromatic collimated neutron beam (40 mm horizontal slits, 0.5 mm vertical slits) at a wavelength of 0.466 nm. 3He pencil detectors, offset from each other by 0.44° in 2, are used for recording the scattered intensity in the specular and background channel, respectively15. The reflectivity data were corrected for footprint and background. The measurements at Amor were realized with the Selene setup

18,19

. For standard

NR methods either the angle or the wavelength is fixed and the other parameter is varied. With the Selene setup both quantities can be varied at the same time. The wavelength is determined by timeof-flight, and the scattering angle with a position sensitive detector16. The raw data (intensity, time of flight and detector position) are transformed into intensity as a function of qz by a data reduction program giving the reflectivity curve. The SIMS measurements were carried out with a Cameca imf 3f/4f machine using an O2+ primary ion beam (5 keV) in depth profiling mode. The sputtered area was about 250 m x 250 m. Depth calibration was done by measuring the crater depth with a stylus profilometer (Tencor, Alphastep).

Results and Discussion Experiments with Secondary Ion Mass Spectrometry (SIMS) In Fig. 1(a) a characteristic galvanostatic charging curve and in Fig. 1(b) a discharging curve of the first lithiation cycle as well as partly the charging curve of the second cycle measured on a characteristic silicon electrode after assembling the cell are given. They show the change of the cell potential measured between working electrode and reference electrode for a current density of about 30 A/cm2. The initial thickness of the silicon electrodes was about 450 nm resulting in a lithiation rate of  C/9. The lithiation time of the first half cycle is about 9 h (Fig. 1(a)). This time period varies for different samples prepared within 10 % due to variations in the initial thickness (see below). The potential decreases as a function of charging time from about 2.7 V (initial open circuit voltage) down to 0.0 V during lithiation and the reverse process is observed during 5 ACS Paragon Plus Environment


delithiation. As shown in Fig. 1, charging/discharging is interrupted at certain stages of the lithiation/delithiation process for SIMS analyses. The cell is opened in the glove-box and the electrodes are rinsed and cleaned with salt free propylene carbonate and afterwards with isopropanol. Afterwards the electrode is transferred into the SIMS machine without any contact to air and analysed afterwards. Post mortem SIMS measurements are done after certain time steps on different samples, which were prepared under identical sputter deposition conditions.

3.0

3.0

(a)

(b)

2.5

2.5

2.0

2.0

Potential (V)

Potential (V)

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

Page 6 of 39

1.5 1.0 0.5 ooo o o

o

0.0 0

1

2

3

4

o

5

7

8

1.5 1.0 0.5

o o

0.0

o

6

o

0

9

1

2

3

4

5

6

7

8

9 10 11

Time (h)

Time (h)

Figure 1 (a) Lithiation curve of a silicon thin film electrode (450 nm) galvanostatically lihiated at a current of 30 A/cm2 (about C/9). The potential (vs. Li/Li+) is plotted as a function lithiation time. (b) Corresponding delithiation curve (black line) and second lithiation curve (blue dotted line). The red circles indicate times where lithiation/delithiation was interrupted and SIMS measurements were done (for different samples of the same type).

A SIMS depth profile of a silicon electrode in the as-deposited state without any lithiation is given in Fig. 2. The secondary ion intensities of 28Si and 65Cu masses are shown. No measurable signal of lithium is detected. We observe a constant signal of 28Si over the whole film thickness, which indicates a homogeneous silicon film with a thickness of 450 nm as obtained from Fig. 2. The Si/Cu interface is indicated by a strong increase of the

65Cu


signal. The initial thickness of the

Page 7 of 39

different silicon electrodes deposited in this study with the same method and the same deposition parameters varies between 430 and 470 nm.

Secondary ion 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


10

4

28

Si Cu

65

10

3

10

2

10

1

0

100

200

300

400

Sputter depth (nm)

Figure 2 Depth profiles of the isotopes 28Si and 65Cu of a silicon electrode (450 nm) right after sputter deposition.

In Fig. 3 the results of SIMS experiments after lithiation of 3 min, 15 min, 30 min, 1h, 1h 40 min, 4 h, and 5 h 30 min (after starting) as well as for full lithiation are given. Full lithiation was achieved by lithiating the sample without potential limitation until a constant potential close to zero is established (here: -0.02 V). Shown are the intensities of the 7Li and 28Si masses. The situation after lithiation of 3 min only is given in Fig. 3(a). During this short time period where the cell potential decreases significantly to 0.27 V (see Fig. 1) the surface is enriched with lithium on a length scale of 10 - 20 nm. As can be seen in Fig. 3(b) and (c) this Li enriched region is penetrating the electrode continuously increasing its thickness and a plateau with a constant Li signal is established. This behaviour can be interpreted with the existence of an amorphous LixSi phase that penetrates the electrode in form of a moving phase boundary. At this stage of the lithiation process we have the 7 ACS Paragon Plus Environment


situation that two phases are present, the pure silicon and the penetrating LixSi phase. During lithiation the thickness of pure silicon decreases while the thickness of the LixSi phase increases. After more than 1 h of lithiation (extrapolated from Fig. 3(d)) this phase boundary is reaching the Cu current collector. A comparison of the measured Li to Si ratio of the SIMS secondary signals to experiments with X-ray photo electron spectroscopy (XPS) done on samples that were prepared in the same way 20 allows to assess the relative Li content of the penetrating phase roughly to about x  0.3 in LixSi. This value is confirmed by the results of the electrochemical measurements. The relative Li content can be calculated to x = I t M/(F m) (M = 28.09 g/mol: molar mass of silicon, m  3.5 x 10-4 g: actual mass of the thin silicon film, F = 96487 C/mol: Faraday constant, I = 100 A: current, t: lithiation time) by considering the consumed charge (Q = I t). From Fig. 3(d) after 1 h of lithiation, where the electrode is nearly completely filled with the Li-poor phase we get a value of x  0.3 in good agreement to the XPS measurement. Using Fig. 3(e) after 1h 40 min for calculation we get x  0.5 also still in acceptable agreement. In comparison to the maximum amount of Li that can be stored in silicon of x  4.2 this phase is termed Li-poor phase, as indicated in Fig. 3 as LPP. This part of the overall lithiation process corresponds mainly to the presence of a constant potential of 0.27 V seen in Fig. 1(a) for the first 90 min of lithiation. After this period it is expected that the sample (in good approximation) is completely and homogenously filled with a lithium concentration of about x  0.3 (despite of a Lirich surface layer). A significant modification of the electrode thickness is not observed. This first lithiation step is a relatively fast process leading to roughly 10 % of the maximum capacity. Generally, amorphous silicon is distinguished by the presence of dangling bonds and also small amounts of impurity atoms like oxygen in the percentage range. Consequently, we suggest to attribute this first fast lithiation step tentatively to a fast interstitial diffusion of Li into silicon and the trapping of Li at dangling bonds and impurities or also interfaces to pores. The plateau and sharp phase boundary indicates a reaction limited diffusion process with strong trapping as also 8 ACS Paragon Plus Environment

Page 8 of 39

Page 9 of 39

observed for hydrogen in silicon or silicon carbide for certain conditions 21. The question whether really an “amorphous phase” or amorphous silicon saturated with lithium is present remains to the future. In Fig. 3(d) after 1 h of lithiation the formation of a second phase at the surface is indicated, that continuously and slowly penetrates the electrode until complete lithiation and a potential close to 0.0 V is reached. Here, also a phase boundary is penetrating the electrode with increasing lithiation time (Fig. 3(e-h)). The higher amount of Li associated with this phase is reflected by a one order of magnitude higher Li to Si ratio. According to the results of the SIMS analyses the Li to Si ratio in the plateau is roughly constant during phase penetration. This indicates that the relative Li fraction x in LixSi is also constant. We assume here the maximum possible Li concentration of x  4.2 to be present in the Li-rich phase (LRP in Fig. 3). Phase penetration takes place by a continuous transformation of the Li-poor phase into the Li-rich phase. This can be realized by a process where the necessary Li is provided by the electrolyte and diffuses through the Li-rich phase to the interface between the Li-rich and Li-poor phase, where the reaction with silicon takes place. This results in a continuous increase of the thickness of the Li phase by consumption of the Lipoor phase and an increase of the overall electrode thickness. As visible in Fig. 3(d) the formation and initial movement of the Li-rich phase seems to be present at a stage the Li-poor phase did not reach the current collector. At later stages, see e.g. Fig. 3(f) the Li-poor phase is present across the

6

10

5

(a) LPP

3 min

SEI

10

4

10

3

10

2

7

Li Si

28

10

6

10

5

LPP

(b) 15 min

SEI

10

4

9

7

Li Si

28

3

10 ACS Paragon Plus Environment 10

2


10


whole electrode.


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


10

6

10

5

10

4

10

3

10

2

LPP

(c) 30 min

SEI

7

Li Si

28

10

6

LPP

(d) 1h

5

LRP

10

4 7

Li Si

28

3

10 SEI 10

2

10

1

0

100

200

300



10

400

10

6

10

5

1 h 40 min (e) LRP

10

4 7

Li Si

28

10

3

SEI

10

2

10

1

0

200

Sputter depth (nm)

10

6

10

5

10

4

10

5 h 30 min (g) LRP

SEI

3

7

10

2

10

1

Li 28 Si

0

500

1000

400

600

800

1000

Sputter depth (nm)

1500



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

10

6

10

5

10

4

10

3

Page 10 of 39



10

6

10

5

10

4

10

3

10

2

10

1

4 h (f) LRP

SEI

7

Li Si

28

0

500

1000

1500

Sputter depth (nm)

full lithiation (h)

LRP

SEI

7

Li Si

28

10

2

10

1

0

500

Sputter depth (nm)

1000

1500

2000

2500

Sputter depth (nm)

Figure 3 Depth profiles of the isotopes 28Si and 7Li of a silicon electrode after different stages of lithiation at 30 A/cm2. Given are results after a lithiation time of (a) 3 min, (b) 15 min, (c) 30 min, (d) 1 h, (e) 1 h 40 min, (f) 4 h, (g) 5 h 30 min, (h) complete lithiation (not potential limited). See also red circles in Fig. 1(a). Note the logarithmic scaling of the intensity. LLP: Li-poor phase, LRP: Li-rich phase, SEI: solid electrolyte interphase. From (e) to (h) the LPP extents over the whole electrode and is no longer indicated. The fact that the penetrating Li-rich phase is not reflected in a further increase of the Li signal but in a decrease of the Si signal can be understood in the following way: The SIMS secondary ion intensity is proportional to the concentration of the corresponding mass/element multiplied by the 10 ACS Paragon Plus Environment

Page 11 of 39 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


ionization probability. In this context, a direct quantification is not possible without detailed calibration standards because of the matrix effect 22. The ionization probability of a certain element depends critically on the concentration of the surrounding elements and bonding properties in a non-linear way. Consequently, the secondary ion intensity of the two elements is different in the Li-poor and Li-rich phases. The relatively sharp interface of the penetrating phases does not only reflect modifications in concentration but also in ionization probabilities. Considering relative concentration only, we expect for the Li-rich phase a higher Li signal (as for the Li-poor phase) what is not observed and a lower Si signal what is observed. The relative Si concentration in the Li-rich phase is lower than in the Li-poor phase. The unexpected nearly unchanged Li signal in the Li-rich phase compared to the Li-poor phase is due to the fact that an increase in Li concentration is masked by a strong decrease in the ionization provability. Note that the interface roughness of both penetrating phases can be assessed by SIMS to be about 10-15 nm for the initial stages and higher for later stages. Lithiation is also associated with a high expansion of the electrode thickness ranging from 450 nm in the initial state to more than 2500 nm in the fully lithiated state (increase in distance between the film surface and the Si/Cu interface (Figs. 2 and 3(h))). This indicates a significant volume expansion of the film during lithiation as reported in the literature

12,13,

which is discussed in detail

below together with the NR experiments. A further peculiarity of the present measurements is that in the maximum lithiated state (Fig. 3(h)) is still a region with a reduced Li fraction at the interface between electrode and current collector also while the main part of the electrode is highly lithiated. This is difficult to interpret in light of the fact that close to interfaces the ionization probability of all species and also the sputter velocity may change drastically in SIMS analysis. The Si-rich (Li-poor) region located around 2000 nm might be in reality much less extended than displayed. However, our results indicate the presence of a Li-poor region at the interface to the current collector which cannot be fully lithiated. A 11 ACS Paragon Plus Environment


Page 12 of 39

possible explanation might be the presence of large stresses close to the interface that slow down lithiation (see 9,23,24), even to a point where it ceases as discussed in 23 . A similar behaviour was observed for the lithiation of silicon nano-particles/-wires where also stress was made responsible for such a self-limiting effect25–27. Alternatively, a modification of the lithium electrochemical potential (without stress) close to the Si/Cu interface may result in a similar effect. An alternative interpretation of the SIMS results might be that we have no real penetration of the Li-rich phase accompanied by a simultaneous consumption of the Li-poor phase. It is also possible that the thickness of the Li-poor phase remains roughly constant (or decreases only slightly) during lithiation while the Li-rich phase expands. This can only be realized if simultaneously the relative Li fraction x in LixSi of the Li-rich phase increases from x  0.3 to 4.2 continuously during the lithiation process. This will result in a pseudo-movement of the Li-rich/Li-poor interface relative to the surface. This scenario is supported by the observation (Fig. 3) that the distance between the interface and the Cu current collector is not verifiable reduced in the whole process. It varies between about 400 to 800 nm with no clear trend. However, the fact that the Li to Si ratio in the plateau of the Li-rich phase remains constant contradicts this scenario. Possibly, a combination of both scenarios is present. For both scenarios, the data show that we have a strong variation of the Li concentration on the length scale of only some tens of nanometre inside the electrode that moves relative to the surface. In conclusion, the results of the SIMS experiments of the first lithiation cycle indicate the existence of a two-step lithiation mechanism via moving phase boundaries. First a Li-poor phase penetrates the silicon electrode relatively fast leading to about 10 % of full capacity. Afterwards a Li-rich phase moves into the electrode leading to complete lithiation in a slower process. The electrode is not completely filled with lithium while a Li-poor region at the interface to the current collector remains. During lithiation, a Li-poor phase and a Li-rich phase separated by a relatively sharp interface are present. The existence of a moving phase boundary during lithiation of amorphous 12 ACS Paragon Plus Environment

Page 13 of 39 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


silicon was also indicated by in-situ TEM studies, however, at undefined high current densities 28,29. Further details will be discussed below. The occurrence of a lithiation mechanism with a sharp

phase boundary in amorphous materials is quite surprising. “Amorphous phases“ are not clearly defined and a priori not expected. The reasons may lie in the characteristics of the reaction kinetics. It can be assumed that the reaction rate constant for the solid state lithiation reaction is low and characterizes the reaction as rather sluggish as it is the case with more complicated reactions involving significant molecular rearrangement upon electron transfer or multistep mechanisms30. Given the fact that lithium diffusivities 31 and also permeabilities (diffusivity x solubility)32 in the amorphous Li-rich phase are several orders of magnitude higher compared to pure silicon and also to the Li-poor phase it becomes apparent that the found sharp phase boundary is the result of a reaction controlled lithiation mechanism of amorphous silicon. A high number of Li atoms can be transported fast in the Li-rich phase, while Li migration is stopped at the interface between the two phases coexisting during lithiation. At the interface the chemical solid state reaction takes place converting the phases into each other. The slow permeation of Li in silicon and in a Li-poor phase prevents a homogeneous lithiation of the whole electrode32. In Fig. 4 additional SIMS depth profiles are given which were recorded (a) after complete lithiation and subsequent delithiation of 5 h, (b) after the first cycle and (c) after the first cycle and additional lithiation of 4 h (see also Fig. 1(b)). As obvious from Fig. 4(a) during delithiation there is also a strongly heterogeneous mechanism acting. The Li-rich phase introduced into the electrode during lithiation is now removed during delithiation in the same way by a moving phase boundary. This result is in contrast to statements based on TEM measurements given in literature

28

which

postulate a homogeneous delithiation mechanism. After complete delithiation we see again a homogeneous Li distribution within the whole electrode (Fig. 4(b)). The difference to the virgin state in Fig. 3(a) is that Li in small amounts is still present, homogeneously distributed in the electrode. This is the origin of irreversible capacity losses widely discussed in literature 9,32–34 and 13 ACS Paragon Plus Environment


recently found to be also due to Li irreversible stored in the whole electrode and not only in a surface layer 16. The concentration of the irreversibly stored Li can also be assessed roughly to x  0.3. This is seen in the SIMS depth profiles from a comparison of the ratio of the secondary ion signals of Li and Si. In Fig. 3(e) or (f) at high depths of 1000 nm where the Li-poor phase is well established over the whole electrode thickness this ratio is about ILi/ISi  5, which is the same as given in Fig. 4 (b) for the delithiated state. This estimation indicates that the relative Li concentration in the Li-poor phase occurring during first lithiation is roughly the same as in the delithiated state. This indicates that the lithiation process of the Li-rich phase is mainly reversible while that of the Li-poor phase is irreversible. Delithiation is done in a one-step process and the Li introduced by the Li-poor phase during lithiation remains irreversibly trapped in the electrode. Fig. 4(c) shows depth profiles after the first cycle and additional lithiation of 4 h during the second cycle. The result is very similar to the curve in Fig. 4(a) and points again to a lithiation mechanism of a Li-rich phase via a moving phase boundary. Possibly, the interface between both lithiated phases smears out further as indicated by a comparison of Fig. 4(a) and Fig. 4(c) probably due to increasing interface roughness. The heterogeneous mechanism is also present for higher cycles than the first. This is also contrary to literature, where a homogeneous mechanism is postulated for higher cycles28. Finally, SIMS depth profiles were recorded after lithiation of 31 h of the same electrode as discussed above, however, for a significantly lower current density of about 3 A/cm2, as shown in Fig. 4(d). Qualitatively, nearly the same depth profile is present as visible in Fig. 3(f) or 4(a). Only the Li to Si ratio in the Li-poor phase is lower, indicating a lower Li content to be formed during the first lithiation step. The results prove that also at very low current densities a heterogeneous mechanism is acting. 10

6

10

5

10

4

(a) LRP

7

10

3

10

2

10

1

Li Si

SEI

28



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

Page 14 of 39

10

6

10

5

10

4

10

3

10

2

200

400

600

800 1000 1200

Sputter depth (nm)

SEI

7

14

0

(b)

10

Li Si

28

1

0

100

200

ACS Paragon Plus Environment

300

400

Sputter depth (nm)

500

10

6

10

5

10

4

(c)


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



Page 15 of 39

LRP

7

Li 28 Si

SEI

10

3

10

2

10

1

0

200

400

600

800

1000

10

6

10

5

10

4

10

3

10

2

10

1

(d) LRP

7

Li Si

SEI

Sputter depth (nm)

0

200

28

400

600

800 1000 1200

Sputter depth (nm)

Figure 4 Depth profiles of the isotopes

28Si

and 7Li of a silicon electrode during cycling at 30

A/cm2. Given are results after (a) 5 h of delithiation, (b) complete delithiation (after first cycle), and (c) after the first cycle and additional 4 h of lithiation (second cycle). See also red circles in Fig. 1(b). In (d) profiles for lithiation (first cycle) at about 3 A/cm2 for 31 h are given. Note the logarithmic scaling of the intensity. LLP: Li-poor phase, LRP: Li-rich phase, SEI: solid electrolyte interphase.

Also clearly visible in Fig. 3 and 4 is that at the surface there exists an additional thin surface layer with a high Li to Si intensity ratio which is distinguished by the fact that the Si signal reaches nearly zero. We interpret this layer as the solid electrolyte interphase (SEI), which is a passivation layer between the electrolyte and electrode and is formed by a dissociation of the electrolyte9,35. The initial thickness in the virgin state is about 10 nm or less. During lithiation this layer is expanding to a value of 80-90 nm in the fully lithiated state and reduces again to 10 nm after the first cycle. Such a behavior is also found in other studies36–39, however, with strongly varying thicknesses of the SEI. The absolute thickness of the SEI is a complex problem because it may depend on various parameters like, electrolyte composition, structural state, layer thickness of silicon, cycle number, and impurities etc.



Page 16 of 39

In preliminary studies, SIMS measurements were also carried out for other current densities of 7.8 A/cm2 and 100 A/cm2 for a single given lithiation time 36. There, also a penetration of a Li phase boundary into silicon is observed, indicating the presence of a heterogeneous mechanism.

Experiments with Neutron Reflectometry (NR) In contrast to the ex-situ SIMS experiments, the NR experiments were done in-operando, meaning reflectivity patterns are continuously recorded in real time during galvanostatic cycling without disassembling the cell or stopping the current. Measurements at low current densities were done at HZB, Berlin using the reflectometer V6 while measurements at high current densities were recorded at PSI, Villigen using AMOR/Selene because faster data acquisition rates are possible there. Galvanostatic charge-discharge curves similar to that given in Fig. 1 were recorded. Different current densities were investigated. Figs. 5-7 give results for an electrode with an initial thickness of 65 nm and a current density of 1.6 A/cm2, resulting in a lithiation rate of  C/50 done at V6. The recording of a reflectivity pattern lasted about 58 min. Note that for NR measurements an electrode thickness of about 500 nm like for SIMS is not preferable here, because the weakly expressed interference effects do not lead to meaningful results.

10

-1

10

-2

10

-3

10

-4

10

(a)

Virgin

0.00

0.02

0.04

0.06

0

10

-1

10

-2

10

-3

10

-4

Lithiation 5.80 h

0.00

0.08

0.02

0.04

0.06

-1

-1

qz (Å )

qz (Å )

16 -8 ACS 10 Paragon Plus Environment

-8

-4

Å )

-4

Å )

10

10

(b)

Reflectivity R

0

Reflectivity R

Reflectivity R

10

0.08

0

10

-1

10

-2

10

-3

10

-4

(c)

Lithiation 19.33 h

0.00

0.02

0.04 -1

qz (Å )

0.06

0.08

-8

10

-9

4

-4

10

10

-10

0.00

0.02

0.04

0.06

-1

qz (Å )

Figure 5 Neutron reflectometry (R) patterns (open circles) measured on the 655 Å electrode (current density 1.6 A/cm2) together with fitting results using Motofit (lines). Also shown is a R qz4 representation for a better visualization of the fringes (bottom). Given are results (a) for the virgin state, (b) after 5.80 h of lithiation and (c) after 19.33 h of lithiation. The Motofit fits are based on the homogenous model (red) and the heterogeneous model (blue) as described in the text. Fig. 5 shows different examples of reflectivity curves recorded during first lithiation. In addition to the edge of total reflection at 0.01 Å-1, characteristic slight fringes can be seen resulting from the low difference in neutron scattering length density (SLD) between silicon (2.00 x 10-6 Å-2) and the electrolyte (1.75 x 10-6 Å-2). If during lithiation Li is incorporated into the electrode the reflectivity R is modified. This reflects mainly changes in thickness of the electrode but also in SLD. The modification of the patterns is better visualized in the bottom of Fig. 5, where the quantity R qz4 is shown. In Fig. 6 selected reflectivity patterns obtained during first delithiation are given. All reflectivity patterns recorded during the first cycle are given in Fig. 7 in form of a contour plot, where scattering vector and lithiation/delithiation time is plotted against reflectivity. It becomes clear already from optical inspection that the patterns are continuously modified during cycling, due to the modification of volume/electrode thickness and SLD.

0

10

-1

10

-2

10

-3

10

-4

10

(a)

Delithiation 19.33 h

17

Reflectivity R

10

Reflectivity R

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


R qz (Å )

Page 17 of 39

0

10

-1

10

-2

10

-3

Delithiation 33.83 h

ACS Paragon Plus Environment -4 10

(b)

0.08

-9

10

-8

10

-9

Page 18 of 39

4

-4

10

R qz (Å )

-8

-4

10

4

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

R qz (Å )


10

-10

0.00

0.02

0.04

0.06

10

0.08

-10

0.00

-1

0.02

0.04

0.06

0.08

-1

qz (Å )

qz (Å )

Figure 6 Neutron reflectometry (R) patterns (open circles) measured on the 655 Å electrode (current density 1.6 A/cm2) together with fitting results using Motofit (lines). Also shown is a R qz4 representation for a better visualization of the fringes (bottom). Given are results (a) after 19.33 h of delithiation and (b) after 33.83 h of delithiation. The Motofit fits are based on the heterogeneous model as described in the text.


Page 19 of 39 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


Figure 7 Contour plot of the reflectivity (R qz4 representation) for the first cycle (red: high reflectivity; green: medium reflectivity; blue: low reflectivity) as a function of scatting vector qz and measurement time (each measurement number corresponds to a time period of 0.97 h). Measurements are done at the 655 Å electrode (current density 1.6 A/cm2). Lithiation takes place from number 1 to 53 and delithiation from 54 to 89, respectively.

Each NR pattern experimentally recorded was fitted by the program Motofit using a layer-model as described in detail in Ref. 15. Corresponding, results are also shown in Fig. 5 and 6 as lines. For the virgin state directly after assembling the cell without any current applied, a sequence of layers corresponding to the Cu current collector and the Si active material sandwiched between the SiO2 support and the electrolyte is used. The pattern in Fig. 5(a) can be described well using a SLD of the copper layer of 6.45 x 10-6 Å-2, of the quartz support of 4.15 x 10-6 Å-2, and of the electrolyte of 1.75 x 10-6 Å-2 as well as a thickness of the copper layer of 4000 Å. These quantities were kept constant for all the following fitting procedures and analyses. For the SLD of silicon a value of (2.00 ± 0.14) x 10-6 Å-2 was obtained. Errors correspond to a 10 % increase of 2 of the best fit. The SLD of silicon corresponds to a mass density of 2.23 g /cm3 which is in accordance to literature (2.19 - 2.29 g/cm3) 40. The initial layer thickness of silicon is (655 ± 20) Å. The interface roughness between the single layers was fixed to 3 Å, only that between Si and Cu was a fitting parameter, resulting in a value of (11 ± 2) Å. In literature 37,38,41, for an adequate description of experimentally derived NR data by simulation tools, often a thin surface layer with a differing SLD is included in order to consider an initial oxide 19 ACS Paragon Plus Environment


Page 20 of 39

layer or the SEI. For the present case, a significant improvement of the fits cannot be achieved by doing so. This can be traced back to the fact that our experimental arrangement (SLD of electrolyte, thickness of Cu layer) is not very sensitive to this feature in the qz range investigated. For example, the SLD difference between the interface layer and the electrolyte is too small to result in a sufficient influence of the reflectivity. Consequently, we focused on the systematic SLD modification of the silicon layer and neglected such a surface layer in approximation. In a first attempt to describe the experimental data a procedure was carried out where the reflectivity patterns modified during lithiation/delithiation were fitted with the SLD and the layer thickness as free parameters. The SLD was assumed to be constant across the whole electrode layer. This assumption is in contradiction to the heterogeneous lithiation mechanism indicated by the SIMS experiments discussed above. Consequently, the SLD obtained from the fits is a heavily averaged SLD. The aim of this fitting approach is (a) to achieve an easy determination of the modification of the electrode layer thickness during cycling and (b) to see whether a homogenous lithiation model can explain our NR data. Typical examples of reflectivity curves recorded in-operando continuously during cycling at different potentials and the corresponding fits according to the homogeneous model are given in Fig. 5(b,c) (red curve). The reflectivity patterns can be well described with this homogeneous model assumption in contrast to the SIMS results. The relative volume modification V/V0 = L/L0 (L0 = 655 Å) and the averaged SLD are given in Fig. 8 for two cycles. Concerning relative volume, in Fig 8(a) we observe a continuous volume expansion during lithiation up to 380 % and a reduction in volume during delithiation for both cycles. During first lithiation we can separate two regions: an initial region where only slight thickness modifications are found (t < 8 h) and a second region, where a continuous, approximately linear increase of the relative volume is visible (t > 8 h). A more detailed discussion will be given below. Note that the relative volume after the first complete cycle is higher than the initial relative volume by a factor of 1.15. This is associated with the 20 ACS Paragon Plus Environment

Page 21 of 39

irreversible Li loss already mentioned above. A linear modification of the volume during lithiation/delithiation was already reported by our group 12 and others 13,42. During Li incorporation the silicon layer expands in the direction perpendicular to the surface 43. In contrast, the expansion in the direction parallel to the surface is avoided by adhesion of silicon to the substrate. Note also that with increasing volume the number of fringes becomes more and they are less pronounced (Fig. 5). For a relative volume V/V0 > 3 the situation came aware that the thickness of the lithiated silicon layer could no longer be unambiguously extracted from the reflectivity pattern. The number of fringes were too high and the intensity too low and the fitting procedure results in several different thicknesses with similar 2 values. Consequently, we proceeded in the following way by a stepwise procedure: First, we calculated the thickness Ly that would be expected from a linear extrapolation of the data already measured. Then, the fit of the new pattern was restricted to the range Ly ± 200 Å and the resulting fit parameter was used as the final value. The procedure was repeated subsequently for higher thicknesses. The results are marked in Fig. 8 by open circles. In Fig. 8(b) the corresponding behaviour of the SLD is displayed. There is a decrease of the average SLD during lithiation and a reverse process during delithiation. This indicates that considerable amounts of lithium are incorporated in and released from the electrode during cycling. Due to the negative coherent scattering length of Li, an incorporation of Li into the silicon layer results in a decrease of SLD. Because the SLD displayed in Fig. 8(b) is expected to be a heavily averaged quantity it is not analysed in further detail.

2.7

(a)

4

(b)

2.4 2.1 1.8

3 SLD

1.5

V/Vo

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


2

1.2 0.9

21

0.6

ACS Paragon Plus Environment 0.3

1

0.0

0

20

40

60

80

100 120 140 160

0

20

40

60

80

100

120

140

160


Page 22 of 39

Figure 8 (a) Relative volume V/V0 and (b) SLD modification of the 655 Å electrode (current density 1.6 A/cm2) as a function of time for the first two cycles. In (a) results are shown as obtained by fitting the homogeneous (black dots) and the heterogeneous (red cross) lithiation model. The meaning of the black circles is explained in the text.

During the next step of the fitting approach, the assumption of a homogeneous SLD distribution in the electrode layer is diminished. Based on the results obtained by SIMS a heterogeneous mechanism is more reliable. For fitting we use now a layer model based on moving phase boundaries as sketched in Fig. 9. There are two regions: the first region termed A where the Lipoor phase with a constant SLDLi-p = 1.2 x 10-6 Å-2 (x  0.5) is penetrating pure Si with SLDSi = 2.0 x 10-6 Å -2 and a second region B where a Li-rich phase with a constant SLDLi-r = -0.5 x 10-6 Å -2 (x  4.2) is penetrating the Li-poor phase with SLDLi-p = 1.2 x 10-6 Å-2. The SLD values are fixed during fitting. During the whole lithiation process 54 reflectivity patterns were recorded, where the second pattern after 0.97 h is identical to that of the virgin state (first pattern) and no measurable lithiation takes place. After the lithiation process in region A is finished (11th pattern) a state is assumed where the whole electrode is nearly completely lithiated with LixSi (x  0.5) corresponding to SLDLi-p = 1.2 x 10-6 Å-2. Note that x  0.3 as estimated by SIMS/XPS is also possible but give slightly worse fits. In order to simulate the penetration of the Li-poor phase, region A is subdivided into certain steps, where each step corresponds to a reflectivity pattern recorded. For region A these are nine steps (ti ; i = 1…9) and during each step a certain volume of silicon with a thickness of lSii and SLDSi is transformed into the Li-poor phase of lLi-pi > lSii and 22 ACS Paragon Plus Environment

Page 23 of 39

SLDLi-p. This adds up into a continuous penetration of the Li-poor phase into silicon. The overall thickness of the electrode L is a function of lithiation time. As mentioned, it continuously increases during lithiation. The overall thickness of the Li-poor phase is termed LLi-p which also increases as function of lithiation time and overall thickness of silicon is termed LSi which decreases as function of lithiation time. We get L(t) = LLi-p (t) + LSi (t) corresponding to

Li = LLi-pi + LSii

for time step ti (i = 1…9) with LLi-pi =  lLi-pi and LSii = L0 -  lSii with lLi-pi = 86 Å. Consequently, LLi-p (t) is fixed during the process as well as the interface roughness between the Cu current collector and the SiO2 support. A critical parameter is the roughness between the Li-poor phase and the silicon. For adequate fitting it has to be fixed to 80 Å. Such a high value is in agreement to the results of the SIMS experiments described above. Based on this model we have the following free fit parameters: the thickness of the residual silicon LSi (t), the interface roughness between the Cu current collector and Si and the surface roughness. The roughness of the Cu/Si interface varies during fitting statistically between 11 and 15 Å and is no further discussed. The same fitting procedure is done for region B in Fig. 9 assuming L(t) = LLi-r (t) + LLi-p (t), lLi-ri = 57 Å and Li4.2Si. Examples of fitting results can be found in Figs. 5 and 6 as blue curves. 3.0 2.5

A

SLDLi-p

SLDSi

B

SLDLi-r

SLDLi-p

2.0

Potential (V)

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


1.5 1.0

SLDSi = 2.0 x 10-6 Å-2 SLDLi-p = 1.2 x 10-6 Å-2 SLDLi-r = -0.5 x 10-6 Å-2

0.5 0.0

A 0

B 20

40

Time (h)



Figure 9 Schematic representation of the regions with different SLD used for fitting the NR data in the heterogeneous model. In Fig. 10 the SLD distribution as a function of lithiation time and spatial coordinate is given illustrating the penetration of the two phases in a multistep process. Interfaces are indicated. The penetration of the moving phase boundaries are clearly visible. The surface roughness is a parameter that increases from initially 3 Å continuously to 60 Å after the 12th measurement after 10.6 h of lithiation and remains constant afterwards statistically varying between 50 and 70 Å. It is clear that the model given in Fig. 10 is simplified compared to the results of SIMS analysis given above in order to reduce the number of free parameters during fitting. However, it demonstrates that a heterogeneous model is able to explain the measurement data. The data free region (white area) in Fig. 10 between measurement # 41 and # 67 could not be described unambiguously by the fits due to the fact that the layer thickness became too high to produce adequate fringes (see above and Fig. 8(a), circles).

-0.5 -0.5100 -0.4100 -0.3100 -0.2100 -0.1100 0 -0.5 -0.01000 0.09000 0.1900 0.2900 0.3900 0.4900 0.5900 0.6900 0.7900 0.8900 0.9900 1.090 1.190 1.290 1.390 1.490 1.590 1.690 1.790 1.890 1.990 2.090 -0.5 2.190 2.290 2.390 2.490 2.590 2.690 2.790 2.890 2.990 3.090 3.190 3.290 3.390 3.490 3.590 3.690 3.790 3.890 3.990 4.090 4.190 4.290 4.390 4.490 4.590 4.690 4.790 4.890 4.990 5.090 5.190 5.290 5.390 5.490 5.590 5.690 5.790 5.890 5.990 6.090 6.190 6.290 6.390 6.490 6.5

90.0 80.0 70.0

Measurement number

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

Page 24 of 39

60.0 50.0

Cu interface

40.0

6.5

Li-rich/Li-poor interface

30.0 Electrolyte interface

20.0 10.0 Li-poor/Si interface

0.0 5500

6000

6500

7000

7500

8000

8500

d (A)


9000

Page 25 of 39 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


Figure 10 Contour plot of the scattering length density (SLD) for the first cycle as a function of electrode thickness and measurement time (each measurement number corresponds to a time period of 0.97 h). Measurements are done at the 655 Å electrode (current density 1.6 A/cm2). The region in white between measurement 41 and 67 could not be described by the fits unambiguously due to the fact that the layer thickness became too high to produce adequate fringes. For details it is referred to the text.

The main result of these fittings based on the heterogeneous model is that the reflectivity patterns can also well described with identical 2 values assuming a heterogeneous lithiation mechanism (see Fig. 5, blue curves). Necessary prerequisite is that the interface width/roughness between the Li-poor phase and silicon (region A in Fig. 9) and also between the Li-rich phase and the Li-poor phase (region B) has to be relatively high. We have to fix it during fitting to 80 Å. Such a high value is in agreement to the results of the SIMS experiments described above. Further, the surface roughness has to increase from initially low values of 3 Å to higher values around 60 Å. Consequently, a blurred region between the two phases and a rough surface are necessary for adequate description. A description of the reflectivities with sharp interfaces is not possible and can be excluded. Note that a modified model as described in the SIMS section with an increasing Li concentration in the Li-rich phase and a nearly constant thickness of the Li-poor phase can also be used to describe the NR results (not shown). From the heterogenous model also the relative volume change during cycling can be extracted, which is also given in Fig. 8(a) as red crosses. The results are in agreement to the relative volume change obtained from the homogeneous model. This shows again (as already stated in ref.12) that the assumption of a concrete model is not necessary in order to extract reliable volume expansion and reduction data. Based on the results of the heterogeneous model the only weak volume expansion in Fig. 8(a) for low lithiation times is due to the first fast lithiation step describing the



Page 26 of 39

penetration of the Li-poor phase. The second penetration of the Li-rich phase corresponds to the linear part at higher times. An increase in volume during lithiation of amorphous silicon was also observed by AFM was quantified by Neutron Reflectometry Sectioning

43,

12

Synchrotron X-ray Imaging

(our group), AFM 46,

42,45

44

and

, Focussed Ion Beam Cross

Optical Diffraction Microscopy

47,

and X ‑ ray

Reflectivity 48. A maximum expansion up to about 400 % was found in agreement to our results and an (at least approximate linear) change in volume after initial effects 12,42,43,45,47,48 The reflectivity patterns recorded during delithiation of the first cycle and also during lithiation and delithiation of the second cycle can also be described by a heterogeneous mechanism. Based on the SIMS results only region B in Fig. 9 is present. Regarding the results of both fitting procedures based on a homogenous or alternatively on a heterogeneous model, it can be stated that from the NR measurements in the present configuration alone the extraction of the lithiation mechanism is not possible. Additional SIMS experiments are necessary which indicate qualitatively the dominating mechanism. Based on this qualitative result, adequate models for describing the reflectivity data can be constructed. During a second set of NR experiments done at Amor PSI, Villigen galvanostatic in-operando measurements were carried out using a higher current density of 4.7 A/cm2 and a lower electrode thickness of 125 Å resulting in significant higher lithiation rates of about C/3. The lower thickness was also chosen in order to obtain a reduced number of fringes in the accessible qz range in the lithiated state enduring a better fitting also at high volume expansion. In order to enhance the SLD contrast between Si and the electrolyte deuterated propylene carbonate was used (SLD = 5.8 x 10-6 Å-2). For these experiments taking place on a time scale of 180 minutes for the first cycle instead of 50 hours the time resolution has to be enhanced. Due to the Selene set-up installed at AMOR 18 a single reflectivity pattern can be measured within 2 min. Examples for reflectivity patterns and 26 ACS Paragon Plus Environment

Page 27 of 39

corresponding fits with Motofit (homogeneous and heterogeneous model) are given in Fig. 11. As obvious, only 1-2 fringes are present during the whole lithiation process. The results are in fundamental agreement to that described for a current density of 1.6 A/cm2 already discussed. The modification in volume and average SLD as well as the surface roughness as obtained using the homogenous fitting model is given in Fig. 12. The initial layer thickness of silicon is (125 ± 10) Å. The interface roughness between the single layers was fixed to 15 Å. Here, also the (partly) reversible expansion and contraction of the relative volume up to 380 % (Fig. 12(a)) and of the SLD (Fig. 12 (b)) are clearly visible. The roughening of the surface during lithiation can nicely be seen in Fig. 12(c). Fits of the reflectivity with a heterogeneous model as described above is also possible. In Fig. 11 such fits are also seen as a blue curve done with the parameters given in the figure caption.

10

0

10

(a)

0

(b)

10

-1

10

-2

10 10

76 min of lithiation

10

-1

10

-2

-3

10

-3

-4

10

-4

Reflectivity R

Reflectivity R

virgin

0.02

0.04

0.06

0.08

0.02

-1

10

(c) 158 min of lithiation

-1

10

-2

10

-3

10

-4

27 0.02

0.04 -1

qz (Å )

0.06

qz (Å )

0

10

0.04 -1

qz (Å )

Reflectivity R

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


0.06 Paragon 0.08Plus Environment ACS

0.08


Figure 11 Neutron reflectometry (R) patterns (open circles) measured at the 125 Å electrode (current density 4.7 A/cm2) together with fitting results using Motofit (lines). Given are results (a) in the virgin state, (b) after 76 min of lithiation and (c) after 158 min of lithiation. The Motofit fits are based on the homogenous model (red) and the heterogeneous model (blue) as described in the text. The parameters are for (b) LLi-r =126 Å, LLi-p =117 Å, surface roughness = 46 Å and for (c) LLi-r =372 Å, LLi-p = 18 Å, surface roughness = 58 Å). 2 values are identical for both types of fits.

4

4.0

3

3.0

-1

SLD (10 Å )

2.5 2.0 1.5

1 0

1.0 0.5

2

-6

Relative volume V/V0

(b)

(a)

3.5

end 1st cycle start

0

100

-1

200

300

400

0

100

70

(c)

60 50 40 30 20

200

300

Time (min)

Time (min)

rface roughness (Å)

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

Page 28 of 39


400

500

Page 29 of 39 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


Figure 12 (a) Relative volume V/V0, (b) SLD and (c) surface roughness of the 125 Å electrode (current density 4.7 A/cm2) as a function of time during cycling. After 426 min (overall time) the measurement has to be stopped due to an unexpected neutron beam interruption.

In literature, there are ex-situ and in-operando/in-situ studies available which also address the lithiation mechanism of amorphous silicon. There exist three in-situ studies carried out with transmission electron microscopy (TEM). First, the initial lithiation of a 20 nm thin amorphous silicon layer deposited by CVD on carbon nano-fibers was investigated 29 at high lithiation rates of 18 C using solid Li2O as an electrolyte. A two-step lithiation mechanism was found where in a first step a sharp moving phase boundary penetrates the silicon as derived from contrast variation (electron transparency) and volume expansion of the layer. In a second step fully lithiation was achieved, which however could not visualized because of low contrast. This is in agreement to our results concerning the existence of a two-step heterogeneous lithiation mechanism, however, the Li concentration of the first phase is expected to be higher. The second paper by McDowell et al. 28,

hydrogenated a-Si nano-spheres (about 600 nm) were lithiated/delithiated (Li2O eletrolyte).

Here, also high lithiation rates of 25 C were used. They also find by image contrast the penetration of a single sharp phase boundary during the first lithiation cycle (no two-step process). During delithiation and subsequent cycles such a moving phase boundary was not observed by contrast variations and a homogeneous mechanism was postulated. This later assumption is in disagreement to our findings. The reason is unclear, however, the very high current densities used in the TEM 29 ACS Paragon Plus Environment


Page 30 of 39

study may play a role or problems with contrast variation within TEM imaging. The third study 49 is based on STEM in a fully enclosed liquid electrolyte cell doing galvanostatic lithiation at 6.3 C. In contrast to the studies above they found no phase boundary but a more homogeneous filling of the electrode in form of solid solution formation. They attributed this result to the low contrast while silicon is immersed in the liquid electrolyte or to a large number of defects in this evaporated silicon electrode. The existence of a two-phase mechanism was also suggested by operando attenuated total reflection Fourier transform infrared spectroscopy 50. Bordes et al. 51 did ToF-SIMS investigations combined with XPS. The experiments were carried out on 100 or 300 nm amorphous silicon films and electrochemical experiments were done at C/20 at 10, 30, 50 and 100 % lithiated state. They found also a phase front moving into the electrode. Further, they proved the presence of fast-diffusion paths in the Si active material in which a highly lithiated phase aggregates and leads to lithium segregation at the interface between current collector and active material. Miao and coworkers 52 postulated a model in which a sharp interface between an amorphous LixSi phase and Li-saturated amorphous Si propagates through the film solely based on electrochemical investigations and special peculiarities in the current density vs. time curves. Further in literature there is also a publication where the lithiation of silicon was investigated by ex-situ using Auger Electron Spectroscopy (AES) on sputter deposited 100 nm thick amorphous silicon films lithiated at a rate of C/3 53. From their results they postulate a homogeneous single phase mechanism at least at low lithiation rates in contrast to our findings and also to the XPS data in ref.

20

and ToF-SIMS data in

52,

already mentioned. The reason for these different results is

unclear at the moment. It might be correlated to the measurement method. The fact that the emission depth of Auger electrons is rather low (about 1 nm or less) means that information on Li concentration always come from a surface layer of less than 1 nm, successively modified after each


Page 31 of 39 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


Ar ion etching step. In future, a comparison of SIMS and AES investigation and possible other techniques like XPS on the same sample might be necessary for further clarifications. The conclusion of the literature work available at the present stage of investigation is that the existence of a heterogeneous or a homogeneous mechanism of lithiation of silicon electrodes might be influenced by several factors: current density used, measurement method, sample preparation. Consequently, for future studies the application of different analytical methods on the same sample will be highly desirable.

Conclusion Combined investigations with in-operando Neutron Reflectometry and ex-situ Secondary Ion Mass Spectrometry were carried out in order to study the lithiation and delithiation mechanism of amorphous silicon electrodes. The results show that during the first lithiation half-cycle a heterogeneous two-step process is acting taking pace via moving phase boundaries. First a Li-poor phase penetrates the silicon electrode relatively fast leading to about 10 % of full capacity. Afterwards a Li-rich phase moves into the electrode leading to complete lithiation in a slower process. During lithiation, a Li-poor phase and a Li-rich phase are separated by a relatively sharp interface of several nanometers. A heterogenous mechanism was also found for delithiation, but only a one-step process where the highly lithiated phase is removed from the silicon electrode. The Li introduced in the first process step during lithiation remains in the electrode after delithiation as



Page 32 of 39

irreversible charge. During the second cycle the same process as during the first cycle is indicated, however, in form of a one-step process of the highly lithiated phase only.

Acknowledgements Financial support from the Deutsche Forschungsgemeinschaft (DFG) (Schm 1569/25) is gratefully acknowledged. We thank the Paul Scherrer Institut, Villigen, Switzerland and the HelmholtzZentrum Berlin for provision of neutron beamtime. Further we thank B.-K. Seidlhofer for her assistance during the neutron experiments at Helmholtz-Zentrum Berlin.


Page 33 of 39 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


REFERENCES (1) Aifantis, K. E.; Hackney, S. A.; Kumar, R. V. High energy density lithium batteries: Materials, engineering, applications; Wiley-VCH: Weinheim, 2010. (2) Armand, M.; Tarascon, J.-M. Building better batteries. Nature 2008, 451, 652–657. (3) Yoshio, M.; Brodd, R. J.; Kozawa, A. Lithium-ion batteries: Science and technologies; Springer: New York, NY, 2009. (4) Yuan, X.; Liu, H.; Zhang, J. Lithium-ion batteries: Advanced materials and technologies; Green chemistry and chemical engineering; CRC Press: Boca Raton, FL, 2012. (5) Pistoia, G. Lithium-ion batteries: Advances and applications, First edition; Elsevier: Amsterdam, 2014. (6) Lithium-Ion Batteries: Basics and Applications; Korthauer, R., Ed.; Springer Berlin Heidelberg: Berlin, Heidelberg, 2018. (7) Li, M.; Lu, J.; Chen, Z.; Amine, K. 30 years of lithium-ion batteries. Adv. Mater. 2018, 1800561. (8) Liu, J. Addressing the grand challenges in energy storage. Adv. Funct. Mater. 2013, 23, 924–928. (9) McDowell, M. T.; Lee, S. W.; Nix, W. D.; Cui, Y. Understanding the lithiation of silicon and other alloying anodes for lithium-ion batteries. Adv. Mater. 2013, 25, 4966–4985. (10) Obrovac, M. N.; Chevrier, V. L. Alloy negative electrodes for Li-ion batteries. Chem. Rev. 2014, 114, 11444–11502. (11) Obrovac, M. N.; Christensen, L. Structural changes in silicon anodes during lithium insertion/extraction. Electrochem. Solid-State Lett. 2004, 7, A93. (12) Jerliu, B.; Hüger, E.; Dörrer, L.; Seidlhofer, B.-K.; Steitz, R.; Oberst, V.; Geckle, U.; Bruns, M.; Schmidt, H. Volume expansion during lithiation of amorphous silicon thin film electrodes studied by inoperando neutron reflectometry. J. Phys. Chem. C 2014, 118, 9395–9399. (13) Huang, S.; Zhu, T. Atomistic mechanisms of lithium insertion in amorphous silicon. J. Power Sources 2011, 196, 3664–3668.



Page 34 of 39

(14) Chon, M. J.; Sethuraman, V. A.; McCormick, A.; Srinivasan, V.; Guduru, P. R. Real-time measurement of stress and damage evolution during initial lithiation of crystalline silicon. Phys. Rev. Lett. 2011, 107, 45503. (15) Jerliu, B.; Dorrer, L.; Huger, E.; Borchardt, G.; Steitz, R.; Geckle, U.; Oberst, V.; Bruns, M.; Schneider, O.; Schmidt, H. Neutron reflectometry studies on the lithiation of amorphous silicon electrodes in lithiumion batteries. Phys. Chem. Chem. Phys. 2013, 15, 7777–7784. (16) Jerliu, B.; Hüger, E.; Horisberger, M.; Stahn, J.; Schmidt, H. Irreversible lithium storage during lithiation of amorphous silicon thin film electrodes studied by in-situ neutron reflectometry. J. Power Sources 2017, 359, 415–421. (17) Nelson, A. Co-refinement of multiple-contrast neutron/X-ray reflectivity data using MOTOFIT. J Appl. Crystallogr. 2006, 39, 273–276. (18) Stahn, J.; Glavic, A. Focusing neutron reflectometry: Implementation and experience on the TOFreflectometer Amor. Nucl. Instrim. Meth. A 2016, 821, 44–54. (19) Stahn, J.; Filges, U.; Panzner, T. Focusing specular neutron reflectometry for small samples. Eur. Phys. J. Appl. Phys. 2012, 58, 11001. (20) Hüger, E.; Jerliu, B.; Dörrer, L.; Bruns, M.; Borchardt, G.; Schmidt, H. A Secondary Ion Mass Spectrometry study on the mechanisms of amorphous silicon electrode lithiation in Li-ion batteries. Z. Phys. Chem. 2015, 229.. (21) Janson, M. S.; Hallén, A.; Linnarsson, M. K.; Svensson, B. G. Hydrogen diffusion, complex formation, and dissociation in acceptor-doped silicon carbide. Phys. Rev. B 2001, 64, 13549. (22) Deline, V. R.; Katz, W.; Evans, C. A.; Williams, P. Mechanism of the SIMS matrix effect. Appl. Phys. Lett. 1978, 33, 832–835. (23) Seidlhofer, B.-K.; Jerliu, B.; Trapp, M.; Huger, E.; Risse, S.; Cubitt, R.; Schmidt, H.; Steitz, R.; Ballauff, M. Lithiation of crystalline silicon as analyzed by operando neutron reflectivity. ACS Nano 2016, 10, 7458-7466.


Page 35 of 39 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


(24) Haftbaradaran, H.; Gao, H.; Curtin, W. A. A surface locking instability for atomic intercalation into a solid electrode. Appl. Phys. Lett. 2010, 96, 91909. (25) Liu, X. H.; Fan, F.; Yang, H.; Zhang, S.; Huang, J. Y.; Zhu, T. Self-limiting lithiation in silicon nanowires. ACS Nano 2013, 7, 1495–1503. (26) McDowell, M. T.; Ryu, I.; Lee, S. W.; Wang, C.; Nix, W. D.; Cui, Y. Studying the kinetics of crystalline silicon nanoparticle lithiation with in situ transmission electron microscopy. Adv. Mater. 2012, 24, 6034–6041. (27) Gu, M.; Li, Y.; Li, X.; Hu, S.; Zhang, X.; Xu, W.; Thevuthasan, S.; Baer, D. R.; Zhang, J.-G.; Liu, J. et al. In situ TEM study of lithiation behavior of silicon nanoparticles attached to and embedded in a carbon matrix. ACS Nano 2012, 6, 8439–8447. (28) McDowell, M. T.; Lee, S. W.; Harris, J. T.; Korgel, B. A.; Wang, C.; Nix, W. D.; Cui, Y. In situ TEM of two-phase lithiation of amorphous silicon nanospheres. Nano Lett. 2013, 13, 758–764. (29) Wang, J. W.; He, Y.; Fan, F.; Liu, X. H.; Xia, S.; Liu, Y.; Harris, C. T.; Li, H.; Huang, J. Y.; Mao, S. X. et al. Two-phase electrochemical lithiation in amorphous silicon. Nano Lett. 2013, 13, 709–715. (30) Bard, A. J.; Faulkner, L. R. Electrochemical methods: Fundamentals and applications, 2. ed.; Wiley: Hoboken, NJ, 2001. (31) Hüger, E.; Dörrer, L.; Schmidt, H. Permeation, solubility, diffusion and segregation of lithium in amorphous silicon layers. Chem. Mater. 2018, 30, 3254–3264. (32) Guo, H.; Zhao, H.; Yin, C.; Qiu, W. A. Nanosized silicon thin film as high capacity anode material for Li-ion rechargeable batteries. Mater. Sci. Engin. B 2006, 131, 173–176. (33) Cui, L.-F.; Ruffo, R.; Chan, C. K.; Peng, H.; Cui, Y. Crystalline-amorphous core-shell silicon nanowires for high capacity and high current battery electrodes. Nano Lett. 2009, 9, 491–495. (34) Beattie, S. D.; Loveridge, M. J.; Lain, M. J.; Ferrari, S.; Polzin, B. J.; Bhagat, R.; Dashwood, R. Understanding capacity fade in silicon based electrodes for lithium-ion batteries using three electrode cells and upper cut-off voltage studies. J. Power Sources 2016, 302, 426–430.



Page 36 of 39

(35) Nie, M.; Abraham, D. P.; Chen, Y.; Bose, A.; Lucht, B. L. Silicon solid electrolyte interphase (SEI) of lithium ion battery characterized by microscopy and spectroscopy. J. Phys. Chem. C 2013, 117, 13403– 13412. (36) Jerliu, B.; Dörrer, L.; Hüger, E.; Seidlhofer, B.-K.; Steitz, R.; Borchardt, G.; Schmidt, H. Electrochemical lithiation of silicon electrodes: neutron reflectometry and secondary ion mass spectrometry investigations. Int. J. Mat. Res. 2017, 108, 999 -1007. (37) Veith, G. M.; Doucet, M.; Baldwin, J. K.; Sacci, R. L.; Fears, T. M.; Wang, Y.; Browning, J. F. Direct determination of solid-electrolyte interphase thickness and composition as a function of state of charge on a silicon anode. J. Phys. Chem. C 2015, 119, 20339–20349. (38) Fears, T. M.; Doucet, M.; Browning, J. F.; Baldwin, J. K. S.; Winiarz, J. G.; Kaiser, H.; Taub, H.; Sacci, R. L.; Veith, G. M. Evaluating the solid electrolyte interphase formed on silicon electrodes: a comparison of ex situ X-ray photoelectron spectroscopy and in situ neutron reflectometry. Phys. Chem. Chem. Phys. 2016, 18, 13927–13940. (39) Bordes, A.; Vito, E. de; Haon, C.; Secouard, C.; Montani, A.; Marcus, P. Investigation of lithium insertion mechanisms of a thin-film Si electrode by coupling time-of-flight secondary-ion mass spectrometry, X-ray photoelectron spectroscopy, and focused-ion-beam/SEM. ACS Appl. Mater. Inter. 2015, 7, 27853–27862. (40) Remeš, Z.; Vaněček, M.; Torres, P.; Kroll, U.; Mahan, A.H.; Crandall, R.S. Optical determination of the mass density of amorphous and microcrystalline silicon layers with different hydrogen contents. J. NonCryst. Solids 1998, 227-230, 876–879. (41) Veith, G. M.; Doucet, M.; Sacci, R. L.; Vacaliuc, B.; Baldwin, J. K.; Browning, J. F. Determination of the solid electrolyte interphase structure grown on a silicon electrode using a fluoroethylene carbonate additive. Sci. Rep. 2017, 7, 6326. (42) Beaulieu, L. Y.; Hatchard, T. D.; Bonakdarpour, A.; Fleischauer, M. D.; Dahn, J. R. Reaction of Li with alloy thin films studied by in situ AFM. J. Electrochem. Soc. 2003, 150, A1457.


Page 37 of 39 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


(43) Choi, Y. S.; Pharr, M.; Oh, K. H.; Vlassak, J. J. A simple technique for measuring the fracture energy of lithiated thin-film silicon electrodes at various lithium concentrations. J. Power Sources 2015, 294, 159– 166. (44) Breitung, B.; Baumann, P.; Sommer, H.; Janek, J.; Brezesinski, T. In situ and operando atomic force microscopy of high-capacity nano-silicon based electrodes for lithium-ion batteries. Nanoscale 2016, 8, 14048–14056. (45) Becker, C. R.; Strawhecker, K. E.; McAllister, Q. P.; Lundgren, C. A. In situ atomic force microscopy of lithiation and delithiation of silicon nanostructures for lithium ion batteries. ACS Nano 2013, 7, 9173– 9182. (46) Dong, K.; Markötter, H.; Sun, F.; Hilger, A.; Kardjilov, N.; Banhart, J.; Manke, I. In situ and operando tracking of microstructure and volume evolution of silicon electrodes by using synchrotron X-ray imaging. Chem. Sus. Chem 2019, 12, 261–269. (47) Duay, J.; Schroder, K. W.; Murugesan, S.; Stevenson, K. J. Monitoring volumetric changes in silicon thin-film anodes through in situ optical diffraction microscopy. ACS Appl. Mater. Inter. 2016, 8, 17642– 17650. (48) Cao, C.; Steinrück, H.-G.; Shyam, B.; Stone, K. H.; Toney, M. F. In situ study of silicon electrode lithiation with X-ray reflectivity. Nano Lett. 2016, 16, 7394–7401. (49) Leenheer, A. J.; Jungjohann, K. L.; Zavadil, K. R.; Harris, C. T. Phase boundary propagation in Lialloying battery electrodes revealed by liquid-cell transmission electron microscopy. ACS Nano 2016, 10, 5670–5678. (50) Koo, B. M.; Corte, D. A. D.; Chazalviel, J.-N.; Maroun, F.; Rosso, M.; Ozanam, F. Lithiation mechanism of methylated amorphous silicon unveiled by operando ATR-FTIR spectroscopy. Adv. Energy Mater. 2018, 8, 1702568.



Page 38 of 39

(51) Bordes, A.; Vito, E. de; Haon, C.; Boulineau, A.; Montani, A.; Marcus, P. Multiscale investigation of silicon anode Li insertion mechanisms by time-of-flight secondary ion mass spectrometer imaging performed on an in situ focused ion beam cross section. Chem. Mater. 2016, 28, 1566–1573. (52) Miao, J.; Thompson, C. V. Kinetic study of the initial lithiation of amorphous silicon thin film anodes. J. Electrochem. Soc. 2018, 165, A650-A656. (53) Wang, X.; Singh, S. S.; Ma, T.; Lv, C.; Chawla, N.; Jiang, H. Quantifying electrochemical reactions and properties of amorphous silicon in a conventional lithium-ion battery configuration. Chem. Mater. 2017, 29, 5831–5840.


Page 39 of 39 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


TOC graphic


On the Lithiation Mechanism of Amorphous Silicon ... - ACS Publications

Recommend Documents