Title: Deposition, corrosion and coloration of tungsten trioxide electrochromic thin films
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 Material Information
Title: Deposition, corrosion and coloration of tungsten trioxide electrochromic thin films
Alternate Title: Tungsten trioxide electrochromic thin films
Electrochromic thin films
Physical Description: vii, 141 leaves : ill. ; 28 cm.
Language: English
Creator: Sun, Sey-Shing, 1955-
Copyright Date: 1983
 Subjects
Subject: Thin films   ( lcsh )
Materials Science and Engineering thesis Ph. D
Dissertations, Academic -- Materials Science and Engineering -- UF
Materials Science and Engineering thesis Ph. D
Dissertations, Academic -- Materials Science and Engineering -- UF
Genre: bibliography   ( marcgt )
non-fiction   ( marcgt )
 Notes
Statement of Responsibility: by Sey-Shing Sun.
Thesis: Thesis (Ph. D.)--University of Florida, 1983.
Bibliography: Bibliography: leaves 132-139.
General Note: Typescript.
General Note: Vita.
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Bibliographic ID: UF00099355
Volume ID: VID00001
Source Institution: University of Florida
Holding Location: University of Florida
Rights Management: All rights reserved by the source institution and holding location.
Resource Identifier: alephbibnum - 000447021
oclc - 11386614
notis - ACK8309

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DEPOSITION, CORROSION AND COLORATION OF
TUNGSTEN TRIOXIDE
ELECTROCHROMIC THIN FILMS








BY


SEY-SHING SUN


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF
THE UNIVERSITY OF FLORIDA
IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE
DEGREE OF DOCTOR OF PHILOSOPHY


UNIVERSITY OF FLORIDA




































To my mother

and

Mei-Hwei, my beloved wife

















ACKNOWLEDGEMENTS


The author would like to acknowledge and thank his advisory

committee and, in particular, to express his gratitude to Professor

P. H. Holloway, who convinced the author of the importance of his

research and maintained an open and creative research environment

throughout the course of this work. The author is also grateful to

Professor J. R. Ambrose, J. Ali and R. U. Lee for their assistance

and useful discussion during corrosion studies. In addition, thanks

are extended to Drs. G. E. McGuire and R. T. Tuenge of Tektronix, Inc.,

who provided the author with internship and encouragement for the

development of practical electrochromic display devices.

Thanks are also due to A. Haranahalli, K. Shanker and Y.-X. Wang

for their technical assistance and encouragement.

Finally, acknowledgements are due to the Army Research Office,

Durham, NC., for research support under grant #DAAG-29-80-C-0007.



















TABLE OF CONTENTS

PAGE

ACKNOWLEDGEMENTS .......................... ...................... iii

ABSTRACT ................. .............................. vi

CHAPTER

I. INTRODUCTION................. ............................1

II. CORROSION STUDIES AND MODIFICATION OF TUNGSTEN
TRIOXIDE THIN FILMS BY OXYGEN BACKFILLING .............6

Introduction .............. .... ............ ............. 6
Experimental. .............. .............. .............. 11
Film Preparation.................. .................. 11
Chemical Characterization ............................ 12
Physical Characterization............................. 15
Corrosion Measurements.................... ........... 16
Electrochromic Properties Measurements............... 16
Results. ............... .................. .............. 20
Corrosion Studies........... ........................ 20
Modification of WO3 Films by Oxygen Backfilling......26
Discussion .. ......... ................... ............... 44
Corrosion of WO, Thin Films...........................44
Modification of Stability of WO Films by
Oxygen Backfilling... ... .................... .... 61
Modification of EC Properties of WO3 Films by
Oxygen Backfilling................................. 62
Summary ............ ........................ ............. 80

III. DEVICE FABRICATION........................... ........ .84

Introduction ............... ............................ 84
Experimental ............................................ 85
Preparation........... .. ...... ......... ............. 85
Characterization. .... .................... ............86












PAGE
CHAPTER

Liquid Electrolyte Device.............................88
Results ........................................... 88
Discussion........................................ 97
Solid Electrolyte Device.............................105
Results...........................................108
Discussion.........................................111

IV. SUMMARY................................................ 113

Conclusions..........................................113
Future Developments..................................119

APPENDICES

A. PRELIMINARY STUDIES OF CATHODOCHROMISM IN TUNGSTEN
TRIOXIDE FILMS AND THE EFFECT OF AIR BACKFILLING..121

B. CORROSION OF TUNGSTEN TRIOXIDE FILMS DURING
STORAGE IN MOIST AIR..............................128

BIBLIOGRAPHY ..................................................132

BIOGRAPHICAL SKETCH ...........................................140


















Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirement for the Degree of Doctor of Philosophy

DEPOSITION, CORROSION AND COLORATION OF
TUNGSTEN TRIOXIDE
ELECTROCHROMIC THIN FILMS

By

Sey-Shing Sun

December, 1983


Chairman: Paul H. Holloway
Major Department: Materials Science and Engineering



Electrochromism (reversible color change of a material, induced

by an electric field) has been investigated in tungsten trioxide

thin films. It is known that electrochromic device lifetime is

generally limited by corrosion of WO films in an acid electrolyte.

Two mechanisms of corrosion were observed, viz., general dissolution

and interfacial delamination. It was found that the dissolution of WO3

films in acid could be attributed to high concentration of thermody-

namically unstable species (WO2 and W205) in the as-deposited films.

These species were identified by X-ray photoelectron spectroscopy

data and are consistent with Rutherford backscattering spectroscopy

data which showed an oxygen deficiency (0/W = 2.76). The Pourbaix

diagram for tungsten indicated that WO3 was the thermodynamically

stable specie for the present storage condition (pH = 0.5,










E = 0.4 VSHE). A corrosion mechanism was proposed consisting of

dissolution of WO2 and W205 and precipitation of crystalline WO3 and

its hydrates. Interfacial delamination occurred when WO3 and its

hydrates precipitated back onto the original films.

The oxygen content in the WO3 films was increased by oxygen back-

filling during evaporation. Dissolution and interfacial delamination

of the oxygen enriched films were reduced to negligible rates due to

reduced concentration of W02 and W205. However, the electrochromic

properties were degraded by oxygen enrichment. For example, increased

resistivity and decreased optical efficiency in the oxygen enriched

films resulted in slower coloration speed. The resistivity increase

and decreased optical efficiency were explained by postulating an

increased density of inactive electron trapping sites. The porosity

of the films could be increased by deposition at high background

pressure, resulting in increased surface area and absorbed water.

The bleaching speeds and self-erasure rates were increased since the

rates of removal of protons were increased by the increases in

porosity and absorbed water.

In another approach to increase the electrochromic device lifetime,

the electrolyte was modified. Devices using a solution of LiCIO4 in

propylene carbonate exhibited excellent lifetime. Switching speeds

were increased by increased porosity, deposition of MgF2 overlayers,

and more conductive indium-tin-oxide layers. In addition, solid state

electrochromic devices using a hydrated MgF3 film were fabricated.
















CHAPTER I
INTRODUCTION


Electrochromism is broadly defined as the ability of a material

or system to change color reversibly in response to an applied

potential. A great variety of electrochromic materials or systems

are know and can be conveniently classified into the following

categories according to the physical mechanisms involved:

reversible electrodeposition systems

organic ion-insertion materials

inorganic ion-insertion materials.
1 2
Viologens, anthraquinones and electroplating of very thin films of

metals such as silver comprise the first class. Diphthalocyaniens

and tetrathiafulvalenes5 are included in the second class. Class

three is comprised of transition metal oxides, among which WO3 thin
6-39
films have received by far the most attention. Other materials
40-44
such as anodically deposited iridium oxide films of IrO2, molyb-
45-47 48
denium oxide (MoO3), vanadium pentoxide (V205), and nickel

oxide (NiO)49 have also been shown to color electrochromically, and

would be placed in this class. Besides coloring electrochromically,

some of the transition metal oxides may also be colored by heating




Although the vapor deposited tungsten trioxide films are substoichio-
metric as reported below, the term 'WO3 thin films' will be used for
simplicity unless otherwise noted.












(thermochromism), by exposure to ultra-violet light or X-ray

(photochromism)7 or by electron beam irradiation (cathodochromism).5052

The WO3 electrochromic (EC) display device typically consists of a

transparent electrode such as a SnO2 coated glass substrate, a thin

layer of transparent WO3, an electrolyte and a counter electrode

(Pt or Au). When a small negative voltage ( 1.0 V) is applied to

the transparent electrode, the WO3 film is colored blue in trans-

mission.89 The induced optical change is proportional to the charge

injected during the coloring pulse. If the voltage is removed, the

WO3 film remains colored. Reversing the polarity of the applied

voltage will cause bleaching of the films. The electrochromic

coloration of WO3 in acidic aqueous solution involves the injec-

tion of electrons and protons into WO3 films. The blue coloration

in transmission is associated with the broad optical absorption band

which gives an absorption minimum at wavelength of 0.4 pm and an

absorption maximum at 0.95 pm.89 The absorption band is generally

explained by electron transfer between neighboring tungsten atoms,

i.e., optical absorption by excitation of polaron hopping.1012

The protons injected are for the purpose of charge compensation.
+ 13-21 + 25-28
Therefore, any monovalent cation can be used, e.g., Li Na ,
+ 29
or Ag .

The electrochromic device exhibits many attractive features as

displays, such as

good color contrast especially under high level of ambient
light;

wide-angle viewability which is an advantage over most
liquid crystal displays (LCD);












low voltage operation (1.5 V down to approximately 0.5 V);
2
low power, %9 mW/cm2, somewhere between LCD and light-
emitting diode (LED) displays;

storage of the display without power dissipation;

potentially low cost.

The response time has been reported to be approximately 50 ms with
30
a modified device configuration in acid electrolyte, and is

comparable with LCD. However, there is a principal obstacle to

the commercialization of EC displays, viz., short device lifetime.
5 31
The general lifetime reported was about ',10 cycles at 0.5 Hz or

less than two weeks for static condition. To understand the cause

of degradation and to find a way to improve the lifetime was the

major goal in the present research.

The degradation of EC devices was reported to result from
32-35
corrosion of WO3 thin films in the acid electrolyte. In

Chapter II, the corrosion mechanism and the results of using oxygen

backfilling during WO3 evaporation to prevent corrosion will be

discussed. The durability of the deposited films in sulfuric acid

solution will be shown to be limited by two mechanisms-- a general

uniform film dissolution and an interfacial delamination. Due to

the deficiency of oxygen (0/W = 2.76), some of the configurations of

tungsten atoms in as-deposited films may be thermodynamically unstable.

A conversion process, viz., dissolution of these unstable

species and precipitation of the stable crystalline-WO3 and its

hydrates, leads to the general film dissolution and interfacial











delamination of WO3 films. The delamination of films from the sub-

strates occurs when the internal stresses induced by precipitation

exceed the interfacial adhesion strength.

In order to improve the lifetime, two basic approaches were

adopted in the present study. In one approach, oxygen backfilling

during WO3 evaporation was used to modify film structure and

reduce the attack by the acid electrolyte. By this method, the

O/W ratio was increased and the amount of degradation was reduced

due to the increased concentration of the thermodynamically stable

specie, WO3. The current injection and optical efficiency during

coloration were lower, however, while the bleaching speed and

self-erasure rate were faster for oxygen-enriched films. The

mechanisms by which these modifications of EC properties were

achieved are discussed.

The second approach to improve the device lifetime involved

the use of non-aqueous electrolytes. Non-aqueous electrolytes

have been shown to provide an inert environment where the corrosion
15
of WO3 was not observed. In Chapter III, trial devices using

LiCIO4 dissolved in propylene carbonate (Li/PC) and solid MgF2

thin film electrolytes are discussed. Previous investigators

have reported that devices using Li/PC as an electrolyte exhibited

slow response speed. In this study, an improvement in speed was

achieved by nitrogen backfilling during WO3 deposition and by the

use of a MgF2 overlayer. Layers of MgF2 were also used as elec-

trolyte to fabricate solid state display devices. These experi-

ments indicated that the increased water content in MgF, films





5




resulting from air backfilling increased the coloration/bleaching

(C/B) speed, while the operational area of the device was limited

by the air backfilling during deposition of WO3 films.

Preliminary data on cathodochromism (electron beam induced

coloration) in WO3 thin films are reported in Appendix A. The

coloration density in WO3 films was found to be increased by air

backfilling, especially for higher energy electron beams.

Results indicate that cathodochromism in amorphous WO3 thin

films may still be caused by polaron hopping absorption as in

electrochromism.
















CHAPTER II
CORROSION STUDIES AND MODIFICATION OF TUNGSTEN TRIOXIDE
THIN FILMS BY OXYGEN BACKFILLING


Introduction

Electrochromism (EC) has been investigated in tungstic oxide

films for possible application in display devices. The devices

typically used diluted sulfuric acid solutions as the electrolyte

With a modified device configuration (Figure 2.1), viz., a porous gold

overlayer on the WO3 films, it is possible to achieve a switching

time of approximately 50 ms.30 The time is comparable to that

of liquid crystal displays (LCD). However, the lifetime of EC

devices in such an electrolyte is normally short. The device

lifetime was found to be limited by the degradation of WO3 films

in the acid electrolyte in the form of general dissolution and
32-35
delamination.32-35 Short device lifetime is a principal obstacle

to the commercialization of EC displays.

Degradation of WO3 thin films has been studied and was
32-35
correlated with corrosion, However, the mechanisms) by

which corrosion occurs and the thermodynamics of WO3 corrosion

have not been considered in detail. Faughnan and Crandall35

studied evaporated W03 films in H2SO /water solution (pH = 0) and

found the dissolution rate was approximately 300 A per day. Cyclic

coloring and bleaching of the films led to faster dissolution rates.



























Au- \ \\ In
In ..............
7SnO2 .. 3:: : 3

Glass








Figure 2.1. Schematic representation of the
electrochromic cell with thin
gold overlayer on WO3.











32
Randin32 replaced the water with glycerin as the solvent in the

electrolyte and found the dissolution rate was lowered by an order

of magnitude (approximately 20 A per day). An increase in the

dissolution rate was also observed under cycling condition.

The film was eventually destroyed as a result of delamination

from the substrate. To explain this behavior, Randin listed a

number of possible mechanisms of degradation, e.g.:

a) proton insertion process causing hydrogen embrittlement
which undermined the internal strength of WO3,

b) anodic disintegration of the films,

c) oxidative degradation occurr ng reagdly on semiconductor
electrodes, such as WO + 6h -- W (sol'n) + 3/2 02,
where h designates hole carriers.

However, no specific data and discussion supporting the above

postulates were given. He also found WO3 thin films were more

stable in some non-aqueous solvents such as propylene carbonate,

acetone and Y-butyrolactone. However, the response speed of the

device using an electrolyte with such a solvent was too slow to

be practical.

Reichman and Bard33 reported that both the kinetic behavior

and the stability of the films in acid electrolyte depend on the

amount of water and the porosity of the films. From infrared

absorption spectra, they found the water content was increased

when the films were cyclically colored and bleached. On the basis of

these observations, they speculated that during operation, uptake of

water into the films caused expansion of the lattice, making it

more porous and allowing even greater water penetration. The











result was higher water content in the films after cyclic colo-

ration and bleaching. They also measured the dissolution rate of

WO3 film in acid by monitoring the decrease in current injected

during cyclic coloration and correlated it with the increase in

water content of WO3 films. They concluded that the dissolution

was due to the formation of WO3 hydrates (WO3.H20 or WO3.2H20)

since these hydrates are more soluble than WO3. Hence, the higher

the water content of the film, the more WO3 dissolved. No model

of dissolution was proposed, nor was saturation of the electrolyte

discussed. In addition, no data were shown demonstrating the

formation of hydrates. Reichman and Bard also observed slow

dissolution in glycerin/H2SO4 (10:1) as did Randin; they attri-

buted this to the lack of water pickup during cycling.
34
Arnoldussen34 studied the dissolution of the evaporated W03

films in deionized (DI) water and concluded the films dissolve

to form metatungstate or pseudometatungstate ions. On the basis of

previous study of vapors generated by subliming WO337 he postu-

lated that as-deposited WO3 films are composed chiefly of trimeric

clusters weakly bound to one another through water-bridge or

hydrogen bonds. The dissolution of the film was therefore pos-

tulated to occur by these trimeric clusters' being dissolved as

3-, 6-, or 12-mer ionized complexes directly from the dissolving

surface. He also suggested cyclic conditions were just a natural

extension of the dissolution process, i.e., voltage enhanced

dissolution. Degradation in non-aqueous solvents was suggested












to result from formation of negative complexes with electrolyte

anions or from tungstate formation due to H20 incorporated into

the films during manufacture. However, all of Arnoldussen's

data dealt with electrolyte of pH > 4, and the corrosion reaction

proposed was based on metatungstate ions; it has been shown that


2
WO is the dominant specie for dissolution of tungstic oxides

in acid solution of pH < 1.38 The O/W ratio in Arnoldussen's

microstructure model is always 2 3, since the oxygen is bound

either in W309 or H20. It will be shown in this study that the

O/W ratio was always < 3. In addition, the idea of H20 stabilized

W309 rings is unreasonable, since the disintegration of the films

would be expected if H20 was removed. However, Zeller amd Beye-
39
ler39 reported H20 was removed at 1700C without detectable change

in the structure of WO3 films. In fact, no structural change
32
was observed until crystallization at 3500C, which is well

above the H20 removal temperature.

As is evident from reported studies, a variety of mechanisms

have been proposed for degradation of WO3 films. All of these

investigators have speculated that water is a major cause of

corrosion. The mechanisms proposed are normally not supported

by strong and direct evidence, and are sometimes contradictory

to existing data (e.g., Arnoldussen's H20 model discussed above).

However, it is generally recognized that WO3 thin films in acidic

electrolyte degrade in the form of general dissolution and dela-

mination. In cycling conditions the dissolution was increased.













In the following, new mechanisms which can fully explain these

phenomena without the shortcomings of the existing postulated

mechanisms will be discussed.

There are several ways to counter corrosion problems, e.g.,

1. modify the electrolyte, 2. put protective coating on the WO3

film, 3. modify the WO3 thin film microstructure. In the second

half of the chapter, stability data will be reported for WO3 films

evaporated with oxygen backfilling (OB). The O/W ratio will be shown

to be increased by backfilling. This led to reduced degradation.

However, the coloration speed and optical efficiency were decreased

while the bleaching and self-erasure rates were faster for the oxygen

enriched films. The mechanisms by which these EC properties were

modified are also discussed.



Experimental



Film Preparation

Thin films of WO3 were prepared by evaporation of 99.9% pure

WO3 powder (Cerac, Inc.) from a resistivity-heated, alumina-coated

tungsten spiral boat (Sylvania, Inc.). The substrate was normally

tin oxide coated glass (NESA glass, Pittsburgh Plate and Glass Co.)

with a resistance of 100 2/[, but soda-silicate glass slides and pure

graphite substrates (Ernest Fullan, Inc.) were used occasionally.

Prior to deposition, the substrates were cleaned with Alconox/water

solution, rinsed in deionized water for 10 minutes, etched in











Chromerge (a registered brand name of Monostat Labs representing a

mixture of chromic acid and sulfuric acid, 1:10) and degreased in

boiling methanol vapors for 5 minutes. The evaporation system was
-6
evacuated to a base pressure of 2 x 106 Torr and WO3 was evaporated
-5
either at a residual pressure (P ) of 1 x 10 Torr. or with a
res.
-5 -3
partial pressure of oxygen (PO ) between 5 x 10 and 1 x 10 Torr.
2
Oxygen (99% pure) partial pressure was manually controlled by a needle

valve. The deposition rate was typically 10 A/sec. The thickness

of the films ranged from 2,000 A to 4,000 A. The source to sub--

strate distance was 30 cm. No substrate heating was used during

evaporation.



Chemical Characterization

Rutherford backscattering spectroscopy (RBS) was used to deter-

mine the oxygen to tungsten ratio (0/W) of the evaporated films.

Briefly, a beam (300 nA, 300-600 pC) of monochromatic (2 MeV) alpha

( He ) particles was focused (rectangular area 2 mm x 4 mm) onto a

WO3 films on a graphite substrates. Alpha particles were scattered

through 1550 by W and O atoms in the films losing an energy which

depends on the masses of these atoms.53 Analysis of the number of

backscattered alpha particles versus their energy allowed the O/W

ratio to be determined.

A typical RBS spectrum is shown in Figure 2.2, which is plotted

as the backscattered yield (counts per second) versus the channel


























8








4
















on graphite substrate using alpha particle of 2.0
MeV.
6



















MeV.












number. The channel number is related to the backscattered particle

energy by a conversion constant (eV/channel number). The concentra-

tion of the elements in the films appears as square-top peaks in the

spectrum. The ratio of the different elements in the films can be

obtained by calculating the ratio of the area under the respective

peaks and correcting for variations in their scattering cross sections.

Since the width of the peak corresponds to the energy lost by alpha

particles which traveled through the films, the thickness can be

obtained by converting the peak width to depth using calculated
40
stopping powers.40 The shape of the peaks also provide information

about the structure of the films (uniformity, interdiffusion, etc.).

Compositional information about light constituents such as

alkali and water were obtained from secondary ion mass spectroscopy

(SIMS).54 The experimental system incorporated a UTI quadrupole

mass filter (UTI 100C) with a 3M prefilter (Model 610), a 3M mini-

beam ion gun, and a sample manipulator (Huntington PM-600XYZTRC).
-9
In this study, the vacuum system was pumped down to 1 x 10 Torr

and was then backfilled to 5 x 10-5 Torr with argon. A primary ions

energy of 4 KeV and beam current of 100 nA was used during the analysis.

X-ray photoelectron spectroscopy (XPS)55 was used to obtain

the information on the concentration of different oxidation states

of tungsten ions in WO3 thin films. The analysis was performed with

a Krato XPS system (XSAM 800) using Mg K-a X-rays (1253.6 eV). The

samples were maintained at room temperature and a pressure of 5 x 10-

Torr.











Physical Characterization

The relative density of evaporated WO3 films was changed by

oxygen backfilling. A simple method using a quartz oscillator

thickness monitor (Piezo Technology Inc. Model 990) was used to

measure the areal mass change. The areal mass measured by the

monitor was fixed for deposition at all pressure. The physical

thickness of the deposited films was then measured by a profilo-

meter (Sloan Dektek II). An arbiturary index was obtained when the

physical thickness was devided by the thickness calculated from the

areal mass by assuming the theoretical density of WO3. This index

was used to compare the relative densities of the films. Thin films

of WO3, with and without oxygen backfilling, were studied for surface

damage (before and after corrosion) with a JOEL scanning electron

microscope (Model JSM-35C).

X-ray diffraction spectra were obtained from as deposited and

corroded films. Nickel-filtered Cu-Ka radiation at 40 KV and 20 mA

from a Phillips Electronics Instrument Model 690 spectrometer were

used. A range of 10 < 29 < 90' was scanned at a speed of 2'/min.,

with a counting time constant of 1.0 second. A Bicron model SR-

BP941 scintillation detector was used.

Films of WO3 deposited on glass substrates could be colored to

blue by touching an indium wire to the film surface through an acid

electrolyte (3.6 N H2S04). The coloration occurs by the injection
3+
of electrons provided by the corrosion reaction: In -- In3+ + 3e ,

and protons being injected into the film from the electrolyte to










maintain charge neutrality. The coloration progressed uniformly

outwards from the point of contact of the In wire with a circular

color front. This phenomenon prompted Crandall and Faughnan56 to propose

a simple method for determining the electron diffusion coefficient.

In this method, the radius of the color front was plotted as a function

of the square root of the time. The electron diffusion coefficient

(De) was obtained from the slope which is equal to 4-D .

Resistivity was measured immediately after deposition without

breaking the vacuum. The electrodes were parallel etched stripes of

SnO2, 0.1 cm wide with a gap of 0.5 cm. The current upon application

of 10 V was measured by an electrometer.



Corrosion Measurement

In the corrosion study, the WO3 films deposited on graphite

substrates were immersed horizontally in the unstirred acid electro-

lyte (3.6 N H2SO4 in water) contained in glass dishes at room

temperature. The electrolyte was not deaerated during the experiment.

The samples were removed at selected time intervals and examined by

RBS to measure the WO3 thickness. The average dissolution rate was

obtained by dividing the change in thickness by time.



Electrochromic Properties Measurements

Optical absorption spectra were obtained from films deposited

on glass slides both before and after coloration. The equipment












consisted of a quartz halogen lamp (Oriel Model 6140), a mono-

chromator (Oriel Model 7240) and a silicon (for 400-1000 nm) or

PbS (for 1000-2000 nm) detector. The range scanned was 300 to

1500 nm with measurements being taken in transmission.

The electrochromic (EC) properties (coloration/bleaching

speed, optical efficiency and open circuit memory) were charac-

terized in the electrochromic cell shown in Figure 2.3. EC

properties are defined as follows

Coloration time (t ) is defined as the time needed to color
a film to a contrast of 50% absorption;

Bleaching time (tb) is that time required to erase the film
from 50% to 25% absorption;

Optical efficiency is defined as optical density change per
unit charge injected;

Open circuit memory is defined as the time needed to bleach
the film to half of the original contrast after the power was
disconnected.

In the following discussion, the C/B speed will be used which is

defined as the rate of coloration or bleaching to reach defined

contrast, and is inversely proportional to C/B times. Electrical

contact to the SnO2 layer was made with indium solder outside the

electrolyte. The active area was 1.5 cm". The cell was made of

Teflon spacers and stainless steel holders for the glass plates.

An optical path through the cell allowed electro-optical measurements.

The electrochemical data were taken using the circuit shown in

Figure 2.4. A potentiostat (Princeton Applied Research Model 173)


















Luggin -- Reference
Capillary Electrode





Bolt

Steel Spacer
Plate



Glass
NESA /Electrolyte
Glass


Working Pt Counter
Electrode Electrode


Figure 2.3. Schematic representation of the
electrochromic cell.


























































Figure 2.4. Circuit diagram showing the connection between
EC cell, potentiostat and measuring devices.
WE: working electrode, RE: reference electrode,
CE: counter electrode.












was used as the power supply. Platinum wire and a calomel electrode

were used as the counter electrode and reference electrode, respect-

ively. Monochromatic light (700 nm) was used throughout the experi-

ment.



Results




Corrosion Studies

The durability of vapor deposited WO3 films in acid electrolyte

was limited by two mechanisms a general uniform film dissolution

and an interfacial delamination. As prepared, WO3 films exhibited

featureless smooth surfaces with some isolated pinholes detected with

SEM at magnification of x500 (Figure 2.5). After storage in the acid

electrolyte, WO3 films prepared without OB begin to degrade by general

film dissolution, formation of crystallite-like particles, wrinkles

spreading in a rosette-like pattern (Figure 2.6), and inter-

facial delamination of the films from the glass substrates (Figure 2.7

(A)). The dissolution rate was very high at '16 A/hr (-350 A/day)

as shown in Figure 2.8. However, the dissolution rate in the static

condition (i.e.,without electrochemical cycling) may have been

enhanced by the graphite substrate, as will be discussed in the

following section. The static dissolution rate is expected to be

lower for films on glass substrates.

The peaks in the RBS spectra of WO3 films also exhibited a

significant change after corrosion (Figure 2.9). The original W































(A)






















(B)

Figure 2.5. Scanning electron micrographs of as-deposited WO3
films on glass subs rates;
(A) P = 1 x 10 Torr, x 500,
(B) Pres. = 1 x 10 Torr, x 3,000.
2















FLA : a

wwJ8r ^ffw


V .- (

.A$?
0*R^


Figure 2.6. Scanning electron micrographs of corroded surface_5
(3.6 N H SO 4 days) of WO films (P = 1 x 10
Torr) showing the forms of interfacial delamination;
(A) wrinkles, (B) rosette pattern.





















































Figure 2.7. Scanning electron micrographs of WO3 films after storage
in 3.6 N H SO; -
(A) P e 1 x 10 Torr, 2 days, x 500,
(B) Pres. 1 x 10 Torr, 8 days, x 3,000.
O
2





















25






20






15














5






0







Figure 2.8. Dissolution rate of WO3 films in 3.6 N H2SO4
solution versus the oxygen partial pressure
during deposition.
i-4






















during deposition.





























S (a) (


6 (b)

o
64- 1






2.


100 300 500 700 900

Channel Number

Figure 2.9. Effect of corrosion on the RBS spectra for WO3 films
deposited without oxygen backfilling;
(a) as-deposited,
(b) after 5 hr. in 3.6 N H SO
(c) after 20 hr. in 3.6 N S .
2 '












peak was rectangular in shape indicating a uniformly thick film.

After corrosion, the peak was triangular in shape which indicates a

variable film thickness consistent with micropore formation by corro-

sion.



Modification of WO, Films by Oxygen Backfilling



Chemical and physical properties

The compositional changes caused by oxygen backfilling are

shown in Figure 2.10 and 2.11. From RBS data, the average O/W ratio
-5
increased from 2.7 with no oxygen backfilling (P = 1 x 10 Torr)
res.
-3
to 2.9 at P = 1 x 10 Torr. Since the films were hydrated as shown

later, the oxygen detected may have included that from the incorporated

water. The actual number of oxygen atoms that were directly bonded

between tungsten atoms might have been even lower. SIMS data indicated

that the content of water in the films increased by about a factor

of two between the two extreme cases (without oxygen backfilling and
-3 *
P = 1 x 10 Torr) The films prepared at the two extreme conditions

were examined by XPS. The spectra of the W4f core levels for these

films are shown in Figure 2.12. For comparison, the core level spectra

for WO3 on W representing a W state and the positions of the reduced

W states (W 4+) are also shown. The multiple split W4f peaks





For simplicity in the fol owing discussion, 'WO 7 (without OB)
and 'WO, (P = 1 x 10 Torr) are used to den6te WO films prepared
at these two 2 extreme oxygen partial pressure conditions.


















3.00









2.90

0
o






2.8C









2.7C \ I

10-5 10-4 10-3

PO2(Torr)

Figure 2.10. Oxygen to tungsten atomic ratio of WO3 films from
RBS data versus the oxygen partial pressure during
deposition.




28
















---1-------------------------

3.0






2.0

a




1.0



C .D









Figure 2.11. Relative signal intensity for H20 from secondary ion
mass spectroscopy (SIMS) versus the oxygen partial
pressure during deposition of WO3 thin films.
oo .3






















































40 35 30

Binding Energy (eV)


XPS W(4f) core level spectra for WO3 films
prepared at two extreme conditions, i.e.,
'WO2 and 'WO 9'. For comparison,,6he
spectra for WO on W representing a W stas 5
and the position of the reduced W states (W W5
are also shown.


Figure 2.12.











of 'W02.9' spectra are located at the same energies and exhibit

similar intensity as those peaks from W+ in the reference spectra.

The spectra of 'WO2.9' therefore indicate that the tungsten exsited

mostly as W6+ in these films. On the other hand, the spectra of

'W2.7' exhibit a filled valley between and a shoulder on the low
6+
energy side of the multiple split W4f peaks from W This indicates
4+ 5+
that the concentration of subvalent W and W were high in 'WO
2.7

films. Therefore, the oxygen backfilling increased the oxygen content

in the films as indicated by RBS data, and led to a concentration

increase in W states which was identified by XPS data.

The relative density of WO3 films was found to decrease

as a result of OB as shown in Figure 2.13. The unmodified films

exhibited a relative density of 0.8 when compared to the normalized
-3
bulk density. For films prepared at P = 1 x 103 Torr, the

relative density decreased to 0.5 and the porosity was expected to

be substantially higher.

The surface topography of WO3 films was modified by oxygen

backfilling as shown by SEM photomicrographs taken at high magnifi-

cation (x 3,000, see Figure 2.5(B)) for films deposited at P 2= 5 x 10

Torr. The rough surface is consistent with the higher porosity in

these films.




























- *


I II I -r I

10-5 10-4 10-3

S0 (Torr)


Figure 2.13.


The relative density of WO films versus the partial
pressure of oxygen during deposition


i










Stability

Interfacial delamination was not observed for films prepared at
-4
higher partial pressures of oxygen, i.e., >5 x 10 Torr. However,

a slightly corroded surface which was rougher than the original

surface was observed at higher magnification (x 3,000, see Figure

2.7(B)). The dissolution rates for these films were negligible as

shown in Figure 2.8.

The X-ray diffraction pattern for the as-deposited films with or

without OB showed a very broad diffused peak typical of an amorphous

structure (Figure 2.14(A)). After 18 days storage in acid electrolytes,

the X-ray diffraction patterns exhibited crystalline peaks on top of

the diffuse scattering from the amorphous film. These data are

interpreted to indicate a partial transformation into crystalline

phases of WO3'57 WO3H2058 and WO3.2H2059 (Figure 2.14(B) and 2.14(C)).

The only observable difference between the spectra from the two

extreme OB conditions is that there is a distinct peak 3.01 A

which is associated with WO3.2H20 for films without OB (Figure

2.14(B)). This peak is absent from the spectra from films deposited
-3
at P0 = 1 x 10- Torr (Figure 2.14(C)). On the other hand, there is

a peak at 5.37 A in Figure 2.14(C) (associated with WO3.H20) which is

absent in Figure 2.14(B). This seems to indicate that the WO3.2H20

phase was the dominant phase in the films without backfilling while

WO3.H20 dominated the films with OB.







(A)


(5)


L


cwo3
1-WO
WO- H20
*WO3 2H0


4
o
n

j I
W 'i~i


29 Idegreet


X-ray diffraction spectra: (A) As deposited, showing
amorphous pattern, and (B), (C) from WO3 films after
18 days in H SO solution where the amorphous back-
ground intensity was substracted to emphasize the
crystalline peaks. The spectra are from samples which
were deposited with (B) residual pressure of 1 x 10
Torr, (C) oxygen partial pressure of 1 x 10 Torr.
The symbols (o, +, *) show the origins of specific
diffraction peaks. (-: WO3, o: 03 .H20, +: WO 32H20).


Figure 2.14.











Electrical and electrochromic properties

The coloration times, tc, for films deposited at various pressure

are shown in Figure 2.15. The t increased from 4.2 sec. for films
c

prepared without OB to about 5.8 sec. for films prepared at P =
02
-3
1 x 103 Torr. The increase in the coloration times may be due to a

decrease in charge injection rate and/or a decrease in optical effi-

ciency. Figure 2.16 shows injected current versus time (j vs. t)

during coloration for films prepared at the two extreme OB conditions.

The films exhibited similar behavior i.e.,the current was almost

independent of t at short time (t < 0.2 sec.), and followed a t1/

dependence at long time (t > 2.0 sec.). However, the magnitude of

the current was lower for films with OB. The charge injected versus

time may be calculated from Figure 2.16 and is shown in Figure 2.17.

Again it demonstrates the lower charge injection at a specified

time for the high oxygen backfilled pressure. In Figure 2.18 the

electron diffusion coefficients (D ) which were measured by the

color front growth rate technique are plotted versus the partial

pressure of oxygen. The D was decreased by about a factor of three
-3
when the films were prepared at PO = 1 x 10 Torr, as compared to

the films without OB. The thermal activation energy (Eth ) for elec-

tron conduction in the uncolored films was also found to have been

changed by OB as shown in Figure 2.19. In this Figure, the tempera-

ture dependence of current is plotted versus T- for films prepared

at the two extreme OB conditions. From the slope of the plot Eth






































5.0
0
O








4.






1 o
10-5 10-4 10-3


P (Torr)



Figure 2.15. Time needed to color WO films to obtain 50% absorp
tion versus oxygen partial pressure during deposi-
tion. (E = -0.6 VSCE, Film thickness: 3,000 A).



























10



5




I



y 1
E = -0.6 VCE





0.5 1.0 5.0 10.0

Log Time (sec.)

Figure 2.16. Coloration current density versus time for WO3 films
prepared at two extreme conditions; solid circles:
'WO2.7' and open circles: 'WO2.9'.
2.7 2.9
























































8 10 12

Time (sec.)


Injected charge versus time for '.03
films prepared at two extreme conditions;
solid circles: 'WO and open circles:
'w .92.7
02.9


20






15

N



10



1


5






0


0 4




Figure 2.17.





























o 3.0 -








2.0
U













1.0









0-5 10--4 10- 3


P (Torr)
00
o
























Figure.2.18. Electron diffusion coefficient in WO3 films versus
S
- *


U .
[-1

0.--^-----------------------
10 5 0" l '

P Tor
3"

Fiue21.Eeto ifso cefcet nIO flsvru
pata rsueo xgndrn eoiin


























-6
10


10-7
10


2.5 3.0

1000/T ("Kl)


Figure 2.19.


Temperature dependence of current in WO films pre-
pared at two extreme conditions: solid circles:
'WO2 and open circles: 'WO .
2.7 7.9











can be calculated, i.e.,

I = I0 exp (-Eth/kT) (2.1)

where I is the current measured, k is Boltzman's constant and I0 is

the pre-exponential constant. For films without OB, Eth was 0.52 eV
th
but was increased to 0.67 eV for films prepared at Po = 1 x 10-

Torr. In Figure 2.20, the optical density change (OD) is plotted

as a function of the injected charge (Q) for the films deposited

at various oxygen pressure. The optical efficiency is defined as the

change in optical density per unit charge injected, and is the slope

of the plot AOD versus Q. The optical efficiency was decreased by

OB as shown in Figure 2.21. The absorption spectra for films prepared

at the two extreme OB conditions for the same thickness (2 zm) and

colored by the same injected charge (k6 mC/cm ) are shown in Figure

2.22. For ease of calculation, the absorption intensity and the

wavelength were converted into relative absorption coefficient and

photon energy, respectively. The decrease in optical efficiency by

OB is again apparent in these spectra since the relative absorption

coefficient (normalized to unity for 'WO ') at the absorption peak
2.7

of 'WO2.9' is only 65% of tof 'Wf 'WO2.7'. The energies of the

absorption maximum, Ep and the widths of the absorption band were

not changed significantly by OB, being 1.47 eV vs. 1.45 eV and 1.6

vs. 1.5 for 'WO2.7' and 'WO2.9', respectively.

The variation of bleaching time and color retention time

(open circuit memory) with oxygen backfilling are shown in Figure

























E = -0.6VE
SCE


0.8_


n-f


0.4.


5 10 15 20

Charge (mC/cm-)


Optical density change as a function
charge for WO films deposited with:
pressure of 1 10 Torr,_ B), (C),
partial pressure of 5 x 10 Torr, 5
and 1 x 10 Torr, respectively.


of injected
(A) residual
(D) oxygen
:10 Torr


0



Figure 2.20.


r





























40 L


0


P (Torr)
02


Optical efficiency of WO films versus oxygen
partial pressure during deposition.


Figure 2.21.


I I i







Energy (eV)


1.25


700 800 900


1,000


Wavelength (nm)


Absorption spectra for WO films prepared at two extreme
conditions. '--' represents '0 and '-
represents 'WO The films were of same thickness
(2.0 m) and colored with same injected charge ( -6
mC/cm ).


Figure 2.22.










2.23 and 2.24, respectively. The bleaching times were found to
-3
decrease from 4.7 sec. without OB to 1.6 sec. at P = x 103

Torr. After coloration, the power supply was disconnected and the

films would self-erase at different rates depending on the back-

filling pressure. The time for self-erasure to half of the original

OD decreased from 1,200 sec. for films prepared without OB
-3
to 460 sec. for films prepared at P0 = 1 x 10- Torr, as shown

in Figure 2.24.




Discussion

The corrosion of WO3 thin films in H2SO4 electrolyte proceeds

in the form of general dissolution and interfacial delamination.

Previously, it was speculated that the water in electrolyte and/or

WO3 films allowed or caused the corrosion. However, the mechanism

proposed could not fully explain corrosion in the acid-organic

solvent systems, nor was it compatible with data showing that WO3

thin films are oxygen-deficient and stable up to 350 C (crystalliza-

tion temperature). A new corrosion model will be proposed below.




Corrosion of WO, Thin Films
-------- -- ----



General dissolution

The thermodynamic equilibrium between crystalline WO and water

predicts that WO dissolves by hydrolysis to form tungstate ion;
3


ilt'l


























6.0





S4.0





" 2.0





0.0
10-5 10-4 10-3

P0 (Torr)




Figure 2.23. Time needed to bleach the WO films to 50% of
the original optical density versus oxygen
partial pressure during deposition.
(E: +0.6 V S Thickness: 3,000 A)
SCbL
































800L


u

0 600


C-.


2001
Sn-5 104 10-3


P (Torr)


Time for self-erasure to half of the original OD
of WO films versus oxygen partial pressure during
deposition.


Figure 2.24.


nnn


, UUU


-1


I


I


I











WO3 + H20 = WO2- + 2H (2.2)

log (W042-) = -14.05 + 2pH. (2.3)
-4 9
The solubility at 25 C is 10 moles/liter in neutral pure

water and decreases exponentially with decreasing pH. In acid

solution of pH 1 0.5 such as used in the current studies, the

solubility would be decreased to %10- moles/liter if this reaction

is still valid at such a low pH. The equilibrium potential of WO3

in this solution is 0.4 V with respect to the standard hydrogen

evolution potential (SHE).60 According to the Pourbaix diagram

for W (Figure 2.25),61 the experimental condition for storage

condition (pH = 0.5, E = 0.4 VSHE) is located in the WO3

stability region. That is, WO3 is the thermodynamically stable

specie under these storage condition. In contrast, the vapor

deposited WO3 thin films dissolved comparatively fast even in the

small volume of low pH solution used in the present tests. To

explain this discrepancy, a thorough understanding of structure and

composition of the film and of the thermodynamics of corrosion is

necessary.

Previous studies and present results have shown that thermally

evaporated WO3 films are

1. amorphous X-ray diffraction pattern of these films
exhibited broad diffuse peaks typical of amorphous
structure;

2. oxygen deficient XPS data also indicated that the
trngsten atoms exhibited various oxidation states of
w W and W due to the substoichiometry in these
films;

























-2 -1 C 2 3 i 6 7 'a 9 IC 1I :2 1. 5











S .
*,L I0 i.1
.0,L I












-lL :-i




-cs

-O Z 1 2 2 4 5 5 8 52 .2 5- E-






-,1 -


Fi Polc5nla oH r ul-lr lum i Iagram ionic senl; : 'C




Figure 2.25. Pourbaix Diagram for tungsten from Atlas of Electro-
chemical Equilibria in Aqueous Solutions, by M. Fourbaix,
Pergamon Press (1966). Reproduced with permission.








3. porous the relative density is approximately 50% to
80% of that of crystalline WO3;

4. hydrated films were reported to contain as much as 0.5
H20 per WO 3.

The water might have been incorporated during WO3 evaporation since

the partial pressure of water is the primary contributor to the mea-

sured total residual pressure.35 Water could also originate from

absorption of moisture from air after deposition since the films

were porous.

From the X-ray scattering experiments, Zeller and Beyeler39

obtained the radial distribution function of these films and

concluded that the amorphous WO films consist of a disordered

network of corner sharing W-0 octahedra. The substoichiometry and
4+ 5+ 6+
the existence of the W W and W oxidation states in these

films suggests that tungsten atoms may exist with a configuration not

only as W03 but also as WO2 and W205. Therefore, it is logical to

consider the evaporated WO3 films as a random solution of W02, W205

and WO3. These molecules are connected with each other through

corner sharing 0-bonds39 and forming a network-like structure (Figure

2.26). The water incorporated during evaporation or after deposition

is not expected to be important in structure formation, since

no observable structural change was observed at 170 OC when most
39
water was evolved.39 Water may be bonded in the open space between

tungstic oxide molecules by hydrogen bond or van der Waals forces.

The corrosion of these films in acid electrolyte can be

explained by examining the Pourbaix diagram for tungsten. Both









0




O HO
H


W 2 05
O/ \ 25


\
* 0


O --W .


0 W 'N
0 / 0

\ O 2 5


0---

/ 5


Figure 2.26. Proposed microstructure of vapor deposited WO3 films-
a random network composed of WO. W 0 and 0O units
linking each other through corner-snaring oxygen bonds.
Note water molecules may occupy the empty space between
tungstic oxide molecules.


0

0-^


/O


0---


/
WO










WO2 and W205 are not thermodynamically stable in the storage

condition (pH = 0.5, E = 0.4 V SHE). They should be converted to

higher oxidation states, but the conversion processes and products
38
are both unknown. However, Nazarenko et al. have shown that for
2+
dissolution of tungstic oxide in acid (pH < 1), the cationic, WO2
42-
form dominates. The anion forms, e.g.,WO2(OH)3 WO4 only dominate
36,62
at higher pH (> 2). Di Paola et al. 62 studied anodic WO films

on tungsten and proposed that the corrosion of these WO3 films took

place by hydration of WO3:

WO3 + 2H =WO2+ H20 (2.4)

WO22+ + (x I1)H20 = WO3.xH20 + 2H+ (2.5)

where x = 1 or 2.

However, the free energy of hydration for forming both the mono- and

di- hydrates was shown to be positive.63 The hydration reaction was

observed if electrochemically driven.362 Although the mechanisms

for corrosion of WO2 and W205 are unknown, it seems reasonable to

postulate similar reactions. For example, the dissolution processes

may be
2+
WO2 --. WO, 2 + 2e (2.6)

W205 + 2H+ -- 2WO22+ + H20 + 2e (2.7)

pH < 1.
7+
The solubility of WO2~ under these conditions is unknown, but the

ions as postulated precipitate as WO3 and its hydrates when the

solubility limits are reached. The total reaction can then be










written as

WO2 + (x+l)H20 = WO3.xH20 + H2 (2.8)

W205 + (2x+l)H20 = 2WO3.xH20 + H2 (2.9)

where x = 0, 1, or 2

and would result in the crystalline phases detected by X-ray diffra-

ction after 18 days storage in acid.

Even though XPS data (Figure 2.20) indicated the existence of
5+ 4+
high concentration of subvalence states (W5+ W) in the evaporated

W03 films, how much W02 and W205 does this represent? Gerard et al.64

calculated the concentration of each valence states (W4 W5+ and W6+

from XPS spectra of reactively sputtered WO3 with varied O/W ratio
6+
as shown in Figure 2.27. The concentration of W in a 'WO2.9

films was 80% and decreased to only 40% for a 'WO.5' film. It

is reasonable to expect that the same relation could be found in

evaporated WO3 films. For a film prepared without OB, the O/W
64
ratio was about 2.7. Based on the data from Gerard et al. the
6-
concentration of W6+ is expected to be < 60% with the remaining 40%
5+ 4+
tungsten atoms existing as W or W states. As discussed, the

WO films can be considered to be a random network consisting of

WO2' W205 and WO3 molecules linking each other through corner sharing

0-bonds. Then every one out of three tungsten links in the network

exists with a configuration of WO2 or W205. Since the oxygen

detected may have included that from the incorporated water, the

actual number of oxygen atoms that were directly bonded between














4+ 0+
w7 w


WO2.9


5 4 3 2 1 0


O2.5








01.,


Chemical Shift (eV)

6+ 5+ 4+ 0-
Concentration of W W, and W versus
corresponding chemical shifts for che sputtered
WO films with different 0/O ratio. From Gerard
et al., J. Appl. Phys. 48, 4253 (1977). Reprodu-
ced with permission.


Figure 2.27.









tungsten atoms might have been lower as discussed in the last section.

Therefore, the concentration of WO2 and W20 could be even higher than

proposed here. In the acidic environment, these WO2 and W205 links

are selectively attacked and the whole network solid may be broken

down into a distribution of random length of WO3 chains. These

broken WO3 chains presumably can no longer stay in the bulk but drift

into solution. These WO3 chains may be the polymer ions in the
34
solution observed by Arnoldussen.4 As a result, the dissolution

rate is quite high. Once the solubility limits in solution are

reached, precipitation occurs on surfaces as crystalline WO3 or

hydrates, since the free energy of hydration may be negative starting

with WO2 and W205.



Interfacial delamination

In addition to general dissolution, flaking by delamination at

the interface were also observed and may be induced by:

1. the internal stresses caused by precipitation of these cry-
stalline phases on the original sample;

2. gas evolution, e.g., 2H + 2e -- H2, accompanying the
dissolution and re-precipitation processes;

3. surface contamination and defects which weaken the adhesion
strength of the film to the substrate.

The first two mechanisms can be viewed as side effects of the general

dissolution mechanism proposed. It is not uncommon to find some

pinholes (t 10 um) in the evaporated WO3 films. In the corrosion

processes, the dissolved ion concentration in these pinholes will









reach saturation limit much earlier than over the flat surface and

lead to the rapid precipitation of crystallites of hydrated tungstic

oxide. These crystallites serve as nucleation sites for further

precipitation of crystalline phases, and eventually grow to a size

larger than the volume of the pinhole. The induced stresses from the

crystallites may be transmitted radially outward in the films and

along the film-substrate interfaces since the pinholes go completely

through the films. The rosette-shaped wrinkles in Figure 2.6(B) are

typical result of this effect. The delaminated films can be easily

broken into fragments by physical disturbance of the electrolyte

(Figure 2.7(A)). Small bubbles were observed on the samples during

corrosion and are assumed to be hydrogen gas evolved during the

corrosion process. The gas evolution may also enhance the inter-

facial delamination.

Interfacial delamination was rare for films deposited on graphite

substrates (semi-polished). The rough surfaces on these substrates

may provide a mechanical interlocking through pore and valley. Gas

bubbles were most often observed on films deposited on graphite sub-

strates, since the rims of graphite substrates provided sites for

cathodic reactions. However, gas evolution is not thought to be a

major cause for film delamination. This is particularly true for

films deposited on NESA glass substrates because little gas evolution

was observed.








Corrosion during cyclic coloration and bleaching

The reported increase in the dissolution rate for cyclic

coloration and bleaching condition can be explained by mixed potential

theory60'66 as shown in Figure 2.28. Since the WO3 specie in the

films should remain stable in this polarization range according to

the Pourbaix diagram, only WO2 and W205 species are involved in the

dissolution process. Figure 2.28 schematically illustrates the current/

voltage relations for the W02-H2SO4 system, but the W205-H 2S4 system

is expected to be similar. The two reactions occurring are assumed
22 -
to be WO2 dissolution (WO2 -* W02 + 2e ) and hydrogen evolution

(2H + 2e -- H2). The static corrosion potential and current for

this particular system is represented by the intersection between

the polarization curves for WOi dissolution and hydrogen evolution.

This represents the storage conditions discussed above. During the

bleaching pulse, the system is polarized anodically and the anodic

current density (IBW for W02 dissolution) is much higher than the

cathodic current density (IBH for hydrogen evolution). Thus more

rapid dissolution of WO2 (and W205) is expected during bleaching.

However, the WO2 dissolution current, ICW, is much lower than

hydrogen evolution current, ICH, during the coloration pulse. As

a result, WO2 (W205) dissolution is less but gas may be evolved

during coloration.




























I
I, | \.> a











I I I



BH CW ICH BW
Log Current Density

Figure 2.28. Combined activation polarization curves for WO2 disso-



nCr, and I are the overpotential, dissolution
current density of WO2 and reduction current of H
during bleaching pulse, respectively. While nC, ICW
and ICH are those for coloration pulse.
I I








In addition, Faughnan and Crandall observed an increase in

dissolution rate when device were cycled at higher contrast. Higher

contrast would require either operating at increased C/B voltage at the

same frequency,3 or an increased C/B time at the same voltage.3 As

discussed above, an increase in the dissolution rate can be expected

when the bleaching voltage is increased. On the other hand, Randin32

reported shorter lifetimes when the devices were cycled at low

frequencies (< 0.5 Hz). This low frequency effect is consistent

with the contrast effect reported by Faughnan and Crandall if the

higher contrast was obtained by prolonged C/B time at the same

voltage. The causes of increased dissolution rates at low frequency

can be attributed to a number of possibilities. For example, at higher

frequency, the short bleaching pulse may allow a sharp concentration
2+
gradient of WO2 to occur near the W03 film surface (due to limited

diffusion). This limits subsequent dissolution of WO,2 and hence

the kinetics of dissolution would be impeded.



Non-aqueous electrolyte

Attempts to use non-aqueous electrolytes have been reported by
32 33
Randin32 and Reichman et al.33 The results are summarized in Table

2.1. The WO3 thin films dissolved not only in sulfuric acid, but

also in formic acid and acetic acid, but remained stable in organic

solvents such as acetone, y-butyrolactone, and propylene carbonate.

The WO3 films dissolved when these organic solvents were combined

with sulfuric acid, e.g., methanol plus HSO, or acetonitrile plus












Table 2.1

Stability of vapor deposited WO films in
various non-aqueous solvents or electrolytes.


Electrolyte (solvent) Dissolution Reference


Formic acid Yes

Acetic acid Yes

Acetone No Randin32

y-butyrolactone No

Propylene carbonate No

32
Randin, 3Reichman
H2SO4/glycerin (1:10) Yes and Bard33


H2 SO/methanol Yes Reichman and Bard33

H SO acetonitrilee Yes


concentration unknown










H2SO The dissolution rates were not specified for these mixtures.

Randin32 reported the dissolution rate in H2SO/glycerin (1:10) was

an order of magnitude lower. These results indicate that in addition

to the substoichiometry of the films, the pH of the electrolyte is

also the culprit in corrosion. The water in the electrolyte can not
32-35
be the main factor as most investigators claimed, since a) the

corrosion does not occur until the organic solvents were combined with

acid, b) the dissolution rate in H2S04/glycerin was decreased by only

an order of magnitude while the water content in the electrolyte

was reduced by many (Q 20) orders of magnitude lower. The water

incorporated during evaporation is also unlikely to be the cause

of corrosion in non-aqueous electrolyte as Arnoldussen claimed,34

otherwise the corrosion should have been observed in acetone and

propylene carbonate. The lowered dissolution rate of WO3 film in
2+
H2SO4/glycerin could be caused by lowered mobility of WO2 ions

because of the high viscosity and/or by decreased proton concen-

tration in the electrolyte because of the decreased dissociation
67
constant. However, no conclusion can be made as to which cause is

dominant.










Modification of Stability of WO, Films by Oxygen Backfilling

Since the substoichiometry in the WO3 films is responsible for

the corrosion, an improved stability is expected at higher stoichiometry.

By backfilling during WO3 evaporation, the oxygen content of W03 films

was increased. The XPS data also indicated the concentration of W6+
5+ 4+
states in these films were increased relative to W and W Since

there are less dissolution-prone species available, the dissolu-

tion rate of oxygen enriched films in acid electrolyte should decrease.

As shown in Figure 2.28, the dissolution did decrease and became

negligible when the 0/W ratio reached 2.9. Interfacial delamination

was not observed as a result of the reduced corrosion rates (Figure

2.7b). However, the dissolution and precipitation of small amount

of WO2 and W205 still rearranged the surface portion of the films to

give crystalline peaks of WO3 and its hydrates in X-ray diffraction

pattern from OB samples (Figure 2.14c). The bulk of the films remained

intact when OB was used. As indicated by Figure 2.7b, only at higher

magnification can the corrosion effect (the slightly roughened

surface) be observed. The X-ray diffraction data also indicated that

WO .H20 was the dominant phase in oxygen enriched films while W03.2H20

was in the films without 03. This difference may be another reason

for lower dissolution in oxygen enriched films since WO3.H20 dissolves

less readily than WO3.2H20.68 The cause for inducing different

dominant phase by backfilling is unknown.










Modification of EC properties of WO3 Films by Oxygen Backfilling

The performance of electrochromic display devices is generally

determined by four electrochromic properties: 1. coloration speed

2. optical efficiency, 3. bleaching speed, and 4. coloration

retention time (self-erasure rate). The effect of oxygen backfilling

on the EC display performance will be discussed in this order.



Coloration speed

The rate controlling mechanism for coloration was proposed30 3

to be current limited by the WO3 resistance at short time. The

duration of WO3 resistance control increases as the thickness of WO3 was

increased. In addition, the current was affected by the Schottky and

Helmholtz double layer (HDL) barriers. At long time ( >1 sec.) for

thickness of <0.4 fm, the-current control was dominated by the HDL.

As shown in Figure 2.16, the coloration current for films prepared

with or without oxygen backfilling showed a time independent behavior
-1/2
at the onset of coloration, and approached a t dependence at

t >3 sec. This t-1/2 dependence is characteristic of HDL control.30

In general, the charge injection process is not affected by OB.

However, the magnitude of injected current was lowered for oxygen-

enriched films. This can be explained by considering the Butler-

Volmer equation:60
Volmer equation:


S=(l-O eI p -(enxp 0
JP = J [ exp T exp ]
0 T kT


(2.10)









where J is the proton current density across the double layer, J0

is the exchange current density, B is the barrier symmetry factor,

k is Boltzman's constant, T is temperature, e is electronic charge

and n is the overpotential (the potential drop across the HDL when
60
Jp # 0). Since B is a property of the ionic solution, it is not

expected to be affected by changes in the film properties. The

overpotential, n, can be expressed as

nu(x)
= a NF JRf (2.11)

Au(x)
where V is the applied potential, is the reduction in chemical
a NF
potential due to the proton concentration (x) in the W03, N is Avogadro's

number, F is Faraday's constant and Rf is the film resistivity. As

shown in Figure 2.18, the electron diffusion coefficient in WO3 was

found to be decreased by OB. The conductivity (resistivity) measured

at room temperature was also found to be decreased (increased) by a

factor of three (Figure 2.19). Since Rf is higher at the onset of

the coloration (x = 0), n would be smaller for oxygen enriched films

at the same applied bias V The current injected would be lower as
a

indicated by equation 2.13.

Resistance control was dominant at t < 1-2 sec. However, the

effect on the magnitude of current injected persisted into the HDL

control region, i.e., the decrease in current by OB was almost

invariant with time even in the HDL controlling region (Figure 2.16).

However, the optical efficiency, i.e., the change in optical density








per unit charge injected, was also decreased by OB as shown in Figure

2.21. This effect together with the lower current injection led to a

slower coloration speed as shown in Figure 2.15. The mechanism of

degraded optical efficiency is discussed in the following section.



Optical efficiency

An understanding of the electronic structure of WO3 thin films is

necessary before discussing the change in optical efficiency by oxygen

backfilling. In the following, the mechanisms proposed for optical

absorption and DC conduction are reviewed and compared with the

present results.



Electronic structure for WO thin films. Many models have been

proposed to explain the optical transition in electrochromic WO films.

Optical absorption in these films is known to arise from a transition

of an electron from one trapping site to an adjacent site. A trapping

site is a position in the lattice where localization of an electron

will result in a lowering of the total free energy of the system. The

trapping sites are probably W sites as identified by electron spin

resonance studies.69 Faughnan and Crandall proposed an inter-valency

charge transition model which can be expressed as shown below:


W5+(A) + W (B) + hu -- W6(A) + W5 (B) (2.12)

where A and B represent two neighboring sites of tungsten atom,

and hv is the photon energy. This model explains qualitatively ehe

features of the optical absorption including the magnitude of oscillator










strength. However, the XPS data70 have shown the presence of large

concentration of W4+ and W +( fl0%) even in the transparent WO3 thin

films and hence appear to disagree with this model.

An alternative model based on small-polaron theory has been

proposed by Shirmer et al.0-12 In a systematic study of the optical

properties of amorphous and crystalline WO3 films, they found a

shift in the maximum absorption peak from -900 nm to n1,200 nm upon

crystallization. This shift, plus the asymmetry of the high-energy

side of the absorption maximum (Figure 2.22) in amorphous films has

led Shirmer et al. to postulate a self-trapping term (i.e.,a polaron

formation) which would lower the energy of the site with the trapped

electron relative to other W6+ sites. In a disordered system, the

W sites do not all have equivalent energy because of the existence

of a range of W-0 bond lengths and angles.71 The injected electrons

will be trapped primarily at those sites of lower energy. Hence,

optical transitions will occur with a range of energies. This explains

the asymmetry in the optical-absorption peak. The shift to lower

energy for crystalline films occurs because the disorder term

disappear for crystalline films.

The optically excited polaron hopping between two neighboring

sites can be described by a diagram showing the electron potential

energy as a function of a single, one-dimensional lattice configu-

ration coordinate (Figure 2.29). It has long been realized that




























E
Sop





Eth



qA qB

Configurational Coordinate


Figure 2.29. Configurational Coordinate model for the polaron
hopping absorption in the WO3 films.










optically induced transitions between two sites can be considered as
72
Frank-Condon transitions, indicated by the vertical arrow in

Figure 2.29. These processes take energy E which can be measured
op
from the peak energy of the absorption band. The change of the

lattice distortion accompanied by the electron transfer from site'A

to site B is represented by Aq = qA q. The energy Eth is the

thermal energy of the small polaron hopping from site A to site B,

which can be obtained from the measurement of conductivity change

over a temperature variation as discussed below.

The DC conductivity of WO3 amorphous films have been reported

by several authors. Faughnan et al.7 reported a log T-/4

dependence at low coloration (x < 0.3 for H WO ) between 10.2 and
x 3

300 0K, and concluded that the electron transport in this condition

was best described by a variable-range-hopping model (electron
74 75
hopping between trapping sites). Benci et al.74,75 in a similar

study reported a logI T/7 dependence for transparent film
,-l
between 4.2 and 300 "K, although a logI T depencence could

described the data at temperature near 300 K. They concluded the

conduction in transparent WO3 films at T<300 "K is in agreement

with the hopping between two localized states (a donor state and a
76
trap state). Gritsenko et al. studied the DC conductivity of WO3

films at a higher temperature range, 353 773 K. and reported a

logic T-1 dependence and concluded that the overlapping of the

potential of deep centers was responsible for the conductivity.









The mechanism of electron transport between these centers was

thermally assisted tunneling and the conductivity increased exponen-

tially with temperature. Although the detailed nature of these

"donor states" 7475 and "deep centers" was not explained, they are

expected to be associated with the active polaron states where a

localized electron is allowed to contribute to conduction. The

exact nature of active polarons will be discussed in a later section.

It seems, however, that the DC conduction mechanisms of electrons

in WO3 films vary with the temperature range investigated. The

mechanism is associated with polaron hopping at low temperature

(T < 300 K), and with polaron tunneling at high temperature (353 <

T < 773 K). Hence, the thermal activation energy Eth obtained at

high temperature will not have the same physical significance as

that from low temperature measurement. One needs to be cautious

in interpreting data to identify the controlling mechanism.



Effect of oxygen backfilling on EC characteristics and

electronic structure of WO. films. As discussed, the optical

efficiency (or the optical absorption per unit injected charge),

the DC conductivity and the electron diffusion coefficient of 'O3

films were decreased by oxygen enrichment, while the thermal acti-

vacion energy (between 300 450 oK) was increased. The decreased

optical absorption for more oxygen enrichment was also reported by










34
Arnoldussen.34 He observed that oxygen ion implantation of

evaporated WO3 films resulted in films which could no longer be

colored. This is the extreme case. He attributed this phenomenon

to the creation of electron traps by structural change as a result

of oxygen implantation. However, no detailed discussion on the

nature of electron traps or the mechanisms causing decreased optical

absorption was given.

In the discussion of optical properties of the trapped or

localized center as in WO3 thin film, the Smakula's7 equation is

applicable


N'f = 8.7 x 106 n (K L)W (2.13)
(n2 2)2 max 1/2

2
where N' is the number of absorption centers per cm n is the index

of refraction of the film, K is the absorption coefficient at
max
-1
the band peak in cm L is the thickness of the film and W /2 is the

full width at half maximum (FWHM) in eV. The index of refraction, n,

does not vary significantly between films deposited from WO~ or WO3

powders (which were assumed to result in a different 0/W78 ratio) or
9
between amorphous films and single crystal WO The absorption

spectra in Figure 2.22 are from film of the same thickness and

injected charge. As reported earlier, the width of the absorption

bands (W1/2) are nearly the same (1.6 eV for films prepared without
-3
OB, 'WO 1.5 eV for films prepared at P = 1 x 10 Torr,
2.7

'WO 9'). Thus, equation 2.12 can be written as:

K = Constant N' f (2.1.)
max










The analysis of degradation in optical efficiency therefore, can be

initiated by either of the following two hypotheses:

1. The charge in K was caused by the change in f, while N'
max
remained constant at constant charge injection (i.e.,every
electron injected resulted in the formation of a polaron).
The decrease in the absorption coefficient was therefore
caused by a decrease in the oscillator strength of the
absorption center.

2. The change in K was caused by the change in N', while f
max
remained constant (i.e., the oscillator strength of absorp-
tion center remained unchanged by the backfilling). The
decrease in absorption coefficient was caused by a decrease
in the concentration of absorption centers, possibly by
creation of electron traps by OB.
78
The first hypothesis was adopted by Yoshimura et al. to

explain the increase in optical efficiency of tungstic oxide films

deposited from crystalline WO2 powder. They observed the absorption

coefficient K was increased for films deposited from WO2 powder
max

as compared with those from WO3 powder (for the same film thickness.

injected charge and nearly equal FWHM's of the absorption peak). By

assuming every injected electron resulted in an absorption center

(polaron) and the use of equation 2.13, they concluded that the

oscillator strength, f, of films deposited from WO, powder was

higher than that of films deposited from WO3 powder. They also

concluded that the increased oscillator strength was caused by an

increase in the degree of extension of electron wave function for

films from WO2 powder. Using a configurational coordinate model,

they suggested the change of lattice distortion iq was decreased as

a result of an increase in electron wave function extension and










predicted a decrease in both Eop and Eth for films deposited

from WO2 powder versus those from WO3 powder. The predictions

were consistent with experimental observations. Therefore, they

concluded that the decrease in optical efficiency for films

deposited from WO3 powders was caused by a decrease in oscillator

strength of the absorption centers. The decrease in f was caused by

a change in the lattice configuration. Although they did not measure

the 0/W ratio in these films, they speculated that the change in

lattice condition was introduced by the deficiency of oxygen in films

deposited from WO2 powder.

At first thought it would seem that the result in the present

study, i.e., the decrease in optical efficiency in 'WO2.9' films can

be explained by a decrease in oscillator strength, 'f', using the

model of Yoshimura et al. The decrease in f was in turn caused by

a decrease in the degree of electron wave function extension caused

by oxygen enrichment. However, an increase in E was not observed.
op

The increase in E might not have been measurable due to a number
op

of factors that affect such measurements. For example, E was
op

reported to increase in energy as the optical density was increased

(as much as 0.15 eV).9 However, the optical density at constant

charge injection decreased for 'WO2.9' versus 'WO2.7'. Since data

were taken at constant charge injection, the configurational coordinate

model predicts a higher Eop for 'WO2.9' films, but the lower optical
op 2.9











density may have resulted in the wavelength of absorption peaks being

shifted lower and fortuitously coinciding with that of 'WO2 '
2.7

On the other hand, the derivation of the model by Yoshimura et al.

was based on a number of assumptions. One assumption was that the

concentration of injected electrons was equal to that of absorption

center. As will be shown later, there are a number of tungsten

sites where electrons may become highly localized and are not able

to cuase optical absorption or contribute to electronic conduction.

Therefore, the concentration of the electrons which can participate

in optical absorption is expected to be lower than that measured by

current integration.

A second assumption was that the DC conduction mechanism was

assumed to be associated with polaron hopping even for the tempera-

ture range 300 500 K, where the measurements were performed.

As discussed, the DC conduction mechanisms of electrons in WO3 films

varied with the temperature range investigated. The electronic

conduction in the studies of Yoshimura et al. was expected to be

controlled by thermally assisted tunneling as opposed to the hopping

they assumed. In addition, the polaron hopping mechanism predicts that
10
Eth = 1/4 E while their measured relationship was E o, 2 Eth
th op op th
(i.e., Eth versus Eop were 0.75 eV and 0.55 eV versus 1.4 eV and 1.2

eV for films deposited from WO3 and WO2 powders, respectively).

Therefore, the Eth Yoshimura et al. measured does not have the same
th










physical significance as the Eth in the model they proposed (i.e., the

polaron hopping model). The prediction by their model of an increase

in E which was not observed in this study further indicating that
op

the model may be invalid and that a different approach is necessary.

The decrease in the absorption coefficient K of 'WO '
max 2.9

films (for the same charge injection), therefore, could have

been cuased by a decrease in the concentration of optical abosrption

cneters instead of a decrease in oscillator strength. However, the

decrease in the concentration of absorption centers must be associated

with an increase in the concentration of inactive electron traps as

a result of OB.

The existence of inactive trapping sites have been indicated by a

number of authors. Gritsenko et al.76 reported that the calculated

electrically active "deep center" concentration was two orders of magni-

tude lower than that of injected charge in WO3. They speculated that

not all injected charge formed "deep center". However, no explanation

for such behavior was given. In addition, most plots of optical density
35
versus injected charge (e.g., Figure 2.20 and reported data ) indicated

that there is a threshold charge (- 2 mC/cm 2) needed before coloration

can be observed. The threshold charge strongly suggests the presence

of inactive trapping sites. Thus, the decreased absorption center

density could be explained by an active-inactive polaron site model.

The electrons that are trapped at inactive sites will become highly

localized, and the possibility for such electrons to re-escape and










contribute to absorption or electron conduction is very small

compared to active polaron hopping. The energies of such inactive

sites are expected to be lower than those of active sites. Electrons

will preferentially occupy the lower energy states.

The origin of such inactive centers can be correlated with the

increase in water concentration or the decrease in W4 and W5+ by

oxygen backfilling. Deneuville et al. have proposed that cation

sites are improtant to coloration and that there are two types

of sites in WO3 films which protons can occupy. One type is

optically inactive, i.e., the proton does not contribute to colora-

tion. Protons incorporated during evaporation occupy these sites and

the as-deposited films are transparent. The other site is optically

active; coloration may occur when the protons in the optically inactive

sites are redistributed into active sites by changes in the chemical

potential of H within WO3. This in turn may be caused by changes in

the chemical potential of H in the hydrated dielectric electrolyte
79
as found for W03/MgF2/Au devices, or from the absorption of photons
80
as in ultra-violet irradiation. The essentially invariant change

in the H/W ratio in the films before and after coloration (reported

by Deneuville et al. to be a 3% increase) was used to support this

postulate of two sites. However, this model can not be used

verbatim to explain electrochromism since electrically-induced

coloration is accompanied with a large quantity of charge passed

through the cell and an opposite amount of charge passed in bleaching.

This charge is tens of mC/cm2 for deep coloration, whereas in the









redistribution model, the only charge transported could be that to
2 65
charge the double-layer capacitance, i.e., '20 mC/cm Should the

charge be going into some other electrochemical side reaction. it

is extremely unlikely that such a reaction would be completely

reversible while having such a large capacity.

The idea of two types of cation sites is still acceptable, and

supported by the observed aging effect in WO 3/Li device reported

by Knowles.4 The aging consisted of a gradual decrease in optical

contrast obtained during a fixed time potentiostatic coloration

pulse. It was reported65 that some Li was present but did not make

a contribution to coloration in the bleached state. This residual

Li content increased with both the number of C/B cycles and the

depth of coloration, and was dependent on the degree of hydration

of the WO3 films. The dependence of residual Li content on the

degree of hydration of the films was postulated to result from the

evaporated W03 film behaving as an ion-exchange material. The H

ion of an OH- group within the film was postulated to exchange in

a colored film for an Li+ ion from either the electrolyte or from

solid solution within the WO3. In the latter case, a film colored

only by Li+ insertion would then be bleached by extraction of the

same total amount of charge comprised of both Li+ and H leaving a

residual concentration of Li+. This would limit the rate of further

injection and result in the observed aging effect. Knowles14

reported that a solution to this problem was ultra-violet illumination


/i











of WO3 films held in as-bleached state. As a result of UV

illumination, the aged cell could be restored to close to its

initial performance.

Therefore, it seems that the idea of active-inactive cation

site can be related to the active-inactive polaron site model

proposed here. The inactive cation incorporated during deposition

might have introduced an extra lattice distortion on the local

arrangement of W-0 entities. The extra lattice distortion might

cause'an increase in the self-trapping energy for the trapped electron

by Coulumbic interaction.781 The electron trapped at such an

inactive site would require much higher energy to escape and the

possibility for such electron to contribute to conduction and

optical absorption would be low. Since the water concentration was

increased by OB, there were more inactive protons incorporated and

hence an increase in inacitve electron trapping sites.

On the other hand, the electrons that were localized around W
4+ 5+
sites (W 4 W states) in an as-deposited transparent film might cause

a decrease in the self-trapping energy for the electrons trapped

at neighboring W6+ sites (e.g..by Coulumbic interactions). Therefore,

the electrons localized at these neighboring sites might have a higher
5+
tendency to escape and behave as active polaron. For the W sites
t+ 5+
that were not adjacent to a W or W the electrons that were

localized might self-trap in a deeper well and act as the inactive
4+ 5+
polaron. Therefore, when the concentration of W and W was

decreased by OB, there were less active polarons formed upon










electron injection. As a result, the conductivity of the oxygen

enriched films was lowered. The increased thermal activation energy

in oxygen-enriched films might be caused by an increase in lattice

distortion as shown in configurational coordinate model. The lattice

distortion was increased by the decrease in active polaron

density, since the average distance between polarons was increased

and hence an increase in the lattice distortion would be required

for polaron motion between the sites.

In conclusion, the decrease in optical efficiency in oxygen

enriched films is best described as the result of an increase

in the active trap density by OB. However, the oscillator

strength of absorption center might have been decreased by oxygen

enrichment; present data are not sufficient to preclude a decrease

in oscillator strength but do show that the decreased optical

efficiency is not a result of changing only the oscillAtor

strength. While reasons for optical inactivity of selected

sites are uncertain, it may be associated with the lower energy

of the site and/or increased lattice distortion in the configurational

coordinate model. Increased distortion in this model could
4+ 5+
result from the presence of W W and protons.










Bleaching speed

The bleaching has been reported to be limited by proton
35
diffusion in a space-charged region. The current as a function

of time is described by


I(t) = (p3k0 p 1/4Val/2(4t)-3/4 (2.15)


where p is the volume charge density of protons (equal to proton

charge per unit area multiplied by thickness of the film), k is the
-12
relative dielectric constant of WO3 (g0 = 8.85 x 10 Farads/m),

ip is the proton mobility and Va is the applied voltage. The time

for complete erasure is


tf = pL4 / 4kE pVa2 (2.16)


where L is the thickness of the film. If the thickness of the films,

the applied voltage and the injected charge density are all kept

constant, tf is inversely proportional to the proton mobility, p .

There are two major factors which affect the diffusion of protons

in the WO3 thin films:

1. amorphous structure crystallization of the films is
known to decrease the g~ogn diffusion coefficient by
an order of magnitude;

2. water content higher wate53content in the films resulted
in faster proton diffusion.

The microstructure of oxygen-enriched WO3 films remained amorphous,

but the water content was higher in these films as shown in Figure

2.11. Thus, proton mobility is expected to be higher. In addition,










the porosity in the oxygen enriched films was significantly increased

(Figure 2.13). The increased surface area and decreased diffusion

length caused by porosity will also facilitate the bleaching

process. As shown in Figure 2.23, the bleaching time was reduced
-3
by a factor of three for the films prepared at P0 = 1 x 10 Torr

as compared to the films without backfilling.



Self-erasure

The effect of increased proton mobility and reduced diffusion

length is also reflected in the self-erasure rate of OB films.

Hitchman82 has proposed that self-erasure is caused by the oxidation

of the so-called 'hydrogen tungsten bronze' HWO3 (which can also be

viewed as a solid solution of protons in WO3) back to WO3 according

to the reaction:


2HWO3 + 1/202 = 2WO3 + H20 (2.17)


This is, however, a general description when the process is considered

from a thermodynamically point of view, e.g., it could be written as


2H1 + 2e + 1/202 -- 2H20 (2.18)


where the H+ and the e are both removed from WO3 films. In

reality, the injected protons remained ionized and the colored WO3
35
films contained a large concentration of proton. A chemical

potential is developed by the proton in the films resulting in a









back emf of 0.7 V for an OD of 1.0,83 and this is the driving force

for the self-erasure process. Therefore, similar to bleaching, self-

erasure is controlled by concentration-dependent proton diffusion in

WO3 films. Since the proton mobility was increased and the

diffusion path was reduced in the oxygen enriched films, the

coloration retention time was decreased by approximately a

factor of three (Figure 2.24). The magnitude of the reduction

was the same for bleaching and self-erasure. This again indi-

cates that proton transport was the controlling mechanism in both

processes and was modified by the oxygen backfilling.



Summary

1. Vapor deposited WO3 films have been confirmed to be amorphous,

oxygen deficient (O/W = 2.76), and porous (%80% of bulk density).

2. X-ray photoelectron spectroscopy data have confirmed that the

tungsten atoms in the as-deposited films existed not only
6+ 4+ 5+
as W but also as W and W5. The total concentration of

subvalence states can be %40% or higher.

3. The tungsten atoms in the as-deposited films are postulated to be

present as WO3 W205 and WO2 according to their valence states.

The films structure can be visualized as a random network

composed of WO and lower valent oxide molecules linking with

each other through corner sharing oxygen bond.

4. The Pourbaix diagram for W indicates that WO3 is the thermo-

dynamically stable specie for WO3 films stored in the acid










electrolyte (pH = 0.5, E = 0.4 VSHE), while WO2 and W205 are not.
SHn' 5 Z

The conversion of WO2 and W205 into WO3 were observed as dissolu-

tion of WO2 adn W 205 in acid to form crystalline WO3 and its

hydrates (i.e., W03.H20 and WO3.2H20). The reactions are
2+
uncertain, but appear to involve the formation of WO ions in

solution and saturation of these solution.

5. The dissolution of as seposited films was rapid ( 116 A/hr.) due

to the presence of these unstable species. In addition to

the general dissolution, interracial delamination between

the film and substrate was observed. The precipitation of

crystalline WO and hydrates on the original film, particularly

in the pinholes is postulated to cause delamination by intro-

ducing internal stresses which exceed interfacial adhesion

strength between film and substrate. The hydrogen evolution

accompanying the corrosion processes might also contribute to

the delamination.

6. Corrosion was increased when films were cycled between colored

and bleached states because the dissolution rates of WO2 and W205

were greatly increased by anodic polarization during bleaching.

The corrosion was still observed when the H20 in the acid electro-

lyte was replaced by glycerin indicating pH rather than HO content

of the electrolyte defined whether corrosion would occur.

7. The oxygen content in WO films was increased by oxygen back-
filling. The concentration of subvalence sates (4+ 5+
filling. The concentration of subvalence states (W W










decreased with respect to W6+ states. The stability of the films

in acid was improved becuase of decreased concentrations of

dissolution-prone species in the films. The dissolution and

delamination was negligible for films prepared at PO = 1 x 10

Torr, (O/W = 2.9).

8. The coloration speed was slower for oxygen enriched films

because of lower current injection and reduced optical efficiency.

The current injection was found to be controlled by the same

mechanisms as for films without backfilling (i.e.,WO3 resistivity

and Helmholtz double layer). However, the magnitude of the

current passed was lower because the effective potential across the

HDL was decreased by the increased film resitivity from back-

filling. The increased resistivity and decreased optical

efficiency in the oxygen enriched films was explained by postu-

lating an increased density of inactive electron trapping sites.

This increased density of inactive trapping sites limited

the formation of active polaron for optical absorption and DC

conductivity. However, the possibility that a decreased oscilla-

tor strength resulted in a degraded optical efficiency could not

be ruled out.

9. The porosity of the films was increased by deposition at high

background pressure. Due to increased effective surface area,

the films absorbed higher amounts of water after deposition as

shown by secondary ion mass sepctrometry. The bleaching speed











and self-erasure rates were increased since proton mobility was

increased and the diffusion path was shortened by this porosity

effect.














CHAPTER III
DEVICE FABRICATION



Introduction

In Chapter II, the corrosion of EC display system based

on WO3-H2SO /H20 was investigated. A solution to corrosion,

viz., oxygen backfilling, was studied and shown to increase the device

lifetime. However, the improvement was accompanied with degredation

of EC properties.

In this chapter, modification of the electrolyte to increase

the device lifetime is considered. Two types of ECD systems were

investigated and fabricated. The first type involved the use of a

liquid electrolyte consisting of LiCIO4 dissolved in propylene

carbonate (Li/PC). Since the devices using this electrolyte were

reported to exhibit slow switching speeds,15-17 two approaches were

adopted to increase the speed, viz., nitrogen-backfilling and MgF2 over-

layer. The switching speed was found to be increased by depositing the

WO3 films at high partial pressure of nitrogen and by depositing a thick,

dense MgF2 overlayer. The second device configuration involved the

use of a hydrated solid dielectric layer, i.e., a MgF2 film deposited

with air-backfilling. The results indicated that air backfilling

during MgF2 film deposition was effective in achieving fast

switching, while air backfilling during deposition of WO3










films limited the operational area of the device and may not be

desirable. The mechanisms will be proposed to explain these results.



Experimental



Preparation

WO3 films were prepared by thermal evaporation of pure WO3

polycrystalline powder (Cerac, Inc., 99.9%) from a resistively

heated source onto indium-tin-oxide (ITO) coated glass. The source

used was an alumina coated tungsten basket (Sylvania CS-1010). The

surface resistance of the ITO coated glasses ranged from 15 0/ to

100 0/ However, during the experiment, the resistance of the

ITO within the group of the samples to be analyzed was kept constant.

The deposition rate was varied between 3 A/s and 30 A/s using a

thickness monitor (Sloan DDC1000) and manually adjusting the power

to the evaporator. Before evaporation, the chamber was evacuated
-6
to a base pressure of 6 x 106 Torr. The evaporation was accomplished

either with a partial pressure of air or nitrogen controlled by a

gas flow controller (Vacuum General) or directly without gas back-

filling. Thin films of MgF2 were prepared by thermal evaporation

MgF2 powder from Ta boat onto WO3 coated ITO glass. Deposition rate

was kept between 5-10 A/s. The pressure during the deposition was

varied between 4 x 10-6 and 5 x 104 Torr of air or nitrogen. Thin

gold films were prepared by DC sputter deposition from a small coating











system for SEM sample preparation, or by vapor deposition from an

electron beam deposition system. The substrates were cleaned by

rinsing with soapy water (Alconox with water) for 15 minutes

followed by hand scrubbing with cotton wool. The substrates were

then rinsed in deionized water for approximately 10 minutes,

isopropyl alcohol (IPA) for one minute, and vapor degreased for

5 minutes in a IPA/Freon tank. The liquid electrolyte was prepared

by dissolving LiCIO1 in propylene carbonate (PC) to make a 1 M

solution. Neither LiCIO4 nor PC was processed to eliminate

moisture before or after mixing.



Characterization

The relative density of the evaporated WO3 films was changed by

nitrogen backfilling and was measured by the same method reported in

Chapter II, i.e. mass change versus surface step height. The

surface topography of the films was examined by scanning electron

microscopy (Cambridge XSEM-9). For Li/PC EC devices, the glass

substrates deposited with WO3 films or WO3/MgF2 films were assembled

into an EC cell as shown in Figure 2.3. The coloration/bleaching

(C/B) speeds were measured with the apparatus shown in Figure 3.1.

A gold wire loop was used as the counter electrode. A constant

potential, 2.75 V between SnO2 and Au electrodes, was used to

drive the cell (Heathkit IP-28). An Ar-Ne laser (633 nm) was used

in the measurements. The operating area was 3.88 cm The solid

























---- -\ o ^ Cons c a-,
VloL-age
~ 5uppli

?------


:C- e ce


Si- 3.soIay ::e::~~c


Figure 3.1. Electrochromic cell circuit diagram for
the measurements of C/B speeds.


,,











state devices were fabricated by successive deposition of WO3,

MgF2 and Au films as described above (Figure 3.2). The thickness

of WO3, MgF2 and Au was 4.0 KA, 1.5 KA and 0.2 KA, respectively.

The C/B speed was measured in the reflection mode.



Liquid Electrolyte Device



Results

As shown on Table 3.1, the relative density of the WO thin

films increased as the deposition rate (DR) increased (for a fixed

partial pressure of nitrogen), or increased with decreasing partial

pressure of nitrogen (PN ) for the same DR. As shown in Figure 3.3,
2-4
the films prepared at high background pressure (PN = 7 x 10 Torr)

exhibited a more rough surface topography than those prepared at

low pressure. The films prepared without backfilling were feature-

less, even at a magnification of x 20,000.

The effect of deposition rate, on the coloration time (t ) for
c

2,000 A thick WO films to a contrast of 30% absorption are shown in

Figure 3.4. The t remained nearly constant up to a DR of 10 A/s

then increased drastically for a DR of 30 A/s. The data in the same

figure also show that t decreased as the partial pressure increased
-5 -4
from the residual pressure (1 x 10-5 Torr) to PN2 = 7 x 10-4 Torr.

This effect of backfilling pressure on tc is again illustrated in

Figure 3.5 for films deposited at the same DR (5 A/s) at three

















ITO























gF,. (1.5 KA)








Figure 3.2. Schematic representation of the ITO/
WO3/ gF2/Au EC device.











Table 3.1


Effect of nitrogen partial pressure and deposition rate
on the relative density of WO3 film


There are several possibilities for the relative densities of
the WO films to be greater than one; (1) Calibration errors,
(2) The presence of lower oxides, e.g., WO and W205 (bulk
density of WO2 is 12.1 g/cm ).





















































(B)

Figure 3.3. Scanning electron micrographs of WO films showing
the effect of nitrogen backfilling during deposition
on the surface topography of the films;
(A) P = 1 x 10 Torr, x 20,000,
(B) Pres= 7 x 10 Torr, x 20,000.
N2


















60 -



50 -




40

E-
















3 5 10 30

Deposition Rate (A/S)

Figure 3.4. Effect of deposition rate variation on coloration time for
films prepared at two extreme conditions. solid circles:
residual pressure of 1 x 0- Tort and_4open circles:

partial pressure of nitrogen at 7 x 10 Torr. The films
were colored to a contrast of 30% absorption (E = -2.75 V
on WO03). Substrates used were high resistance ITO coated
glasses (75
0

o 20 -









3 5 10 30

Deposition Rate (A/S)

Figure 3.4. Effect of deposition rate variation on coloration time for
films prepared at two extreme conditions. solid circles:
residual pressure of 1 x 10 Torr and ,open circles:
partial pressure of nitrogen at 7 x 10 Torr. The films
were colored to a contrast of 30% absorption (E = -2.75 V
on WO3). Substrates used were high resistance ITO coated
glasses (75 2/a ) .




















12




10



8







4
4 -
o



2





2 x 104 7 x 100 2 x 103

PN (Torr)



Figure 3.5. Effect of variation of nitrogen partial pressure during
deposition on coloration time for WO films colored to
different contrast: solid circles: 4% and open
circles: 30% absorption (E 2.75 V on WO ). Substra-
tes used were low resistance ITO glasses ( 15 2/0).




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