Diplomarbeit - TU Wien · 2019. 12. 11. · Technische Universität Wien Diplomarbeit On the Electrochemical Properties of Li-ion Conducting Li 7 La 3 Zr 2 O 12 and its Utilization

Technische Universität Wien

Diplomarbeit

On the Electrochemical Properties of Li-ion Conducting

Li7La3Zr2O12 and its Utilization in All Solid State Li-ion

Batteries

Ausgeführt am

Institut für Chemische Technologien und Analytik der

Technischen Universität Wien unter Anleitung von

Univ.-Prof. Dipl.-Phys. Dr. rer. nat. Jürgen Fleig

Projektass. Dipl.-Ing. Stefan Smetaczek

durch

Joseph Ring, BSc.

15. März 2019

Die approbierte Originalversion dieser Diplom-/ Masterarbeit ist in der Hauptbibliothek der Tech-nischen Universität Wien aufgestellt und zugänglich.

http://www.ub.tuwien.ac.at

The approved original version of this diploma or master thesis is available at the main library of the Vienna University of Technology.

http://www.ub.tuwien.ac.at/eng

Abstract

In this work, the impact of eld stress on Li7La3Zr2O12 (LLZO) polycrystalline pel-

lets and single crystals was investigated. Voltages of up to a few V were applied

using metal electrodes at elevated temperatures in dierent gas atmospheres (argon

and air). The electrochemical processes induced by eld stress were investigated

by a combination of impedance spectroscopy and chemical analysis. Local con-

ductivities were determined using circular thin lm micro-electrodes prepared by

photolithography and ion-beam etching. Laterally resolved elemental compositions

of the samples were measured by laser-induced breakdown spectroscopy (LIBS).

The correlation between changes in electrical properties and chemical composition

is discussed.

Moreover, electrochemical properties of Li containing oxide thin lm electrodes

on LLZO were investigated by means of impedance spectroscopy and DC-cycling.

In order to prepare such electrodes (e.g. LiMn2O4 and Li4Ti5O12) a RF sputter-

ing device was built and tested for dierent materials. Inuence of the substrate

temperature on the properties of the thin lms were investigated. All solid state

cells were built with Ta-doped LLZO single crystals and polycrystals and reversible

electrochemical (dis)charging of the oxide electrodes was demonstrated.

1

Acknowledgments

The contribution of all LLZO single crystals by S. Ganschow1 and S. Berendts2

is gratefully acknowledged. Polycrystalline pellets were prepared by the group of

D. Rettenwander3, whose contribution to this work is also gratefully acknowledged.

Samples of alumina powders (CT1200 SG and CT3000 SG), courtesy of Almatis

Germany, were used for the fabrication of heating elements.

1Leibnitz Institute for Crystal Growth, Germany2Berlin University of Technology, Germany3TU Graz, Austria

2

Contents

1 Introduction 5

1.1 A Brief Introduction to Lithium ion Batteries . . . . . . . . . . . . . 5

1.2 Solid Li Ion Conducting Electrolytes . . . . . . . . . . . . . . . . . . 11

1.2.1 The Garnet-Type Material Li7La3Zr2O12 . . . . . . . . . . . . 13

1.3 Motivation and Aim . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

I Eects of Field Stress on the Electrochemical Properties

of the Solid State Electrolyte LLZO 19

2 Experimental 20

2.1 Sample Preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.2 Field Stress Experiments with stripe electrodes . . . . . . . . . . . . 21

2.2.1 Electrochemical Methods . . . . . . . . . . . . . . . . . . . . . 21

2.2.2 Chemical Analysis . . . . . . . . . . . . . . . . . . . . . . . . 23

2.3 Field Stress Experiments with Microelectrodes . . . . . . . . . . . . . 24

3 Results and Discussion 26

3.1 Results from Polarization of Stripe Electrodes . . . . . . . . . . . . . 26

3.1.1 ME-EIS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

3.1.2 LIBS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

3.2 Results from Polarization of Microelectrodes . . . . . . . . . . . . . . 33

3.2.1 Polarization Current Measurements . . . . . . . . . . . . . . . 33

3.2.2 Inuence of the Atmosphere . . . . . . . . . . . . . . . . . . . 36

3.2.3 Inuence of the Temperature . . . . . . . . . . . . . . . . . . 37

3.2.4 Results from a Broader Voltage Range . . . . . . . . . . . . . 39

3

3.2.5 U-I Characteristic . . . . . . . . . . . . . . . . . . . . . . . . . 40

3.2.6 LIBS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

II Sputtering and Testing of LIB Electrode Materials 48

4 Construction of a Sputtering Device 49

4.1 Preparation and Characterization of Thin Films . . . . . . . . . . . . 50

5 Results and Discussion 54

5.1 Film Thickness and Deposition Rates . . . . . . . . . . . . . . . . . . 54

5.2 Crystal Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56

5.3 Electrochemical Properties of Sputtered Thin Films . . . . . . . . . . 59

5.3.1 Impedance Spectroscopy . . . . . . . . . . . . . . . . . . . . . 59

5.3.2 DC Cycling . . . . . . . . . . . . . . . . . . . . . . . . . . . . 60

6 Conclusion 69

4

Chapter 1

Introduction

1.1 A Brief Introduction to Lithium ion Batteries

Energy conversion and storage are key challenges of modern life. Despite the great

eorts of recent years to establish renewable, emission-free energy technologies [1

5], conventional fuels are still necessary to fulll today's energy demands. Many

renewable energy sources are subject to natural uctuations (time of day, weather,

...) which do not match the power demand. Unable to meet the demand at one time

and providing excess power which cannot be consumed at another time, renewable

energy sources lack ecient energy storage technologies in order to replace fossil

fuels.

Mobility is another challenge for renewable energy sources. Many renewable

energy sources are not mobile (wind power, etc.), whereas others cannot meet the

power demand (e.g. solar cells). It is therefore essential for mobile applications, to

store energy from renewable sources in a way that provides high energy densities.

Especially for mobility applications (e.g. electric cars), safety of energy storage

devices is a crucial requirement.

Energy can be stored by various dierent approaches. These include physical

(e.g. pumping water uphill), electric (e.g. loading capacitors) or chemical (e.g. elec-

trolysis) methods and many more, each of which have some advantages, depending

on the application. Long-term energy storage, usable for mobile applications, is

best achieved chemically. All chemical energy storage methods are based on the

same concept: An external power source drives an endergonic reaction, which can

5

later be reversed in order to release energy, when needed. Depending on the appli-

cation, the reactions involved can be quite dierent. Electrochemical reactions are

especially useful, because they can release the stored energy as electric energy. The

released energy can therefore be used directly to drive an electric engine, making the

entire system very energy ecient, compared to combustion engine systems. Two

important electrochemical energy storage concepts are:

1. Electrochemical synthesis of fuels: Electrolysis of water is one example of

converting electric energy into fuel. Hydrogen, which is released from the

water splitting reaction can then be used in fuel cells to release the stored

energy.

2. Batteries: Reduction and oxidation of electrode materials can be used to store

electric energy. The reverse reaction then releases the stored energy.

Fuels produced by electrolysis can be used for various mobile and stationary

applications. One advantage of renewable fuels, such as hydrogen, compared to

batteries is the fact that not all chemicals have to be carried alongside the fuel.

When e.g. hydrogen is used in a fuel cell, air from the environment can be used as

an oxygen source, improving the energy density. Further, reaction products, such

as water can readily be released into the environment. Despite its high gravimetric

energy density (about 30 kWh/kg), hydrogen has a rather low volumetric energy

density (<1 kWh/L), which can be improved by compression (several hundred bars)

or chemical modications (producing e.g. methane). The need for an infrastructure

for transporting and distributing the fuel can be challenging, especially for high-

pressure hydrogen.

Batteries are galvanic cells, which can store and release electric energy by chem-

ical redox reactions. Depending on the type of battery, the storage and release of

energy can be repeated over many cycles (several hundred), retaining a satisfactory

energy density. Whereas fuel cells only need fuel and air for operation, all reactants

and reaction products need to remain inside a rechargeable battery. This reduces

the energy density of the device. One great advantage of batteries, however, is that

they can be charged directly with electric power, circumventing the need of a sepa-

rate production plant and fuel transport.

6

Lithium ion batteries (LIBs) have become the most important energy storage

device for portable electronics. They have both high volumetric and gravimetric

energy densities, because Li is a light element (6.94 g · mol−1) with a very high

reduction potential (3.04 V vs. H2/H+). In 1980, Mizushima and coworkers [6]

designed the rst, rechargeable Lithium ion battery utilizing LiCoO2 as a cath-

ode material and Li metal as an anode. They showed high open circuit voltages

(4-5 V) due to the high oxidation potential of the couple Co3+/Co4+ and overall,

very promising electrochemical properties for practical applications. In 1991, Sony

released the rst commercial product, containing a rechargeable LIBs and revolu-

tionized modern battery technology. Since then, great research eort was made to

improve capacity, safety, cycle life, and other properties of LIBs. The performance

of LIBs has indeed improved tremendously over the last decades, but the basic cell

design has not changed much. In principle, a LIB consist of:

1. Cathode: Presently, most commercially used cathode materials are layered

oxides, such as LiCoO2, LiNi0.33Mn0.33Co0.33O2 or spinel-type LiMnO2.

2. Electrolyte: Liquid electrolytes, absorbed in a membrane separator are used

in most batteries. They consist of a lithium salt, such as LiPF6 dissolved in a

liquid, organic solvent, such as ethylene carbonate. The membrane separator

does not aect the conductivity but protects the cell mechanically from short

circuits.

3. Anode: Graphite is the most important anode material of today`s batteries

with Li4Ti5O12 acting as one alternative.

Common cathode and anode materials are listed in Tables 1.1 and 1.2, respectively.

Table 1.1: Energy densities of some common cathode materials [7].

CompoundSpecic capacity(mAh· g−1)(theoretical)

Volumetric capacity(mAh· cm−3)(theoretical)

Average voltage(V)(Vs. Li/Li+)

LiCoO2 274 1363 3.8LiMn2O4 148 596 4.1LiMnO2 285 1148 3.3LiFePO4 170 589 3.4LiNi0.8Co0.15Al0.05O2 279 1284 3.7

7

Table 1.2: Properties of some common anode materials [7].

CompoundSpecic capacity(mAh· g−1)

Volumetric capacity(mAh· cm−3)

Average voltage(V)(Vs. Li/Li+)

Graphite 726-946 330-430 0.1Lithium metal 3862 7232 0Li4Ti5O12 320 600 1.55

The working principle of LIBs can be summarized as follows (s. Fig. 1.1): In

charged state, the anode material (e.g. LiC6) is lled with Li-ions and the cathode

(e.g. Li0.5CoO2) is depleted of Li-ions. In this state, the Li has a high chemical

potential inside the anode and a low chemical potential inside the cathode. This

gradient in chemical potential serves as the driving forces for the discharging reac-

tion. The electrolyte, however, separates the two electrodes, allowing only Li-ions

to pass through the inside of the battery. Therefore, electrons must be transported

through an external circuit when charging or discharging. During discharge, the an-

ode is oxidized, releasing electrons through external leads. The cathode is reduced,

receiving electrons from the external leads. The electrochemical energy of the re-

dox reaction is converted to useful electric energy because the electrolyte does not

conduct electrons, but only Li-ions through the inside of the battery. The chemical

reaction taking place during discharge can be written as two half-reactions.

Anodic half-reaction: LiC6 −−→ Li+ + e +C6

Cathodic half-reaction: 2 Li0.5CoO2 + Li+ + e −−→ 2 LiCoO2

The discharging reaction can be reversed by an external bias, recharging the battery

back to charged state.

At rst glance, the working principle of LIBs appears simple. In practice, how-

ever, various other processes can take place inside a LIB, which need to be consid-

ered. In order to make a useful energy storage device, a LIB needs to meet certain

criteria: High energy densities (gravimetric and volumetric), long cycle life with

8

Figure 1.1: Schematic illustrating the working principle of lithium ion batteries.Upon discharge, Li-ions (red spheres) traverse the inside of the battery throughthe electrolyte (blue), oxidizing the graphite anode (left) and reducing the layeredoxide cathode (right). Electrons pass through the current collectors (hatched bars)and the resistive load, doing useful work. Recharging of the battery is achieved byapplying an external bias, driving the reaction in reverse direction. The pictureelements are not drawn in correct scale.

little degradation, high charging and discharging rates, high energy eciency, good

stability (thermal, mechanical and chemical), and safety. These requirements impose

high demands on cell design and materials properties, making LIBs sophisticated

devices. Obviously, cathode and anode materials need to have a high, respectively

low potential vs. Li/Li+ in order to make a high voltage cell. Further, lithiation

and delithiation of the materials needs to be reversible over many cycles. Layered

compounds are widely used as electrode materials because they can be (de)lithiated

by (de)intercalation of Li. Open cell voltages of modern LIBs are as high as 5 V [8],

but energy densities of LIBs have to be increased further for mobility applications.

Cathode materials must contain at least one other element, which can be re-

versibly reduced, giving a high reduction potential. Layered transition metal oxides

are the most suitable materials for this. Since these compounds contain not only

Li, but numerous other atoms per formula unit, the energy density is fairly low.

9

Further, most cathode materials cannot be fully delithiated, because the resulting

compound would decompose irreversibly into more stable phases, decreasing the en-

ergy density even more. Anode materials, on the other hand, should have a low

reduction potential. Naturally, Li metal has the lowest possible reduction potential

versus Li/Li+ (0 V). Since there are no atoms other than Li inside Li metal, the

energy density of Li metal is very high. However, certain problems (dendrite forma-

tion, reactions with electrolyte) arising from Li metal anodes are hard to overcome.

As a result, graphite is the most commonly used anode material of today's batter-

ies, because it has comparably good characteristics, but does not suer from the

problems mentioned above.

The electrolyte must be a good Li-ion conductor (σLi > 10−4 S/cm), but a poor

electronic conductor (σe < 10−10 S/cm). Furthermore, it needs to be stable against

chemical reactions with the electrodes. Aqueous electrolytes e.g. cannot be used

in high-voltage devices, because water is not stable at voltages exceeding 1.23 V.

Presently, most LIBs use Li-salts (such as LiPF6 or LiClO4) dissolved in organic,

aprotic solvents (e.g. ethylene carbonate, dimethyl carbonate) as an electrolyte.

These electrolytes have high ionic conductivities (> 10−2 S/cm) and practically no

electronic conductivity. Organic electrolytes, however, are ammable, which can

lead to devices bursting into ames or even exploding in case of failure [9]. This

safety hazard is a major obstacle for the application of LIBs in electric vehicles and

other technology elds. However, ammability is not the only aw of liquid organic

electrolytes. Li metal cannot be used as an anode material in combination with liquid

organic electrolytes. This is partly because they are not chemically stable against Li

metal, which can lead to irreversible capacity losses. Moreover short circuits due to

Li-dendrite formation may occur. The issues imposed on liquid, organic electrolytes

raise the question for alternative electrolytes, such as ionic liquids, solid polymer

electrolytes and inorganic solid state electrolytes. Much research eort has been

focused on solid state, inorganic Li-electrolytes [1012], because they are promising

to overcome the problems of organic-based electrolytes. Further, inorganic solid

electrolytes could even enable the use of Li metal as an anode material and high

voltage cathodes within one cell. In recent years, great progress was made [1316]

in the search for highly Li-conductive (> 10−3 S/cm) ceramic electrolytes, making

10

a big step towards the utilization of solid state electrolytes in LIBs.

1.2 Solid Li Ion Conducting Electrolytes

Most research on solid Li electrolytes is focused on crystalline materials. It is noted,

however, that highly conductive amorphous Li electrolytes have been reported [17

19]. Most solid Li electrolytes are either oxide [10, 20, 21] or sulde [18, 22, 23] based.

In general, sulde based electrolytes exhibit higher Li conductivities, whereas oxides

are chemically more stable. Both types of electrolytes, however, can reach high Li-ion

conductivities (> 10−4 S/cm) at room temperature as well as wide electrochemical

stability windows (exceeding 5 V) [24], making them promising electrolyte materials

for all solid state type LIBs.

Beside high ionic and low electronic conductivity, electrochemical stability of the

electrolyte against the electrodes is a crucial requirement for its implementation in

a LIB. Electrochemical stability is usually referred to the stability against oxidation

or reduction of the electrolyte by the cathode or anode, respectively. Oxidation

of the electrolyte can occur if the electrolyte contains electrons in higher energy

states than the lowest vacant electronic state inside the cathode. This can, from a

thermodynamic point of view, result in the transfer of electrons from the electrolyte

into the cathode material (i.e. oxidation of the electrolyte). Naturally, the electron

transfer is accompanied by chemical reactions, which are irreversible and undesired.

Similar to the oxidation, the electrolyte can be reduced by the anode. This case

requires a vacant electronic state inside the electrolyte and an electron with higher

energy inside the anode. Both, reduction and oxidation reactions are undesired

because they result in an irreversible capacity loss. Further, safety hazards can result

from uncontrolled chemical reactions between electrolyte and electrode materials.

Compatibility of the electrolyte and the electrode materials is therefore an important

aspect of cell design.

The compatibility of electrolytes and electrode materials can be described by a

simple picture (see Fig. 1.2). The highest occupied electronic state and the lowest

unoccupied electronic state of the electrolyte mark the limits of its electrochemical

stability window. In order to make a thermodynamically stable cell, the occupied,

11

respectively unoccupied, electronic states of the electrode materials have to lie in-

side this window. If this is not the case, undesired chemical reactions can occur,

as described above. This means, that the electrochemical stability window of the

electrolyte limits the voltage of the cell. The width as well as the energetic location

of the stability window is therefore an important material property for electrolytes.

Figure 1.2: Schematic energy diagram of a LIB. The electrochemical stability win-dow (Eg) of the electrolyte (orange) spreads from about 0.5 V to around 4 V (withrespect to Li metal). The cell drawn in this schematic is thermodynamically unsta-ble, because the electrode materials do not lie within the stability window. However,the formation of an electronically insulating interface layer (SEI) can stabilize the cellkinetically, preventing the cell from further degradation due to energy mismatches,thus enabling high open-circuit voltages (VOC).

When considering an organic electrolyte, the stability window is dened by the

molecular orbitals (HOMO and LUMO) of the organic molecules of the solvents. For

crystalline solids, this corresponds to the valence band maximum and conduction

band minimum, respectively. Ideally, an electrolyte has a wide electrochemical sta-

bility window, that extends from less than 0 V to a voltage beyond 5 V. However,

even when energetic mismatches between the electrolyte and electrode materials

12

occur, the cell can become stable. This is the case, when the reaction between elec-

trolyte and electrode leads to the formation of a stable solid/electrolyte interface

layer (SEI) [25], which has good ionic but negligible electronic conductivity, thus

preventing the oxidation/reduction from proceeding further, without blocking the

Li-ions from passing between electrolyte and electrode. The formation of an SEI,

of course, is accompanied by a capacity loss, but when a stable SEI is formed, this

loss does not recur after the rst charging of the device. The formation of an SEI

can therefore extend the electrochemical stability window of an electrolyte, enabling

higher cell voltages. When the SEI is chemically or mechanically unstable, however,

a continuous capacity loss takes place, reducing the cycle life of the cell. The impor-

tance of SEIs for the cell performance makes the interface between electrolyte and

electrodes a very interesting research subject [2630]. Some researchers even argue,

that ion transport across an electrode-electrolyte is more important than the bulk

ionic conductivity of the electrolyte [31].

1.2.1 The Garnet-Type Material Li7La3Zr2O12

Presently, many dierent solid Li-electrolytes are known [24]. Due to their high

chemical stability and good ionic conductivity, oxide materials are among the most

promising candidates for the utilization in LIBs. Owing their exceptionally high

bulk ionic conductivity, garnet-type Li-oxides have received special attention [16,

3244] since they were introduced to the LIB community in 2003 by Thangadurai

et al. [45].

Most research eorts on garnet-type electrolytes have been focused on Li7La3Zr2O12

(LLZO) and variations (structure, Li-content, doping, substitution) of it. The garnet

structure of LLZO is complex. It can be described by coordination-polyhedra, i.e.

LaO8 dodecahedra and ZrO6 octahedra, that share common edges. The interstice of

this three-dimensional framework can accommodate a maximum of 9 Li+-ions per

formula unit, distributed among the possible tetrahedral and octahedral Li-sites.

At room-temperature, pure LLZO is only stable in its tetragonal crystal structure

(s. Fig. 1.3), which is a slightly distorted form of the cubic garnet structure. The

tetragonal structure belongs to space group No 142 (I41 / acd) and shows a rather

low ionic conductivity (around 10−6 S/cm) at room temperature [32]. In this struc-

13

Figure 1.3: Crystal structure of tetragonal Li7La3Zr2O12. The solid box indicatesthe unit cell. The three dierent types of Li-sites are shown on the right. Graphicreproduced with permission of Elsevier [32].

ture, two thirds of the tetrahedral sites (16 out of 24) are vacant, while all of the

octahedral sites (48) are occupied by Li-atoms, resulting in a total of 56 occupied

(out of a total of 72) Li-sites per unit cell. Note, that one unit cell corresponds

to four complete formula units of LLZO (Li7La3Zr2O12). Partial occupation of the

Li-sites is an important prerequisite for high ionic conductivity, since ionic motion

within a crystal is dependent on the presence of vacant sites for the transported

ion. Ions can then move from their initial site to a vacant site in close proximity,

leaving a new vacancy behind. This ionic motion has to be activated, since there

is an energetic barrier between the two neighboring sites, that has to be overcome.

The probability of a jump from one site to a neighboring site, and hence the ionic

conductivity σion can be related to the activation energy (Ea):

σion ∝ exp(− Ea

kBT)

The height of this activation barrier is inuenced by the local crystal structure (i.e.

bond lengths, ionic radii, etc.). The relative number of vacant sites is, of course, an-

other important factor for ionic conductivity. The mechanism of ionic motion inside

14

the complex structure of LLZO is not only inuenced by the relative occupation of

the Li-sites, but also by the distribution among the dierent sites (tetrahedral or

octahedral). In LLZO, tetrahedral and octahedral Li-sites are interconnected to a

3D network. The Li coordination-polyhedra share common, triangular faces. Each

tetrahedral site is surrounded by four octahedral sites, whereas each octahedral site

connects two dierent tetrahedral sites. As mentioned before, the bulk ionic con-

ductivity of tetragonal LLZO is fairly low, compared to the cubic polymorph (>

10−4 S/cm). This dierence is attributed to small but signicant dierences of the

crystal structure. While both structures have the same basic garnet network, the

distribution of Li-ions among the Li-sites in cubic LLZO is dierent and not as or-

dered as in tetragonal LLZO. This leads to a far higher ionic conductivity in the

cubic polymorph of LLZO. A detailed discussion on the ion migration mechanism

in garnets would be beyond the scope of this thesis, thus the reader is referred to

the literature [34, 36, 46, 47]. As mentioned before, the cubic garnet-phase of LLZO

is no stable at room temperature. However, the cubic phase can be stabilized by

doping with various elements [33, 4852]. Doping with aliovalent atoms can also

change the Li-content of the LLZO, further inuencing the conductivity. The high-

est conductivity yet (1.2 · 10−3 S/cm), amongst LLZO garnets, was achieved by

Rettenwander et al. by doping LLZO with Al 3+ and Ga3+.

Beside the high ionic conductivity of LLZO, the high chemical stability is a key

property for its applicability in LIBs, especially for high voltage cells. LLZO has

wide electrochemical stability window (i.e. band gap) of more than 5 V [33, 53]. The

energetic location of the stability window, with respect to desired cathode/anode

materials, is an important aspect as well. Several works on the stability of LLZO

indicate, that LLZO is stable against Li metal [34, 5356]. The interface between

LLZO and Li metal anodes, however, can give a large resistance. This can be

attributed to poor mechanical contact between the two materials, but small contact

area is not the only contributor to the large interface resistance. LLZO does not

exhibit chemical stability against air (i.e. oxygen, moisture and CO2) [37, 43, 57,

58]. The degradation of LLZO in ambient atmosphere leads to the formation of a

Li2CO3 layer. If such a layer is present prior to electrode deposition, it leads to a

high interface resistance between LLZO and the electrode material. Thus, special

15

care must be taken to prevent this, making the processing of LLZO laborious.

The high resistance of the interface between LLZO and cathode materials, such

as LiCoO2 (LCO), is believed to be a key issue for solid state LIBs [31, 59, 60].

Despite its wide stability window, LLZO forms a resistive reaction layer with LCO at

elevated temperatures [6163]. Sintering or annealing of solid state cells at elevated

temperatures, however, is a common strategy to achieve good contact between the

dierent materials. Recent studies showed, that thermally induced inter-diusion

occurs at temperatures as low as 300 C [64], while other researchers claim, that

solid state reactions between LLZO and LCO do not occur below 700 C [63]. The

reported reaction products include La2Zr2O7, LaCoO3, La2CoO4, La2Zr2O7 and

Li2CO3, if CO2 is present in the atmosphere. The resulting reaction layers show

high area specic resistance of several kΩ · cm−2 [64]. Several approaches to this

problem, such as lower processing temperatures or addition of an interlayer, are

suggested. The critical role of the electrolyte-cathode interface presses for extensive

research, to overcome the problems induced by interface reactions.

A reliable synthesis of highly conductive LLZO is an essential prerequisite for

larger scale production. However, reports on conductivities of LLZO with nominally

equal or very similar composition, diverge [62, 65, 66]. This scattering could be due

to dierent synthesis conditions. There are several dierent synthetic approaches for

the production of LLZO (e.g. solid state reaction, mixed precursor method, sol-gel

methods, etc.). Most synthetic routes to LLZO electrolytes include a sintering or

annealing step for densication of the material. The elevated temperature during

these steps can lead to the evaporation of relatively volatile LiO2. The resulting

Li-loss is often compensated by adding excess Li-containing salts (such as LiO2 or

Li2CO3) prior to the sintering step. Elaborate studies [42, 44] by Wachter-Welzl

et al. on nominally equal LLZO samples showed scattering over almost two orders

of magnitude (2 · 10−5 − 8 · 10−4 S · cm−1) in bulk ionic conductivity. Further,

variations in ionic conductivity within single samples were observed by means of

microelectrode impedance spectroscopy. Detailed chemical analysis showed varia-

tions in the elemental composition of the samples. However, no direct correlation

between local conductivity and composition was apparent. Point defects, such as

oxygen vacancies can also aect the Li-ion conductivity. Kubicek et al. observed

16

signicant oxygen vacancy concentrations in doped cubic LLZO by secondary-ion

mass spectroscopy (SIMS) [41]. Oxygen vacancies can inuence the stoichiometry

as well as the the crystal structure of LLZO. Both eects can have an impact on the

conductivity, thus, a complex relationship between oxygen vacancy concentration

and ionic conductivity was observed. Interestingly, at 350 C, an unexpectedly high

oxygen tracer diusivity (diusion coecient as high as 8.2 · 10−12 cm2 · s−1) was

observed. Despite the great research eorts, the complex interplay of stoichiometry,

defect concentration and conductivity remains unclear.

1.3 Motivation and Aim

The cubic modication of LLZO exhibits excellent properties for a solid state Li

electrolyte. Processing of the material is challenging. Uncontrolled variations of

stoichiometry during synthesis have inuence on the conductivity. The underlying

relation between stoichiometry and Li-ion conductivity is not fully understood. To

get a deeper understanding, the preparation of samples with well dened stoichiom-

etry is crucial. Due to the high temperature steps in most synthetic routes, this is

hardly achievable. Hence, a tool for well-dened modications of the composition

within LLZO samples is desirable.

One aim of the experiments was thus the controlled manipulation of local stoi-

chiometry of Ta-doped LLZO single crystals by application of eld stress. Stoichiom-

etry gradients in solid ionic conductors can be induced by so-called Wagner-Hebb

polarization, referring to the rst observations of this eect by C. Wagner [67] and

M. Hebb [68]. After polarization of a single crystal, by means of metal electrodes lo-

cally resolved conductivity measurements as well as chemical analysis was performed

to check for resulting changes in stoichiometry and conductivity. Results from these

experiments could help to understand the relationship between local conductivity

and stoichiometry in more detail.

These experiments came along with other phenomena, such as electron conduc-

tion due to ion blocking electrodes or chemical decomposition of LLZO upon high

eld stress. To get a better understanding, further eld stress experiments were

conducted under varying conditions (atmosphere, temperature, geometry, etc.). All

17

these experiments are described in Part I (chapters 2 to 3).

The second part of the thesis deals with oxide thin lm electrodes on LLZO.

There is still very little agreement on the properties of the interfaces between typical

LIB electrodes (e.g. LiCoO2, LiMn2O4, Li4Ti5O12) and LLZO. Model thin lm

oxide electrodes on single crystalline LLZO may serve as excellent systems for a

more systematic study of these interfaces and the charging-discharging behavior.

Therefore such model systems were prepared and rst measurements were performed

to show the practicability of this approach. In order to deposit such electrodes, an

novel sputtering device was constructed and used.

18

Part I

Eects of Field Stress on the

Electrochemical Properties of the

Solid State Electrolyte LLZO

19

Chapter 2

Experimental

2.1 Sample Preparation

All experiments were conducted with Ta-doped LLZO (LLZTO) single crystals, cour-

tesy of S. Ganschow1 and S. Berendts2. LLZTO single crystals with the nominal

composition Li6La3ZrTaO12 were grown by the conventional Czochralski technique.

The starting materials were Li2CO3, La2O3, ZrO2, and Ta2O5. Carbonates and

oxides were dried and mixed in the required stoichiometry with a 10% excess of

Li2CO3. Afterwards the powders were uni-axial pelletized, isostatically pressed at

2800 kbar and nally sintered at 1373 K for 16 h in air. Capped magnesia crucibles

were used while covering the pellets with the respective LLZTO powder to avoid Li-

loss during sintering. Because of the high melting temperature, the sintered LLZTO

samples were molten by radio frequency induction heating using a 25 kW microwave

generator. An iridium seed (pulling rate 1.5 mm · h−1, rotation speed 1 rpm) was

used for the crystal growth performed under nitrogen atmosphere. An active af-

terheater was applied to adjust the temperature gradient in the set up. Thermal

insulation was established by an outer alumina ceramic tube lled with zirconia

granules. Inductively coupled plasma optical emission spectroscopy corroborated

that the LLZTO single crystal obtained has indeed the composition Li6La3ZrTaO12.

The crystals were cut into slices (about 7 by 5 mm and 1 mm thick). Since surface

degradation of LLZO in air is a known issue, all crystals were mirror-polished with

1Leibnitz Institute for Crystal Growth, Germany2Berlin University of Technology, Germany

20

SiC grinding paper (#4000) prior to any electrode deposition.

Ionically blocking gold electrodes (200 nm thickness) were deposited onto the

crystals by DC-sputtering (Baltec Med 020) at room temperature. Micro-structuring

was done by photo-lithographic techniques and subsequent ion beam etching. Sam-

ples of two, slightly dierent electrode congurations were prepared this way (s.

Fig. 2.1). The bottom face of both the crystals was completely covered with a gold

electrode. The top faces were structured to get arrays of circular microelectrodes

(100 µm diameter). Additional stripe electrodes were prepared on one sample.

Figure 2.1: Schematic illustration of the prepared samples. The bottom face ofboth the samples is covered with a single electrode. Sample (a) has large stripeelectrodes on the top face, microelectrodes are located between. On sample (b),only microelectrodes were prepared.

2.2 Field Stress Experiments with stripe electrodes

Ta-doped single crystals were subjected to electric eld stress at elevated temper-

atures (300 - 450 C) under lab atmosphere (air). The experimental setup is illus-

trated schematically in Fig. 2.2.

2.2.1 Electrochemical Methods

The polarization voltage (2.5 - 5 V) was applied via the large stripe electrodes (s. Fig.

2.1 (a)), for several hours (15 to 68 h) at elevated temperatures. The samples were

heated from below via a lab-built heating unit connected to a Linkam thermal control

unit. Since the temperature is controlled according to a thermocouple inside the

heating unit, the actual sample temperature is lower than the set temperature. Since

no exact knowledge of the actual sample temperature is required for the experiments,

all temperature indications in this work correspond to the set values, which are

21

Figure 2.2: Schematic illustration of the experimental setup. Field stress is appliedvia two opposing stripe electrodes at elevated temperature (1). Laterally resolvedconductivity measurements are done via microelectrode electrochemical impedancespectroscopy, prior to and after the polarization experiments at room temperature(2).

22

higher (about 20 - 50 K) than the actual sample temperatures. A Keithley, 2611

Source Meter unit was used to apply voltage and measure the resulting current. The

polarization experiments were done at elevated temperatures to ensure high ionic

conductivities of Li-ions and at least some mobility of O-ions [41].

Each stripe electrode was contacted together with the opposite stripe electrode.

Thus, the applied electric eld was only present in the area between the correspond-

ing electrode pairs. After a given time upon voltage the sample was cooled to room

temperature (25 C). The bias was still applied during cooling, to keep possible po-

larization eects from relaxing into the initial state of the sample. Cooled to room

temperature, the ionic mobility was considered so low, that no fast relaxation of po-

larization could take place. Between each of the stripe electrode pairs, three arrays

of circular microelectrodes were used for conductivity measurements at room tem-

perature, prior to and after the polarization experiments. These experiments were

performed by means of electrochemical impedance spectroscopy (EIS). For these

measurements, a single microelectrode was contacted as a working electrode to the

impedance analyzer, whereas the large area backing electrode (on the backside of

the crystal) was contacted as the counter-electrode. For the EIS measurements, a

Novocontrol Alpha Analyzer was used, in the frequency range from 1 MHz to 1 kHz.

Owing the relatively small area of the microelectrodes, the probed volume of the

EIS is located directly underneath the microelectrodes, which allows for laterally

resolved conductivity measurements, as has been shown by J. Fleig [69]. The tech-

nique of microelectrode-EIS (ME-EIS) was applied successfully to investigate local

conductivities in LLZO [42].

2.2.2 Chemical Analysis

After the polarization experiments and subsequent ME-EIS measurements, the ele-

mental composition of the sample was determined by laterally resolved laser induced

breakdown spectroscopy (LIBS). This technique uses a laser beam which is focused

onto the sample surface. Upon impact of the laser beam, material is ablated and

excited. Upon relaxation, the present elements emit characteristic light, which is

measured by a detector. The intensities of the emitted light can be related to rel-

ative elemental compositions within the sample. Line-scans were performed across

23

Figure 2.3: Schematic illustration of a LIBS measurement. While the laser beamis moved over the surface of the sample, material gets ablated and excited. Uponrelaxation, characteristic radiation is emitted, indicating the elemental compositionof the scanned material.

the polarization axes, since a stoichiometry gradient was expected to be present

along the axes (s. Fig. 2.2). Due to matrix eects, the LIBS measurements cannot

be used directly for absolute quantication of elemental composition. A laborious

standardizing procedure would be required for this, which would be beyond the

scope of this work. However, no absolute quantication of the elemental compo-

sition is needed for this experiment. The relative composition within the sample

yields enough information to observe a possible stoichiometry gradient, caused by

the electric eld.

2.3 Field Stress Experiments with Microelectrodes

Further polarization experiments, with more focus on the polarization current, were

done, using a dierent sample geometry (s. Fig. 2.1 (b)). Field stress (from -2 V up

to 2 V) was applied by contacting either two neighboring MEs, or one ME and the

large backside electrode at elevated temperature (300 C) and room temperature (s.

Fig. 2.4). The atmosphere was either air or argon. Using only one or two ME for

each experiment, a large number of experiments could be done on one crystal.

LIBS measurements were done after polarization, to investigate possible changes

in elemental composition due to eld stress. As only MEs were contacted during

polarization, the laser was not scanned over the sample, but focused on individual

24

MEs to obtain information about the local composition.

Figure 2.4: Schematic illustration of the polarization experiments using microelec-trodes. Two neighboring MEs (a) or a single ME and the large counter electrode(b) were contacted. Field stress was applied, measuring the resulting current.

25

Chapter 3

Results and Discussion

3.1 Results from Polarization of Stripe Electrodes

The impact of eld stress on electrochemical properties of LLZTO was investigated

by means of ME-EIS (ionic conductivity) and LIBS (elemental composition). Both

LIBS and EIS measurements were done prior to and after the polarization experi-

ments, to reveal any change caused by eld stress.

3.1.1 ME-EIS

Since most of the voltage between a ME and a large counter electrode (covering the

entire bottom face of the sample) drops very close to the ME, the diameter of the

electrode determines the probed volume beneath [69]. The resistances, measured my

EIS are mostly determined by the conductivity of a hemisphere with a radius of 2d,

where d is the diameter of the ME. Fig. 3.1 shows a typical impedance spectrum of

a ME, obtained prior to any polarization, in form of a Nyquist-plot. Fitting of the

impedance data was done with Zview software (Vers. 3.4f , Scribner Associates).

There are two distinct features in the plot, a high-frequency arc and a low-frequency

capacitive increase. The former corresponds to the charge transport in the probed

sample, which is described by a resistive element (RLLZO) and a parallel constant-

phase element CPELLZO in the equivalent circuit. The low-frequency increase is

caused by the ionically blocking gold microelectrode, described by CPEAu in the

equivalent circuit. A tting of the equivalent circuit to the impedance spectra reveals

the values of RLLZO. Since single crystals were investigated, no charge transfer along

26

Figure 3.1: EIS data (circles) of a ME, measured prior to any experiments at roomtemperature. Fitting results are shown as a dashed line. The used equivalent isshown in the inset.

grain boundaries is possible. Thus, only the local bulk ionic conductivity of the

probed sample contributes to the high-frequency arc. From the resistance of the

high-frequency arc, the local bulk ionic conductivity σLi can be calculated:

σLi =1

2 d RLLZO

(3.1)

Impedance data measured from various microelectrodes prior to polarization show

variations in ionic conductivity (s. Fig. 3.2). After the measurements, the sample

was subjected to thermal treatment (58 hours at 300 C), similar to the polarization

experiments, but without application of any bias. Then, impedance spectra were

measured from the same microelectrodes, again (s. Fig. 3.2). The sample already

shows a scattering of conductivity (2 · 10−5 - 9 · 10−5 S/cm), right after preparation.

An increase of conductivity (up to 2 · 10−4 S/cm) is apparent for most ME, whereas

other ME show no signicant changes in conductivity.

After the application of DC bias (3 V) for 15 hours at elevated temperature (400

C) and subsequent cooling to room temperature (under bias), EIS measurements

were carried out, to investigate the impact of eld stress on the conductivity. After

each polarization experiment, the microelectrodes between the corresponding stripe

27

Figure 3.2: Ionic conductivities, measured via EIS at various microelectrodes. Scat-tering of conductivity is apparent directly after sample preparation (black squares).After the thermal treatment (red diamonds), an increase in conductivity can beobserved.

electrodes were used for ME-EIS (s. Sec. 2.2). Fig. 3.3 shows data from ME-

EIS prior to and after polarization via two opposing stripe electrodes. Close to the

cathode, a slight increase in conductivity is apparent, whereas close to the anode a

decrease can be observed. This general trend was present also on further polarization

experiments, however, the magnitude of the conductivity changes was rather low in

most cases.

Also, the current, owing through the sample during polarization was measured

(s. Fig. 3.4). A rapidly decreasing current is found which reaches 1.3 µA after 5

minutes. After about 60 minutes, the current has dropped to about 1.2 µA and

then remained almost constant for many hours. The relatively high current at

the beginning at the experiment is attributed to a motion of Li-ions. In order to

compensate the gradient in electrochemical potential, induced by the electric eld,

the Li-ions move from the cathodic side to the anodic side of the sample. Assuming

the gold electrodes are blocking the Li-ions, the Li-ion current cannot be retained

for a longer time but drops. After the ionic current has largely diminished, a low

current of about 1µA is still constantly owing. This so-called steady state current is

attributed to the small, but non-zero electronic conductivity of the sample [53]. Since

the electrons can readily pass through the gold electrodes, the electronic current is

28

Figure 3.3: Ionic conductivities, measured via EIS at microelectrodes between thepolarization electrodes, before (black squares) and after polarization (blue circles).Again, scattering of the conductivity is apparent before the polarization. Afterapplication of eld stress via the stripe electrodes (left and right), conductivitieshave changed.

constant over time. In such a case the electronic conductivity can be estimated from

the voltage (U=3 V), current (I=1 µA), length (l=4 mm) and cross section (A=0.63

mm2) of the conductor:

σ =1

R

l

A=I

U

l

A(3.2)

From equation (3.2) an electronic conductivity of about 2·10−5 S/cm is obtained

at 400 C set temperature. This conductivity value, however, is too high to be

attributed to the poor electronic conductivity of LLZO.

Therefore we assume that a continuous decomposition of the sample leads to a

certain Li-ion current which contributes to the 1 µA measured. Existence of some

ionic current is supported by several features observed during and after eld stress.

After polarization, optical changes of the polarization electrodes were apparent (s.

Fig. 3.5), indicating electrochemical reactions at the electrodes. The anode shows a

rough surface after polarization. The roughness appears to be caused by gas bubbles,

possibly arising from the oxidation of oxide ions, which lift o the gold electrode

from the sample surface. On top of the cathode, a dark solid was deposited during

polarization, growing from the edge of the electrode over the entire surface of the

sample. Only a small area of the underlying metal electrode is still visible in the

29

Figure 3.4: Current measured during polarization (3 V) at 400 C. Starting at arelatively high value, the current drops drastically within minutes to a value ofabout 1 µA. a) shows the entire experiment, whereas (b) shows a magnied view ofthe rst 30 minutes.

microscope image. The dark solid most likely consists of Li2CO3, possibly mixed

with Li2O or LiOH. These are expected for cathodic reduction of oxygen in a Li

-ion (i presence of CO2 and H2O, see below). Electrochemical reactions during

polarization could also cause the irregular increase of current in the time interval

from about 6 hours to 11 hours (s. Fig. 3.4). After these irregularities, a slight

decrease of current is observed. This drop can be caused by the gradual lift-o of

the anode, due to the formation of oxygen gas, which leads to a smaller contact

area, thus increasing the resistance of the sample.

The deposition of Li salts and gas formation at the cathode and anode, respec-

tively, can be explained by the following, suggested mechanism (s. Fig. 3.6): Upon

eld stress, oxygen ions are oxidized at the anode, leaving oxygen vacancies inside

the crystal:

Anodic half reaction : OXO →

1

2O2 + V ∗∗O + 2e′

The gaseous oxygen formed upon oxidation is trapped underneath the metal an-

ode, lifting it o. The positive charge of the oxygen vacancies is compensated by

formation of negatively charged Li vacancies (at a ratio of 1:2). The Li-ions are

transported towards the cathode leaving negatively charged vacancies at the an-

ode. These excess Li-ions react with excess electrons, provided by the cathode, and

gaseous species from the ambient atmosphere (i.e. H2O, O2 and CO2), thus reducing

30

the gases and depositing LiOH, Li2O or Li2CO3, e.g.:

Cathodic half reaction : 2Li+ + 2e′ + CO2 +1

2O2 → Li2CO3

Since electrons, Li-ions and the ambient atmosphere are required for this reaction, it

can only take place at the triple-phase-boundary, i.e. the edge of the metal electrode.

Overall, the process leaves Li- and O depleted LLZO and oxygen bubbles underneath

the anode, while the cathode is covered with Li-containing salts.

Figure 3.5: Optical microscopy images of the polarization cathode (left) and anode(right) after a bias of 3 V was applied for 15 hours at 400 C. Prior to polarization,these electrodes had a smooth and reective golden surface. Magnication: 10x,scale bars: 200 µm.

3.1.2 LIBS

After the application of eld stress, LIBS measurements were carried out. The laser

was scanned across the sample in lines between the stripe electrodes (s. Sec. 2.2).

Each line was scanned twice, since the gold electrodes were still present on top of

the sample, possibly aecting the measurement of the elemental composition of the

LLZTO. The rst laser scan of each line completely ablated the gold covering the

sample. The second line-scan is not aected by the electrodes, thus the signals can

be attributed to the LLZTO sample.

The intensity of the emitted, characteristic radiation cannot be directly used for

analysis of the elemental composition. Possible variations in laser-intensity or other

inuence on the signal, can simulate variations in concentration of the elements.

31

Figure 3.6: Schematic illustration of the suggested mechanism. Oxidation of oxygenat the anode leads to oxygen gas evolution. Oxygen gas formed underneath themetal anode leads to bubbles, partially lifting o the anode. At the cathode, gaseousspecies, such as O2, are reduced, yielding Li2CO3. Overall, Li-ions are transportedthrough the sample, leaving vacant Li-sites and O-vacancies at the anodic side.Li-containing salts are formed at the triple phase boundary of the cathode.

To circumvent this problem, a reference signal is required. The element, emitting

the reference signal should have a uniform concentration throughout the sample.

Thus, any variations in signal intensity of this element can be attributed to external

eects, such as laser-intensity, laser-focusing etc. Both La and Zr are expected to be

uniformly concentrated throughout the sample. The relative La/Zr signal (s. Fig.

3.7) is scattered, but shows no signicant changes over the scanned lines. Since it is

highly improbable, that the concentrations of La and Zr are related in a way such

as to yield a uniformly scattered, relative signal, the results presented in Figure 3.7

prove, that La and Zr concentrations can be considered constant. Therefore, a Zr

signal (λ=327 nm) was chosen as a reference signal for all LIBS measurements.

The relative signals of La and O showed no signicant changes over the length

of the line-scans. However, primarily Li was expected to move through the sample,

as its ionic conductivity is highest. The relative Li intensity (λ=610 nm) of two

line-scans are shown in Fig. 3.8.

Between the two stripe electrodes, no distinct gradient in Li content is apparent.

However, depletion and enrichment of Li is apparent underneath the cathode and

anode, respectively. The shift of Li from the anode to the cathode is in agreement

with the suggested polarization mechanism and the polarization current data. The

32

Figure 3.7: La-267 signal, referenced to the Zr-327 signal of a typical LIBS line-scan.Prior to this scan, a preliminary laser scan was performed to ablate the electrodematerial. The positions of the polarization electrodes are indicated by golden stripes.

increased Li content at the cathodic side of the sample indicates, that some can

Li-ions accumulate at the cathode that do not react with the gas phase, thus locally

reducing the LLZTO.

3.2 Results from Polarization of Microelectrodes

The current, owing during polarization of microelectrodes, was measured under

various conditions (gas atmosphere, temperature, voltage), to get a deeper under-

standing of the electrochemical processes occurring during polarization of LLZTO.

Further, the impact of eld stress on the elemental composition was investigated by

LIBS measurements.

3.2.1 Polarization Current Measurements

Polarization of two neighboring MEs gives rise to a current, similar to the current

observed during polarization of large electrodes (cf. Sec. 3.1.1). A typical current

prole, obtained during repeated polarization under Ar-atmosphere at 400 C, is

shown in Fig. 3.9. The sample was polarized repeatedly with increasing bias. Be-

33

Figure 3.8: Li-610 signal, referenced to the Zr-327 signal of two LIBS line-scans.Each line was scanned twice to separate signals of the LLZO from any signals of thegold electrodes and possible precipitates on top. (1) shows the rst line scan overthe sample. After the rst scan all gold electrodes were ablated. The second scan(2) is therefore not aected by electrodes and possible precipitates.

tween each polarization period (10 minutes), the voltage was set to 0 V for 1 minute.

All curves in Fig. 3.9 show two distinct features:

i. A relatively high current (exceeding 1 nA) at the beginning of the experi-

ment, rapidly decreasing to values < 0.1 nA within seconds (depending on the

voltage). Apparently, the positive current approaches a constant, small value,

which is not reached within the time of the experiment. The magnitude of

the constant, so-called steady-state current (SSC), is directly related to the

applied voltage. As in section 3.1.1, the SSC is attributed to a small, but

non-zero electronic conductivity of the sample, since no other charge carrier is

expected ow through the sample steadily.

ii. When the voltage is set to 0 V, a negative current arises. The magnitude of

this negative current is decreasing in a similar manner as the positive current

during polarization. However, the current must approach 0 A, since no voltage

is applied to the sample. Again, the time of the experiments was not long

enough to observe the current reaching 0 A.

The results shown in Fig. 3.9 reveal that application of bias to neighboring MEs (s.

34

Figure 3.9: Typical current data, obtained during polarization of two neighboringmicroelectrodes. 100 (black), 200 (red) and 300 mV (blue) were applied for 10minutes in air at 400 C. Between each polarization period, the bias was set to 0 Vfor 1 minute. The voltage-over-time prole is shown in the inset.

Fig. 2.4 (a)), gives rise to a reversible charging in the sample. Removing the bias

leads to a reverse current, indicating, that the charge carriers, shifted during polar-

ization, migrate back towards their initial positions. Most probably the following

electrochemical processes take place: The bias causes a dierence of the electro-

chemical potential of the Li-ions in LLZO. The gradient of electrochemical potential

serves as the driving force for Li-ion migration. The shift of Li-ions compensates the

gradient in electrochemical potential, approaching equilibrium (i.e. constant electro-

chemical potential). When the bias is removed, the gradient in Li-ion concentration

is no longer compensated for, thus the gradient leads to a reverse migration, ap-

proaching the original Li-ion distribution.

Li-ions are the ionic charge carriers with the highest mobility in LLZTO. There-

fore, the reversible polarization current can be attributed to the reversible migration

of Li-ions. Further, migration of Li-ions due to eld stress was shown in experi-

35

ments described above (s. Sec. 3.1). Since the experiments were conducted under

Ar atmosphere, reactions with the gas phase are highly improbable. Indeed, no de-

composition reactions (cf. Sec. 3.1.1) were observed during any of the polarization

experiments (voltages up to 3 V). Thus, in the true steady state situation (not yet

reached here) only electrons are expected to ow. This is also in accordance with

the much lower currents found in these experiments compared to those in Sec. 3.1.

3.2.2 Inuence of the Atmosphere

The formation of Li2CO3 shows, that reactions of LLZTO with the gas phase can

occur when eld stress is applied via blocking electrodes (s. Sec. 3.1.1). Therefore,

the chemical composition of the gas phase aects the polarization in some way. To

get a better understanding of the inuence of the gas phase on the polarization of

LLZTO, polarization experiments were done under Ar and air atmosphere at 400

C. Both under Ar and air atmosphere, the same voltage prole was applied to a

ME, whereas the large backing electrode served as the counter electrode (s. Fig. 2.4

(b)). The resulting current data are displayed in Fig. 3.10.

Fig. 3.10 shows currents, which are typical for polarization of MEs (cf. Fig. 3.9).

The red and blue curves dier slightly, indicating that the atmosphere has an eect

on the polarization current. During polarization under air, the current is generally

lower in value, and decreasing faster than the current measured in Ar atmosphere.

Further, the current measured under air is still decreasing after 10 minutes of po-

larization, whereas the current under Ar even shows a slight increase after about 7

minutes. The current in Ar after 10 minutes (about 50 pA at 1 V) is in reasonable

agreement with that for two microelectrodes (Fig. 3.9, about 25 pA at 300 mV).

Assuming that steady state current is electronic, we may use the corresponding re-

sistance (R = 1 V50 pA

= 20 GΩ) for calculations on electronic conductivity at 400 C.

From the spreading resistance formula (Eq. 3.1) we get 2.5·10−9 S/cm.. The neg-

ative currents, observed when turning o the applied bias, show an opposite trend.

The negative current under air is generally higher in value, than the current under

Ar. The faster decrease of current observed during polarization under air might

be caused by chemical changes of the sample. Although no distinct evidence for a

decomposition of the sample (i.e. bubble formation or precipitation) was apparent

36

Figure 3.10: Polarization current, resulting from experiments carried out underAr (blue) and air atmosphere (red) at 400 C. The solid lines show results frompolarization with 100 mV, the dashed lines show results from polarization with 1000mV.

at the electrodes, of the LLZO sample could occur during polarization under air.

Since surface degradation of LLZO in air is a known issue [57], this is a possible

explanation for the steady decrease in conductivity. To avoid undesired degradation

eects, most proceeding experiments were carried out under Ar atmosphere.

The current, indicated by the solid red line in Fig. 3.10, becomes negative after

about one minute. Since a positive voltage of 100 mV is applied to the sample,

a negative current, owing against the electric eld is not consistent with basic

physical concepts. Therefore, the slightly negative current values (> -10 pA) are

attributed to measurement errors, possibly arising from thermoelectric phenomena

in the measurement setup.

3.2.3 Inuence of the Temperature

Both ionic and electronic conductivities are dependent on temperature. Generally,

ionic conductivity in solids is enhanced at higher temperature, because the activation

energy for ion migration can be overcome more easily. LLZTO is a wide-band-gap

37

semiconductor, therefore thermal excitation of electrons might also lead to enhanced

electronic conductivity at higher temperature. Therefore, a generally higher current

is expected, when polarization is carried out at higher temperature. Fig. 3.11 shows

current measurements from similar polarization experiments, conducted at room

temperature (25 C) and 400 C under Ar atmosphere.

Figure 3.11: The black lines, solid (100 mV), dashed (200 mV) and dotted (300mV), indicate the measured current during polarization at room temperature. Thecorresponding results obtained at 400 C are indicated by red lines. A generalincrease of current at 400 C is apparent.

The comparison in Fig. 3.11 shows, that the measurements at 400 C are sub-

jected to a distinct, negative oset (s. Sec. 3.2.2), whereas the data obtained at

room temperature show barely any current-oset. This temperature dependence

suggests, that thermoelectric eects contribute to the current-oset. The basic fea-

tures of the current curves are the same, independent of temperature. However, the

red curves do not approach a steady-state during the course of the experiments (10

minutes), whereas the black curves appear constant after about 5 minutes of polar-

ization. This indicates, that a larger number of charge carriers (possibly Li-ions or

O-ions) is shifted reversibly during polarization at higher temperature. Assuming

that the steady state current (not yet reached) corresponds to electron transport

and the additional decaying current is due to Li-ions, we can roughly estimate the

amount of Li transported during the experiment.

38

We may assume a steady state current of 10 pA for 300 mV and subtract this

from the measured current. Then we integrate∫ t1t0I(t)·dt. This gives 1.8·10−8 C and

thus about 1.1·1011 Li-ions. Assuming the density of LLZTO is about 5.3 g/cm3 [70]

this would correspond to a depletion of Li in LLZTO by about 3.7·10−3 per formula

unit in a region of 1 µm beneath the microelectrode (i.e. in ca. 0.8 ·10−8 cm3).

Hence, only a slight stoichiometry gradient builds up.

3.2.4 Results from a Broader Voltage Range

Polarization experiments with MEs were carried out over broader voltage range

(from -2 V up to 2 V) under Ar atmosphere. The voltage was increased/decreased

incrementally, holding the voltage at each step for a given time period (typically

1 hour). After the maximum/minimum voltage of ±2 V was reached, the voltage

was decreased/increased again, retracing the steps. Typical voltage increments were

between 20 mV and 100 mV.

Results from polarization of a ME contacted versus the large backing electrode,

are displayed in Fig. 3.12. Voltage was increased from 0 V to 2 V (100 mV incre-

ments) and decreased to 0 V (Ar atmosphere, 300 C). Currents are muhc lower

than at 400 C. Part (a) shows selected current curves of the increasing part of the

experiment, whereas the decreasing half is shown in (b). Current curves, measured

during polarization with 1100 mV (dark green) and 1200 mV (light green) have

atypical shape. A high current at the beginning decreases rapidly. The curves do

not, however, approach a constant value within the course of polarization time (1

hour). Strangely, the current starts increasing slightly after few minutes.

When the voltage is increased to 1300 mV, a distinct change of the current curve

is apparent: The initially high current decreased over several minutes. Further, the

current drops down below the current from 1200 mV polarization after about 45

minutes and keeps decreasing afterwards. Subsequent current curves (1400 mV -

2000 mV) have the same curve shape again. However, values of the current curves

are declining, despite the increasing voltage.

When the voltage is decreased (b) again, the curves seem to approach a constant

value from below. When voltage is turned o after polarization (c.f. Fig. 3.9), a

negative current, approaching zero, arises because charge carriers migrate back to-

39

wards their initial positions at 0 V. Similarly, when the voltage is decreased during

polarization, a reverse current is induced. However, when the voltage is still applied,

this revers current is overlain by the constant steady-state-current, thus approaching

a positive steady state. In contrast to Fig. 3.12 (a), the curves in Fig. 3.12 (b) seem

almost constant after 60 minutes of polarization. Further, the approached current

values are related to the applied voltage directly. Overall, the results suggest, that

currents, obtained when the voltage is decreased, are more suitable for the deter-

mination of actual steady-state current values. Still, the rather unusual behavior

might indicate that most probably a chemical decomposition process rather than the

electronic conductivity is measured at such high voltages. This is also in agreements

with the measurements shown below.

Polarization of MEs with negative voltages gives rise to dierent current curves,

as described above. Selected results from polarization with up to -2 V are displayed

in Fig. 3.13. The voltage prole was analog to the positive voltage polarizations,

described above. All other experimental parameters were identical to the positive

voltage polarization experiments.

Disregarding the negative signs of current and voltage, the results from nega-

tive polarization of MEs are similar to the results from positive polarizations: The

negative current, arising from negative polarization, is relatively high in value at

the beginning, dropping rapidly to smaller negative values, approaching a steady

state. Increasing the voltage to higher, negative values, results in higher negative

currents. When the value of the negative voltage is decreased again, some charge

carriers migrate in the reverse direction, giving rise to a positive, decreasing current.

The comparison between corresponding dashed and solid lines in Fig. 3.13 indicates

how closely the curves have approached the steady state current.

3.2.5 U-I Characteristic

Numerous polarization experiments in the voltage range from -2 up to 2 V were

carried out in a similar way as described in the previous sections. The voltage

increments and the time of each polarization step was varied, but the basic exper-

iment was unchanged. To investigate the nature of the steady-state current, which

is approached during a polarization experiment, steady-state currents of many ex-

40

Figure 3.12: Current, measured during polarization of a ME at 300 C. The voltagewas increased (a) and decreased (b) incrementally in 100 mV steps, going from 0 Vto 2 V and back to 0 V. For the sake of readability, only selected data is shown. Theblack line, indicating the current measured at 2 V, is displayed in both graphics andcan be used as a reference when comparing the two graphics.

41

Figure 3.13: Current, measured during polarization of a ME at 300 C. The voltagewas decreased (solid lines) and increased (dashed lines) incrementally in 100 mVsteps, going from 0 V to -2 V and back to 0 V. For the sake of readability, onlyselected data is shown. The solid and the corresponding dashed line of each colorapproach the same steady-state-current.

periments were compared (s. Fig. 3.14 & Fig. 3.17), revealing the relation between

polarization voltage and steady-state current. Experiments were carried out under

Ar atmosphere at 300 C. Fig. 3.14 shows the voltage range between -0.5 V and 0.5

V, measured on 4 microelectrodes. with current values in the scale of pA. Exem-

plary current-time curves for two electrodes are shown in Fig. 3.15. For plotting the

steady state I-U-curve of Fig. 3.14, the mean current values of the last 10 minutes of

each voltage step was used. Despite dierences we can state that for all microelec-

trodes the currents are in the few pA range. It is also obvious that there is a voltage

intercept present for extrapolating the curves to zero current. This oset between

50 and 250 mV might partly be due to thermovoltages but might also include a

contribution from any built-in chemical potential dierence of Li.

It is also obvious that the curves are non-linear. Assuming electronic conductivity

as the reason of the current ow, this would be not surprising. Supposed the counter

electrode were hardly polarized (due to its large size) the voltage in steady state

42

translates to a chemical potential dierence of Li beneath the microelectrode. This

should aect the local electron concentration and thus the electronic conductivity.

Figure 3.14: Steady-state current plotted against the polarization voltage, rangingfrom - 0.5 V to 0.5 V at 300 C. Each color represents a single MEs, which wascontacted for polarization. The dierent symbols indicate the dierent measurementseries.

The black boxes show hysteresis, indicating, that the stead-state values obtained

from the increasing part of the series (upper branch) are higher than the ones ob-

tained during decrease of voltage (lower branch). Thus, the steady-state was not

approached closely enough to obtain accurate steady-state current data. However,

since the hysteresis is small (< 1 pA), the data can be used as estimates. The

red diamonds (bottom half lled) show almost no hysteresis at all, suggesting that

steady-state conditions were approached closely during the experiments. Similarly,

the green circles in the negative voltage range show only small hysteresis. Extrap-

olation between the negative and positive voltage data, suggest a current oset of

about -2 pA, which can be attributed to thermoelectric eects within the measure-

ment setup (s. Sec. 3.2.2 and Sec. 3.2.3). The red diamonds (top half lled) are

43

Figure 3.15: Current, measured during polarization of two individual MEs (a) and(b). Steady state values of each voltage step were estimated by taking the averagecurrent of the last 10 minutes of each step. The steady state values of (a) correspondto the red diamonds (bottom half lled) in Fig. 3.14 whereas (b) corresponds to thegreen circles.

from a measurement series which does not include voltage steps going back towards

0 V. Thus, the reliability of the steady-state currents cannot be deduced from any

possible hysteresis. However, the voltage increments as well as the duration of each

voltage step was the same as in the positive voltage range, which gave almost no hys-

teresis. Therefore the red diamonds in the negative voltage range can be expected

to be close to the true steady-state values.

The blue symbols deviate most from a linear I-U relation. The current data

obtained during the positive voltage measurement series is shown in Fig. 3.16.

Starting at 100 mV, the voltage was increased in 20 mV steps. The rst polarization

period was longer (1 hour) to ensure steady-state conditions, since the rst voltage

step was the highest (from 0 V to 100 mV). The subsequent periods were shorter

(10 minutes) since the increments were only 20 mV. From Fig. 3.16, however,

it is apparent, that the 10 minute periods are insucient to obtain steady-state

conditions. Especially at higher voltages, the long time, required to achieve steady-

state conditions, makes measurements of steady-state currents very time consuming.

For obtaining an estimate of the electronic conductivity we used extrapolated slopes

analyzed for U=0 V. We get (1-2 pA)/100 mV and thus a resistance of 50...100 GΩ.

This corresponds to 0.5 - 1·10−9 S/cm. This is slightly lower than the estimate of 400

C (2.5·10−9 S/cm) which might reect the true thermal activation but could also

44

be caused by the problems caused by only apparent steady states, non-linearities

and voltage osets.

Fig. 3.17 shows selected U-I data from a broader voltage range (-2 V to 2 V).

Since higher voltages were applied to the sample, the polarization time periods were

increased (typically 1 hour) to ensure steady-state conditions. In the positive voltage

part, the linear relation between voltage and steady-state current (cf. Fig. 3.14), is

continued up to 2 V. In the negative voltage part, the U-I relation at lower voltages

(s. Fig. 3.14) is continued similarly. However, at about -800 mV, a drastic in-

crease in current is apparent. The current data up to -1 V (red circles) show almost

no hysteresis, thus indicating that steady-state conditions were approached closely

during each polarization period. Therefore, the strong increase of current at about

-800 mV cannot be attributed to insucient polarization time. Rather, it suggests a

change in conduction mechanism at about - 800 mV. At voltages below -1 V (black

boxes), high negative currents are obtained. However, the currents show irregular

uctuations and do not approach steady-state within the polarization time (1 hour).

Steady-state currents could therefore only be estimated. The mean value of the last

10 minutes of each polarization period were used as rough estimates. After the

minimum voltage of -2 V was reached, the voltage was increased again (approach-

ing 0 V). The current data shows strong deviations from the data obtained when

the voltage was decreased towards -2 V (i.e. hysteresis). This strongly suggests

irreversible chemical changes in the sample. Hence precesses may have started that

correspond to those found for macroscopic electrodes in Sec. 3.1. The still dierent

current densities for the microelectrodes (1.3·10−6 A/cm2) and macroscopic elec-

trodes (1.6·10−4 A/cm2) may be caused by either still higher voltage or the higher

temperature.

3.2.6 LIBS

The impact of ME-polarization on the local composition of the sample was investi-

gated via LIBS. 8 dierent MEs were polarized (voltages between -2 V and 2 V) for

1.5 hours under Ar atmosphere at room temperature. After polarization, the rela-

tive, local chemical composition was determined via LIBS. Comparison between the

8 MEs and 4 MEs that were not polarized, showed no signicant dierence in chem-

45

Figure 3.16: Current measurements corresponding to the blue triangles (bottom halflled) in Fig. 3.14. The distinct negative slope at the end of some polarization stepsindicates, that steady-state was not approached closely enough to obtain reliablesteady-state currents. The voltage increments were 20 mV. Since the rst polariza-tion had a greater step (from 0 V to 100 mV), the polarization time was 1 hour,whereas the subsequent polarizations were only 10 minutes each.

ical composition between any of the MEs. However, since the LIBS measurements

were not done directly after polarization of each ME, possible eects of polarization

could have been reversed before the LIBS measurements.

46

Figure 3.17: Steady-state currents between -2 V and 2 V (a) and a magnied view ofthe range from - 1V to 2 V (b) at 300 C. The black boxes indicated rough estimatesof steady-state current values.

47

Part II

Sputtering and Testing of LIB

Electrode Materials

48

Chapter 4

Construction of a Sputtering Device

A spherical vacuum chamber (240 mm diameter), equipped with several anges of

dierent diameters, serves as the basis of the device (s. Fig. 4.1 and Fig. 4.2). The

chamber is mounted on an aluminum frame. A radio-frequency (RF) sputtering

gun (Gencoa 3G Circular Magnetron) is passed through the top ange alongside

with a shutter. The gun is connected to a cooling water cycle, the RF-power source

(Seren, R 301 RF Power Supply) and gas lines (argon and oxygen). The gases are

let into the chamber directly at the sputtering target. The gas ow of argon and

oxygen can be controlled independently by mass-ow-controllers (Tylan, FC 260 for

Ar and FC 2900 M for oxygen). A turbo molecular pump (Leybold Turbovac 250 iX)

combined with an auxiliary scroll pump (Agilent Technologies, SH-110 Dry Scroll

Vacuum Pump) is connected to the chamber via one of the side-anges. Two pressure

sensors (Leybold, Ceravac CTR 100 N and Penningvac PTR90 N) are installed via

another ange to monitor the chamber pressure. Wiring for heating, thermocouples

and possible in-situ measurements are passed through another ange.

The sample stage holds an alumina sample holder which can also serve as a heat-

ing unit. The sample holder is equipped with a type S thermocouple for temperature

control. So far, the highest tmeperature used in operation was 640 C. Sputtering

crystalline thin lms often requires elevated deposition temperatures to aid crystal-

lization. At room temperature, especially ceramic materials tend to form thin lms

with poor crystallinity. The target to sample distance of the device is variable. For

the course of the experiments presented in this work, the distance between target

and sample was set to 6 cm.

49

Figure 4.1: Overview of the sputtering device with labels marking the most essentialparts.

4.1 Preparation and Characterization of Thin Films

First tests of the sputtering device were done with a copper target. Sputtering of

metals is relatively simple, since no reactive atmosphere is required. Further, the

targets can be simply prepared from sheet metal. Since metals are ductile, there is

no danger of metal targets cracking or breaking due to thermal expansion during

sputtering.

Copper was RF-sputtered onto polished YSZ substrates in 3·10−2 mbar argon

atmosphere at room temperature. The substrates were partially covered with a

shadow mask to obtain a relief after deposition. Sputtering power was between 50

W and 110 W and the deposition time was between 5 and 15 minutes. The thickness

of the thin lms was measured with a prolometer (Bruker, Dektak XT) by scanning

over the relief surface.

Targets for deposition of ceramic materials were purchased from Beijing Loyaltar-

gets Technology Co. Active materials for LIBs, i.e. the cathode material LiMn2O4

(99.9 % purity) (LMO) and the anode material Li4Ti5O12 (99.9 % purity) (LTO)

were deposited onto YSZ substrates similarly to the copper thin lms, as described

above. However, the sputtering atmosphere was dierent, since oxide materials re-

50

Table 4.1: Deposition parameters of the samples prepared for structural investiga-tions.

CompoundTime(min)

Power(W)

Thickness(nm)

Pressure(mbar)

pO2

(mbar)Substrate

Li4Ti5O12 200 80 400 1.2 10−2 8.8 10−3 Si [700]LiMn2O4 170 50 90 3.6 10−3 8.3 10−4 Quartz glass

quire a reactive atmosphere. In pure Ar, oxides would be partially reduced during

deposition, resulting in undened stoichiometry and undesired materials properties.

Deposition was done at room temperature rst, since the heating unit had not yet

been installed at that time. Thickness of the lms was determined after deposition

as described earlier.

Further samples were prepared in order to investigate structural and electro-

chemical properties of the oxide thin lms. LTO and LMO lms were deposited

onto silicon wafers and quartz glass substrates at room temperature for structural

investigations (see Table 4.1 for deposition parameters). Crystal structures of the

samples were determined by means of powder X-ray diraction (XRD) in Bragg-

Bretano geometry (PANalytical, X'Pert Pro, Cu K-α, 45 kV/40 mA). Quartz glass is

amorphous and the used Si-wafers are cut in [700]-orientation, thus giving no Bragg-

reections in the measured angle range. After deposition and XRD measurements,

the lms samples were subjected to a thermal treatment (annealing) to improve

crystallinity. The samples were heated in air to 500-700 C for 8-12 hours. After

annealing, the structures of the samples were again measured via XRD.

All electrochemical procedures described below were done using a Novocontrol Al-

pha Analyzer-Pot/Gal 30V-2A. Samples for electrochemical characterization of the

sputtered thin lms were prepared using LLZTO single crystals (s. Sec. 2.1) and

polycrystalline LLZTO pellets of the same nominal composition. Sputtering param-

eters were similar to the parameters described in Table 4.1. Thin lms of LMO (60

nm) were sputtered symmetrically onto a polycrystalline pellet and subsequently a

gold layer (50 nm) was sputtered on top of each side. Prior to and after an annealing

step (3 h at 700 C) for crystallization, the sample was characterized by means of

EIS. Further, the sample was electrochemically cycled repeatedly.

51

Table 4.2: Deposition parameters of the asymmetric sample for electrochemicalcharacterization.

CompoundTime(min)

Power(W)

Temperature(C)

Thickness(nm)

Pressure(mbar)

pO2(mbar)

Li4Ti5O12 20 80 600 31 1.2 10−3 2.0 10−3

LiMn2O4 100 50 400 20 2.0 10−3 3.0 10−3

An asymmetrical sample was prepared using a LLZTO single crystal. LMO was

deposited onto one side and LTO onto the other side by reactive RF-sputtering

at elevated temperature. Deposition details are shown in Table 4.2. Thin gold

layers (20 nm) were sputtered on top of each side as current collectors. The sample

was characterized by EIS and DC-cycling. Film thickness was determined at a

prolometer. Comparison with the thicknesses of layers, which were deposited at

room temperature (s. Table 4.1), shows, that the deposition rates are lower at

elevated temperatures. This can be attributed to a higher density of the lms, since

crystallization and densication is promoted at elevated temperatures.

52

Figure 4.2: From top left to bottom right: Overview; side view, revealing pressuresensors and turbo molecular pump; view inside the chamber during deposition ofLTO, revealing the wiring of the sample stage; sample during deposition.

53

Chapter 5

Results and Discussion

5.1 Film Thickness and Deposition Rates

First deposition rates were determined for sputtering copper at room temperature

in pure Ar atmosphere. The measurement results from the prolometer are shown

in Fig. 5.1. Naturally, longer deposition times and higher sputtering power results

in thicker lms. Roughly, deposition rates can be obtained by extrapolation of the

blue circles and black boxes. However, the lines passing through either one of the

data sets do not intersect the axes at 0/0 but rather intersect the time-axes at about

1 or 8 minutes. Since no shutter had been installed at the time of deposition, it was

not possible to sputter the target with closed shutter prior to deposition. Sputtering

with closed shutter for several minutes is a common procedure in many sputtering

processes to ensure the sputtering target is cleaned of any surface impurities that can

aect the deposition. Since this was not possible at the time of the rst sputtering

experiments, the oset in Fig. 5.1 can be attributed to the missing sputtering

procedure with closed shutter. It is apparent, that the oset value is clearly higher

at lower sputtering power. This results suggests, that detrimental surface impurities

of the copper target are removed faster at higher sputtering power.

Sputtering oxide thin lms generally results in lower sputtering and deposition

rates. Further, oxide materials require a reactive atmosphere for sputtering to cir-

cumvent possible reduction reactions of the sputtered material. Results from pro-

lometer measurements of the oxide lms are shown in Fig. 5.2. LTO deposition

rates at 80 W sputtering power are clearly higher than deposition rates of LMO at

54

Figure 5.1: Film thickness of various Cu thin lms deposited onto polished YSZsingle crystals. Sputtering power as well as the deposition time were varied.

Table 5.1: Deposition rates of Cu and the two oxide materials LMO and LTO.

CompoundPower(W)

Deposition Rate(nm min−1)

Pressure(mbar)

Cu 50 27 3 10−2

90 39 3 10−2

100 41 3 10−2

110 43 3 10−2

LTO 80 50 2 10−3

LMO 50 20 3 10−3

50 W. The blue circles, indicating the LTO lm thickness, show a linear dependence

on the deposition time, as expected.

Deposition rates for Cu, LTO and LMO were estimated from the lm thicknesses

and corresponding deposition times, assuming a linear relationship. The estimated

deposition rates are shown in Table 5.1. The deposition rate of LTO is unexpectedly

high, even higher than all measured deposition rates of Cu. This could be explained

by a lower total pressure during deposition, since higher pressure during sputtering

leads to more collisions between the sputtered material and the gas atmosphere,

thus reducing the deposition eciency. Further, the deposition rates presented in

55

Figure 5.2: Film thickness of various LMO and LTO thin lms deposited onto aSi wafer and quartz glass substrate respectively. Sputtering power as well as thedeposition time were varied. All lms were deposited at room temperature.

Table 5.1 are rough estimates, relying on 3 or less data points for each rate. More

data is required to obtain reliable deposition rates.

5.2 Crystal Structure

The crystal structure of the oxide thin lms was investigated by means of XRD.

The measurements were done in Bragg-Bretano geometry, scanning from 15 to 70

2θ at a step size of 0.020. Each step was scanned for 160 seconds. Since the

thickness of characterized lms is in the nanometer range, long scanning times were

necessary to obtain enough signal from the thin lms. As discussed earlier, thin

lms deposited at room temperature tend to have amorphous or poorly crystalline

structures. Therefore, thin lms are often subjected to thermal treatments after

deposition. X-ray diractograms of the LMO and LTO lms are presented in Fig.

5.3 and Fig. 5.4, respectively. Each gure includes diraction data from the lms as-

deposited and post-annealing. From both gures, a distinct increase of crystallinity

due to the annealing is apparent.

The as-deposited data, presented in Figure 5.3 (black line) shows no dened

56

Figure 5.3: X-ray diractograms of the LMO thin lm, deposited on quartz glass.The black line shows the data from the as-deposited scan, whereas the red lineindicates the data obtain post-annealing. A distinct increase of crystallinity dueto the thermal treatment is apparent. Due to crystallization of the quartz glasssubstrate, reexes of crystalline quartz (*) are visible.

reexes and a large hump in the region at about 20, typical for amorphous samples.

After the annealing (red line), several reexes are apparent in the diractogram.

The reexes marked with an asterisk correspond to the crystal structure of quartz,

whereas the reexes marked with the letter "s" correspond to the crystal structure

of the spinel phase of LiMn2O4. The quartz reexes can be explained by partial

crystallization of the substrate during annealing. Recrystallization of glasses at

elevated temperatures is a well-known phenomenon. Further, oblique inclusions were

apparent inside the quartz glass substrate after annealing. The other reexes can

be clearly attributed to the desired spinel phase of LiMn2O4. However, despite the

long scanning time, the reex intensities are low compared to the amorphous hump,

57

Figure 5.4: Diractograms of the LTO sample, as-deposited (black) and post-annealing(red). The sample shows no sharp reexes as-deposited but only one broadreex at about 43 . Post-annealing, more reexes are apparent and the reex atabout 43 is sharper than before. However, the few reexes could not be assignedto a single phase denitely.

therefore not all reexes of the desired LMO phase are apparent in the diractogram.

For a complete structural analysis of such thin samples, special XRD techniques,

such as grazing-incidence XRD would be required, which is beyond the scope of this

work.

The diraction data of the LTO sample is shown in Fig. 5.4. Only one, broad

reex is apparent in the as-deposited diractogram, whereas post-annealing, 4 re-

exes are visible. Again, the annealing step increased the crystallinity of the sample,

however, no denite structural assignment could be done. The reexes could be as-

signed to various phases of LTO. The weak and broad reexes do not allow for a clear

structural assignment. Possibly a mixture of various titanate phases was present in

the lm. Thicker lms or more advanced XRD techniques would be required for an

unambiguous structural investigation.

58

5.3 Electrochemical Properties of Sputtered Thin

Films

A symmetrical sample, consisting of a polycrystalline LLZTO pellet (1 mm thick,

cross section about 50 mm2) with LMO thin lms (60 nm) and Au contacts (100

nm) deposited onto both sides was prepared (details s. Sec. 4.1). Further, an

asymmetrical sample, consisting of an LLZTO single-crystal (7 mm2, 1 mm thick),

sputtered with LMO (20 nm) and LTO (31 nm) thin lms and Au contacts (20 nm)

was electrochemically characterized via EIS and DC-cycling.

5.3.1 Impedance Spectroscopy

The symmetrical polycrystalline LMO sample was characterized via EIS from 1 MHz

to 100 mHz as deposited and after an annealing step, since the lms were deposited at

room temperature. Fig. 5.5. shows impedance spectra of the as-deposited sample,

measured at room temperature and at 7.8 C as well as an impedance spectrum

of the annealed sample, measured at room temperature. Since the ionic mobility

of the sample is enhanced at higher temperatures, the overall impedance of the

sample is lower at room temperature. The as-deposited measurements show two

semicircles. At room temperature, the minimum between the high frequency and

low frequency semicircle is located at about 3 · 105 Ω. At 7.8C, the two semicircles

partially overlap, giving a relative minimum at about 1.6 ·106 Ω. The high frequency

semicircles can be attributed to the ionic conductivity of the electrolyte, since the

capacitance of the feature is in the range of 10−11 F. With the dimensions of the

electrolyte (50 mm2 by 1 mm), a relative electric permittivity of about 20 can

be estimated, which ts in the general range of oxide materials. For as-deposited

electrodes, the measured resistance of the electrolyte is high and therefore the specic

conductivity low (< 1 µS/cm), which is far lower than expected from a LLZTO

electrolyte. The low conductivity might be caused by bad electric contact between

the electrode materials and the electrolyte. The low frequency features of the spectra

can be attributed to processes at the electrode-electrolyte interfaces. However, the

corresponding capacitance (about 10−7 F) is lower than expected for a chemical

capacitance (i.e. (de)intercalation of Li-ions). Post annealing, the overall impedance

59

of the sample decreased signicantly. The onset of the electrolyte semicircle (s.

inset) is at about 2 kΩ, which corresponds to an electrolyte conductivity of about

0.1 mS/cm, a typical value for LLZO materials. The increased conductivity of

the electrolyte is attributed to better electronic contact between the electrodes and

the electrolyte, since the thermal treatment is not expected to change the specic

conductivity of the LLZTO. The low frequency feature consists of a shoulder and an

arc which does not show any tendency to reach the real axis. Rather, this electrode

arc seems to be followed by an additional arc with probably very high capacitance.

The asymmetric LTO/LLZO/LMO sample was characterized by EIS at room

temperature under argon atmosphere (s. Fig. 5.6). The displayed data was ob-

tained in the frequency range from 1 kHz to 1 mHz. At about 1.6 kΩ, a minimum

is apparent in the high frequency region (s. inset), indicating the impedance of the

electrolyte. Given the dimensions of the electrolyte, this corresponds to a conductiv-

ity of about 1.2·10−4 S/cm. A second semicircle arises from the minimum, ranging

up to 7·107 Ω in the low frequency part. This feature is attributed to processes

at the electrode/electrolyte interfaces due to its small capacitance, which is about

6·10−8 F (for 7 mm2 sample size). Then an additional capacitive contribution seems

to be present (vertical spike). A preliminary equivalent circuit (s. Fig. 5.6) t leads

to 7·10−5 F ·cm−2 which is realistic for a chemical capacitance of a 20 nm layer when

electrodes are in a charging state where large voltage increments correspond to small

charge transfer (e.g. close to LixMn2O4, x = 1). Since the cell was deposited from

a target with x=1, this is a plausible explanation for the small capacitance of the

electrodes.

5.3.2 DC Cycling

The annealed polycrystalline symmetric LMO/LLZTO/LMO sample was further

characterized by potentiostatic DC-cycling. The sample was cycled at room tem-

perature under air. The results are shown in Fig. 5.7. Positive currents arising from

positive voltages are apparent, which are decreasing after the voltage change. De-

creasing the voltage again corresponds to a discharging of the sample, which results

in negative currents. Like the charging currents, the discharging currents are de-

creasing in magnitude at constant voltage. The currents measured during charging

60

Figure 5.5: EIS spectra of the LMO/LLZO/LMO sample measured as-deposited atroom temperature (red) and 7.8 C (blue) and after an annealing step (black) atroom temperature. The equivalent circuit used for estimation of the capacitance ofthe low frequency feature is displayed. A magnied part of the diagram is shown inthe inset.

and discharging of the symmetrical cell are decreasing over time because the migra-

tion of Li-ions and electrons from the anodic to the cathodic side of the sample leads

to a partial compensation of the external electric eld. When voltage is applied, the

Li electrochemical potentials inside anode and cathode are shifted. Due to this gra-

dient, Li ions (and electrons) migrate from anode to cathode, thus compensating the

electrochemical potential gradient. Therefore, the driving force for ion and electron

migration decreases over time. The charging/discharging curves show values in the

range of µA, which is far higher than the current curves obtained from polarization of

LLZTO with blocking electrodes (cf. Part I). Since Li-ions can be (de)-intercalated

out of and into the electrodes relatively easy, ion migration and thus charge transfer

is achieved more easily compared to blocking cells. The resistance estimated from a

voltage of 1 V and a current of 0.12 µA is ca. 8 MΩ and thus, larger than the value

found in the impedance measurements for 1.5·105 Hz (43 kΩ). This also indicates a

61

Figure 5.6: EIS spectrum of the LTO/LLZO/LMO sample measured at room tem-perature under argon. A magnied view of the high frequency region is shown inthe inset.

further process at low frequencies in the impedance (see above). The charge own

in 1 hour is approximately 2·10−4 C corresponding to about 1.3·1015 Li atoms. A

60 nm LMO layer on a 50 mm2 LLZTO contains about 6·1016 Li ions and thus a

signicant amount of charging has occurred already after 1 hour for 1 V and even

more for 1 hour at 2 V (another ca. 5·1015 Li atoms). The subsequent discharging

(at 1 V and short circuiting via 0 V) seems to be strongly hampered since almost no

current ows. However, since negative currents arise during these discharging steps,

it can be concluded that some chemical charging has occurred during the charging

steps at 1 V and 2 V. Possibly the open circuit voltage is still low (due to a plateau

in the discharge curve) and the driving force for discharge is very small.

The asymmetric sample (LTO/LLTZO/LMO), was also characterized via DC

cycling under argon atmosphere at room temperature. Results from potentiostatic

measurements are shown in Fig. 5.8. Typical current proles arise from the poten-

62

Figure 5.7: Results from DC cycling of the polycrystalline symmetric LMO sample.The applied voltage (blue) is shown with the corresponding current (red line).

tiostatic cycling at both positive and negative voltages. The LTO electrode of the

cell was connected to the negative terminal of the voltage source. Therefore, a posi-

tive voltage corresponds to the usual polarity of a cell, that uses a LTO anode versus

a LMO cathode. However, negative voltages correspond to a deeply discharged state

of both electrodes, i.e. LTO depleted of Li to have less than the usual minimum 4 Li

per formula unit and LMO lled to more than the usual maximum of 2 Li per for-

mula unit. Therefore, the negative voltage steps of the potentiostatic DC cycling,

polarize the cell far below the usual minimum of LMO/LTO cells, which can be

accompanied by irreversible phase transitions and other detrimental side reactions

in the electrodes. Despite the relatively high negative voltage, the cell shows no

distinct signs of degradation caused by deep discharging.

The asymmetric cell was subjected to galvanostatic DC cycling under argon

atmosphere at room temperature (s. Fig. 5.9). The results show repeated charging

and discharging of the cell with currents of ±100 nA. Note, that 1 minute at 100 nA

corresponds to 4.8·10−2 Li per formula unit in a 20 nm layer of LiMn2O4 with a 7

mm2 cross-section. Positive current values correspond to a charging of the cell, which

require positive voltages. After each charging and discharging step, the open circuit

voltage (OCV) was measured by setting the current value to zero for 10 minutes.

The results show positive, but decreasing OCV values after positive polarization

63

Figure 5.8: Results from potentiostatic DC cycling of the asymmetricLMO/LLZTO/LTO sample at room temperature. The applied voltage (blue line)is shown with the corresponding current (red line).

and negative, increasing OCV values after negative polarization of the cell. The

high, negative voltages which are applied to meet the negative (i.e. discharging)

currents indicate that the small OCV voltage of the cell after charing, is not a strong

enough driving force to give the demanded discharging current, thus requiring an

additional driving force, i.e. negative voltage. Considering the negative OCV after

the negative polarization steps, the high negative voltages lead to a negative and

therefore inverted polarization of the cell. This inversion of polarity can only be

realized in a LMO/LTO cell, when the electrodes are discharged beyond the usual

cycling limits, being Li4Ti5O12 and Li2Mn2O4, see above. Discharging of the cell

beyond common limits is known to lead to degradation of the cell. Indeed, some

degradation of the cell is apparent in Fig. 5.9, however, the cell can be cycled

repeatedly.

After the galvanostatic cycling described above, the sample was further cycled

at 80C under Ar. Since the Li-ion transfer resistance of the interfaces and all

transport resistances are lowered, more distinct signs of reversible, electrochemical

(dis)charging of the sample were expected from cycling at an elevated temperature.

The sample was, again, subjected to galvanostatic cycling, with ± 100 nA. Between

each 10 minute (dis)charging step, the OCV of the cell was measured for 10 minutes.

64

Figure 5.9: Results from galvanostatic DC cycling of the asymmetricLMO/LLZTO/LTO sample at room temperature. The applied current (red line)is shown with the corresponding voltage (blue line). The voltage limit was set toabout 3 V. Therefore, when this limit is reached during cycling, the current dropsbelow the set value of ± 100 nA.

The results are displayed in Fig. 5.10. The last cell voltage (OCV) at 25 C was

about -1.5 V and not surprisingly the cell thus started at negative voltages (about

-1.1 V) also at 80 C. Surprisingly, however, the 100 nA current did no longer suce

to get positive voltages, as after 10 minutes.Both charging and discharging as well

as the open-circuit parts of the cycling show negative voltages. The OCV after

positive and negative current steps dier by about 0.5 V, lying in the range from

-1 V to -1.8 V. Both OCVs are increasing from cycle to cycle, indicating that some

irreversible reactions occur alongside the reversible dis/charging of the sample. The

results prove, that partially reversible electrochemical cycling of an LTO/LMO type

cell is possible in a negative voltage regime.

Cyclic voltammetry (CV) was carried out (Ar, 80 C) to further elucidate the

state of charge of the sample (sweeping rate = 0.1 mV/s). Since the sample was

found to be in a state of inverse polarity (i.e. -1.4 V) after the last measurements (s.

Fig. 5.10), the rst CV sweep started at -1.4 V, going up to 2.5 V. After the rst

sweep, the sample was cycled between 0 V and 2.5 V. The results are displayed in

Fig. 5.11. The current density arising during the rst sweep is clearly dierent from

the subsequent cycles between 0 V and 2.5 V. This is expected, since the sample is

65

Figure 5.10: Results from galvanostatic DC cycling of the asymmetricLMO/LLZTO/LTO sample at 80C. The applied current (red line) is shown withthe corresponding voltage (blue line). The voltages required to meet the set currentvalues of ± 100 nA are signicantly lower than they were at room temperature (c.f.Fig. 5.9).

changed from negative polarity back to the usual positive polarity. After the sample

is brought back to positive voltages, the cycles show little dierences and share the

same basic features. In the charging parts of the cycles, a very broad and at reaction

peak is apparent at about 1.5 V. After this feature, a steep increase of current occurs.

In the discharging parts, the curves show a broad and at reaction peak at about

1.0 V. The combination of the peaks at 1.5 V and 1.0 V during (dis)charing suggest

a reversible redox reaction taking place during cycling, which is typical for this kind

of battery. However, the steep rise of current at the end of each charging sweep

does not have a corresponding negative peak, thus indicating irreversible reactions

taking place above 2 V. Despite the distinct irreversible reactions at the end of each

sweep, the cycling of the sample shows no strong signs of degradation.

At the end of the CV measurements, the sample was left at a potential 0 V. Sub-

sequently, galvanostatic DC cycling was carried out again with the same parameters

66

Figure 5.11: Results from cyclic voltammetry of the asymmetric LMO/LLZTO/LTOsample at 80C.

as before. The results are shown in Fig. 5.12. The general shape of the voltage curve

is similar to the results from previous galvanostatic measurements (c.f. Fig. 5.10).

However, after charging the sample, a positive OCV is retained over 10 minutes.

When the current is switched to - 100 nA, negative voltages of up to -0.5 V arise.

As described earlier, this is attributed to an insucient discharging current of the

sample, which is compensated by application of negative bias to the sample. After

the negative current steps, the OCV of the sample is negative, quickly approaching

0 V. The results show clearly, that the sample was electrochemically charged, re-

versibly, at positive voltages. Thus proving, that the cell was successfully brought

back to the original polarity of LMO vs. LTO. The OCVs measured after charging

of the sample are between 0.5 and 0.8 V, increasing slightly from cycle to cycle.

Considering the redox peak from Fig. 5.11, the results suggest that the sample was

not charged long enough to reach the voltage plateau of the charging curve at about

1.5 V.

67

Figure 5.12: Results from galvanostatic DC cycling of the asymmetricLMO/LLZTO/LTO sample at 80C. The applied current (red line) is shown withthe corresponding voltage (blue line).

68

Chapter 6

Conclusion

The results from Part I show that the fast Li-ion conductor LLZTO can be sub-

jected to DC voltages of several 100 mV via blocking electrodes without irreversible

reactions occurring. For Au microelectrodes, this leads to electronic steady state

currents. These measurements enable estimates of electronic conductivities and val-

ues in the range of 10−9 S/cm are found for 300-400 C. At higher voltages, however

decomposition of the material occurs. In macroscopic samples this leads to some

local conductivity changes close to the anode. Moreover, drastic chemical changes

were found by LIBS at the anode (Li depletion) and at the cathode (Li accumu-

lation). Moreover, morphological changes take place for higher voltages: Bubble

formation at the anode due to oxygen evolution (i.e. in sum Li-oxide loss) and

growth of an additional phase at the cathode (probably Li2CO3). In experiments

with polarized microelectrodes, on the other hand, irreversible changes are reected

by irreversible current changes.

The results of the second part show that the construction of a sputtering de-

vice for oxide materials was successful. Deposition rates of various metallic (Cu)

and ceramic materials (LMO and LTO) were determined. Inuence of deposition

temperature on the structural and electrochemical properties of active electrode ma-

terials were investigated. The results show a clear correlation between crystallinity

of thin lms and their electrochemical properties. A positive eect of thermal an-

nealing on structural and electrochemical properties was demonstrated.

Sputtered thin lms were used to build all solid state cells with LLZTO sin-

gle crystals and polycrystals. Reversible, electrochemical cycling of an all solid

69

state sample (LMO/LLZTO/LTO) was demonstrated in both positive and nega-

tive voltage regimes. It was shown, that after deep-discharging of the LTO/LMO

cell, reversible electrochemical (dis)charging can be carried out over several cycles

at negative voltage. Further, the cell was successfully cycled repeatedly at positive

voltages, after the cell was subjected to deep discharging.

70

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