University of California Santa...

University of California

Santa Barbara

Integration of Nanostructured Titania into Microsystems

A Dissertation submitted in partial satisfaction of the

requirements for the degree

Doctor of Philosophy in Materials

by

Zuruzi Abu Samah

Committee in charge:

Professor Noel C. MacDonald, Chair

Professor Kimberly L. Turner

Professor Cyrus R. Safinya

Professor Jacob Israelachvili

Professor Anthony G. Evans

June 2005

i

The dissertation of Zuruzi Abu Samah is approved.

Kimberly L. Turner

Cyrus R. Safinya

Jacob Israelachvili

Anthony G. Evans

Noel C. MacDonald, Committee Chair

June 2005

ii


Copyright © 2005

By

Zuruzi Abu Samah

iii

For

My Family

and

In Loving Memory of Mansur Omar

iv

ACKNOWLEDGMENTS

I thank Professor Noel MacDonald for the opportunity to come to Santa

Barbara for graduate studies and the freedom to pursue my research

interests and ideas. My education at UCSB has made me a better engineer

and, more importantly, a better person. Noel has enriched my life

tremendously and for that I am grateful.

I am fortunate to be able to interact with Professor Cyrus Safinya. I thank

Cyrus for access to cell-culture facilities and excellent X-ray tools. I had an

enriching experience interacting with the Safinya Lab, learning and

participating in biomaterials research. I thank Professor Anthony Evans for his

insights and helping me crack problems of my research. Tony has a way to

relieving stressful issues and I learn just by talking to him. I thank Professors

Jacob Israelachvili and Kimberly Turner for serving on my qualifying and

dissertation committees. I thank Professor David Clarke for proof reading the

chapter on oxidation in my dissertation.

Professors Martin Moskovits and Andrei Kolmakov are very much

appreciated for collaborating in sensor research, for sharing their insights and

educating me. Blaine Butler taught me cell-culture and collaborated with me

when I first started research. I enjoyed working with her and I wish her well. I

thank Marcus Ward for working together on nanocomposites and for delightful

conversations between symposiums in San Francisco. Diana DeRosa is

appreciated for spending a summer collaborating on Ti oxidation and sharing

v

her talents with me. She will make an outstanding engineer or medical doctor

one of these days, unless she decides to be a dancer.

There are numerous others who made my stay at UCSB meaningful and

contributed to my education. Listing them all is not possible. They include

friends, numerous clean room staff, administrative personnel in the Materials

and Mechanical Engineering departments and staff of the Microscopy and X-

Ray Labs of the MRL.

Relieving stress from a rough day at work is easy when one has great

folks at home. In this regard, I am particularly blessed. I thank Doris, Erland,

Jennifer, Kathy, Maki and Roy for their friendship. I am grateful to Amina and

family (especially lovely lovely Nancy!!) for their kindness and giving me a

home in theirs. I thank Marley for going on walks with me often. I can’t make

him understand how much I love him. There will be no dog more dear to me.

I am grateful to teachers and supervisors who nurtured my interest in

science and engineering especially Drs. R. S. Chandel, P. Cheang, W. T.

Chen, Z. Chen, D. Z. Chi, C.-h. Chiu, S. K. Lahiri, H. Li, D. Mangelinck, H. M.

Phillips and O. Prabhakar. I thank Rinus Lee, Jofelyn Lye, Kim Shyong and

Wang Weide for being great friends.

I thank my family for their love. My happiest times in the last few years

were those moments when I was home, with bro and sis. I love them all,

forever and always. For Mansur Omar, my ever lasting love and prayers.

vi

VITA OF ZURUZI ABU SAMAH

June 2005

EDUCATION: Bachelor of Applied Science (Materials Engineering) (Honours) Nanyang Technological University, Singapore, 1997 PROFESSIONAL EMPLOYMENT: Student Engineer, Murata Electronics, Singapore Site, 1996 Research Officer, Microelectronic Materials, Processes and Packaging Program, Institute of Materials Research and Engineering, Singapore, 1997-2000 (Dr. Syamal K. Lahiri, Director; Dr. Dominique Mangelinck, Supervisor) Graduate Student Researcher, Materials Department University of California, Santa Barbara, 2000-2005 (Professor Noel C. MacDonald, Advisor) HONORS: International Fellowship, National Science Scholars Program, Science and Engineering Research Council, Agency for Science Technology and Research, Singapore. (2002-2004) Graduate Student Silver Award, Materials Research Society (MRS). (2004)

vii

SELECTED PUBLICATIONS:

1. A. S. Zuruzi, N. C. MacDonald, M. S. Ward, A. Kolmakov, M. Moskovits, C. R. Safinya, “Nanostructured Titania”, University of California Disclosure, UC Case Number 2005-531-1 (2005).

2. A. S. Zuruzi, M. S. Ward, N. C. MacDonald, “Fabrication and

characterization of patterned micrometer scale interpenetrating Au-TiO2 network nanocomposites”, Nanotechnology, 16, 1029 (2005).

3. A. S. Zuruzi, N. C. MacDonald, “Facile fabrication and integration of

patterned nanostructured titania for microsystems applications”, Adv. Func. Mater., 15, 396 (2005).

4. N. F. Bouxsein, L. S. Hirst, Y. Li, C. R. Safinya, Z. Abu Samah, N. C.

MacDonald, R. Pynn, “Alignment of filamentous proteins and associated molecules through confinement in microchannels”, Appl. Phys. Lett. 85, 5775 (2004).

5. D.Z. Chi, D. Mangelinck, A. S. Zuruzi, A. S. W. Wong, S. K. Lahiri,

“Nickel silicide as a contact material for submicron CMOS devices”, J. Electron. Mater., 30, 1483 (2001).

6. A. S. Zuruzi, C. H-. Chiu, S. K. Lahiri. K. N. Tu, “Roughness evolution

of Cu6Sn5 intermetallic during soldering”, J. Appl. Phys., 86, 4916 (1999).

7. A. S. Zuruzi, C. H. Chiu, W. T. Chen, S. K. Lahiri. K. N. Tu,

“Interdiffusion of high-Sn/high-Pb (SnPb) solders in low-temperature flip chip joints under reflow process”, Appl. Phys. Lett., 75, 3635 (1999).

8. A. S. Zuruzi, G. Dong, H. Li, “Diffusion bonding of aluminium alloy

6061 in air using an interface treatment technique”, Mat. Sci. and Eng., A259, 145 (1999).

viii

ABSTRACT


by

Zuruzi Abu Samah

This thesis describes research on a novel process to fabricate

integrated nanostructured titania (NST) features as functional components in

microsystems devices. NST features were formed by oxidizing Ti films in

aqueous hydrogen peroxide followed by thermal annealing. The oxidation

kinetics and properties of NST formed were investigated. The process

developed is compatible with current microelectronics manufacturing

practices for Si and plastic substrates.

Amorphous hydrated titania gels form when hydrogen peroxide (H2O2)

reacts with Ti. Oxidation of a blanket (unpatterned) Ti surface with hydrogen

peroxide results in a titania layer with high crack density. In this study, NST

was formed by reacting pre-patterned Ti thin films with H2O2 solution. Crack

elimination was achieved when exposed Ti films were below a threshold

dimension. Hydrated titania gel crystallizes into anatase after annealing at

300 °C for 8 hr. Crack elimination is thought to result from stress reduction in

titania gels due to patterning.

ix

Oxidation of Ti films occurs by nucleation and growth mechanism.

During growth, oxidation of Ti films with thickness 50 nm and below proceeds

at a constant rate until films are fully consumed. For Ti films with thickness

100 nm or thicker oxidation rate reduces significantly after a period of growth.

This reduction is attributed to a change in mechanism controlling growth of

the hydrated titania gel layer.

Functionality of NST formed and compatibility of the process with

current microelectronics manufacturing practices were demonstrated by

exploring three applications. First, a prototype conductometric gas sensor

was fabricated that used micrometer-scale NST pad arrays as sensing

elements. This sensor is capable of detecting hydrogen and oxygen gas at

concentration of a few parts per million (ppm). Second, micrometer scale Au-

NST interpenetrating network nanocomposite contacts in micro-switches were

fabricated by infiltrating NST features with Au using electroless deposition.

Third, results of cell-culture studies showed that mouse fibroblast cells

exhibited enhanced initial attachment on NST relative to silicon dioxide which

is commonly used in microsystems devices for biological applications.

x

TABLE OF CONTENTS

1. General introduction 1.1 Motivation and objectives 11.2 Nanostructured TiO2

1.2.1 Titania polymorphs 1.2.2 Synthesis routes 1.2.3 Physical properties

557

111.3 Concluding remarks 131.4 References 15

2. Formation and characterization of integrated nanostructured TiO2

2.1 Introduction 182.2 Formation of nanostructured TiO2 202.3 Formation of nanostructured TiO2 by aqueous oxidation of

pre-patterned Ti films 21

2.4 Results and discussion 2.4.1 Morphological study using optical and scanning

electron microscopy 2.4.2 Formation and elimination of cracks in NST pads 2.4.3 Surface chemistry study using X-ray

photoelectron spectroscopy 2.4.4 Phase evolution study using X-ray diffraction 2.4.5 Structural study using transmission electron

microscopy

2324

3033

3438

2.5 Conclusions 452.6 References 46

3. Kinetics of reaction between Titanium and aqueous hydrogen

peroxide 3.1 Introduction 503.2 Experimental procedure 543.3 Results and discussion

3.3.1 Characterization of Ti thin films and calibration of apparatus

3.3.2 Effect of film thickness 3.3.3 Effect of grain size 3.3.4 Effect of temperature 3.3.5 Effect of hydrogen peroxide concentration

58

66747882

3.4 Phenomenological model of Ti oxidation in aqueous hydrogen peroxide

85


xi

4. Gas sensing using nanostructured TiO2 4.1 Introduction 914.2 Gas sensing using nanostructured metal oxides 944.3 Integration of nanostructured TiO2 as sensing elements

on silicon 95

4.4 Fabrication of nanostructured TiO2 on Kapton® 1014.5 Results and discussion

4.5.1 Oxygen sensing 4.5.2 Hydrogen sensing

103104111


5. Fabrication of patterned micrometer scale interpenetrating Au–TiO2

network nanocomposites 5.1 Introduction 1205.2 Interpenetrating network composites 1225.3 Fabrication of interpenetrating Au–TiO2 network

nanocomposites 123

5.4 Results and discussion 1275.5 Integration of Au–TiO2 nanocomposites as contacts in

devices 140


6. Attachment of mouse fibroblasts on nanostructured TiO2

6.1 Introduction 1486.2 Attachment of cells on surfaces 1496.3 Results and discussion

6.3.1 Seeding of fibroblast on various materials 6.3.2 Morphology of fibroblast on patterned

nanostructured titania

152152157


7. Conclusions and future work

7.1 Conclusions 1627.2 Future work 163

xii

Chapter 1: General introduction

1.1 Motivation and objectives

Nanostructured materials are exciting due to their extraordinary physical

and chemical properties brought about by their small grain size (≤ 100

nm)1,2. Their high surface to volume ratio brings about unique properties that

are different from their bulk counterpart. Two examples of such unique

properties are enhanced sensitivity of electrical properties to chemical

species3-5 and enhanced biocompatibility of biological cells6-8. These two

properties alone make nanostructured materials attractive for integration into

future generation of microsystem (MEMS) devices as they render additional

functionality for chemical sensing and biocompatibility. With these enhanced

properties microsystem devices may find applications as electronic noses

and tongues for detection of chemicals9-11 as well as venture into the

biological milieu as implantable electronic (BioMEMS) devices such as drug

delivery12-14. Further motivation for integration of nanostructured materials

into microsystem devices is the enhancement of properties when

nanostructured materials are used in conjunction with other material

systems in the form of nanocomposites. Some of these nanocomposites

1

have lubricating and wear resistant properties15-17 and could be suitable as

contact materials in microsystem devices.

A large variety of nanostructured materials have been studied for

chemical sensing and enhancement of biocompatibility. In the former

application, nanostructured materials investigated have been of the discrete

type such as carbon nanotubes4, silicon nanowires5, palladium mesowire

arrays18, metal oxide nanowires19 and polymeric nanowires20. These

nanostructured materials have very high sensitivity to chemical species.

However, for implementation into practical devices, these nanostructures

need to be individually manipulated and positioned into place at specific

locations on a chip in a highly repeatable manner at the most cost effective

way. Presently, this is a major hurdle as there is no practical method to do

so reproducibly at low cost. Nevertheless, progress has been made towards

this goal.

In addition, any nanostructured material selected and the process used

for materials integration need to be compatible with Complementary Metal-

Oxide-Semiconductor (CMOS) processing. This is because, in most

applications, Si-based CMOS devices will be required to process information

generated. At present, most of the processes used to grow these

nanostructures require high temperatures. High processing temperatures

are undesirable because they result in high residual thermal stresses which

2

may lead to deformation of free standing structures ubiquitous in

microsystem devices. High processing temperatures also result in

interdiffusion and reactions at the interfaces of dissimilar materials which

may cause degradation of CMOS device characteristics.

For enhancement of biocompatibility, composites of nanostructured

materials such as carbon nanotubes and polymers are most commonly

used. However, unlike discrete nanostructures used in sensing applications,

nanostructured materials used to enhance biocompatibility are usually

applied over a larger area. Hence there is no real need for accurate

placement on surfaces of chips. In this aspect, implementation of

nanostructured materials for biocompatibility enhancement is, in relative

terms, more easily achieved.

Nanostructured titania has been widely investigated as a sensing

material due to its stability under adverse conditions and because its

specificity for various gases can be tailored by judicious use of surface

activation and dopants. In addition, it is widely accepted that titania

enhances biocompatibility of medical devices and implants. Therefore,

nanostructured titania is a suitable material for integration into microsystem

devices to render functionalities of gas sensing and biocompatibility.

The purpose of this research is to develop a technique of integrating

nanostructured titania into microsystem devices that is compatible with

3

material and tool sets used in conventional CMOS processing. Morphology

of nanostructured titania fabricated in the present research is porous and

sponge-like with walls consisting of nanocrystals. Nanostructured titania is

formed on rigid substrates such as Si, glass and Ti substrates as well as a

flexible organic (relatively) low-cost substrate; namely Kapton™ which is

commercially available. Emphasis will be placed on the evolution of titania

phases and morphology formed on these substrates. In addition, the

occurrence of cracks in nanostructured titania is closely studied as this

represents a potential source of reliability issue in microsystem device

applications. To demonstrate the functionality of nanostructured titania

fabricated and the compatibility of the process developed with Si-CMOS

processing, a prototype gas sensor is fabricated. In addition, the use of

integrated and patterned nanostructured titania as biocompatible cell

adhesion layers in microsystem is investigated. Also, a method of forming

integrated micrometer nanocomposites by infiltrating metal into the

nanostructured titania is explored. Having defined the scope and objectives,

the remainder of this chapter provides background information on

nanostructured titania, methods that are commonly used to synthesize

titania and selected properties of nanostructured titania that are relevant to

the present research.

4

1.2 Nanostructured TiO2

Titanium (IV) oxide (TiO2) commonly known as titania is widely used in a

number of industrial applications ranging from pigments in paints to coatings

on non-fogging surfaces. It has been recognized that properties of

nanostructured titania are different from the bulk form, which could lead to

new applications or provide better materials for existing ones. Currently,

there is tremendous effort to understand the energetics of various

nanostructured titania polymorphs as it controls size, morphology and

phases of titania formed.

1.2.1 Titania polymorphs

There are three known polymorphs of titania that exists in nature -

namely, rutile, anatase and brookite21. Rutile and anatase have tetragonal

crystal structure while brookite is orthorhombic. Rutile is the stable phase of

titania while anatase and brookite are metastable polymorphs. Hence under

ambient conditions anatase and brookite will transform to rutile when

kinetically permissible. Relative to bulk rutile, the enthalpies of formation of

bulk brookite and bulk anatase are higher by 0.71 ± 0.38 kJ/mol22 and 2.61 ±

0.41 kJ/mol23– indicating that in bulk form, rutile is the most stable phase

followed by brookite and anatase. It is noted that the anatase to rutile and

5

brookite to rutile polymorphic transformations do not occur reversibly, which

is in agreement with anatase and brookite being metastable phases.

The energetics and kinetics of nanostructured anatase, brookite and

rutile polymorphic transformations have recently been investigated by

Ranade et. al.22 as well as Zhang and Banfield24. It was shown that stability

of these TiO2 polymorphs is dependent on size. For spherical TiO2 particles

larger than about 204 nm, rutile is most stable. However between 204 nm to

38 nm TiO2 particles would exist as brookite. Below a critical size of about

38 nm, anatase is the most stable. The presence of these crossovers in

phase stability with variation in particle size is attributed to different surface

enthalpies of the various TiO2 polymorphs. The surface enthalpies of rutile,

brookite and anatase have been estimated to be 2.2 ± 0.2 J/m2, 1.0 ± 0.2

J/m2 and 0.4 ± 0.1 J/m2, respectively22. It is noted that phase stability is

governed by Gibbs free energy (ΔG=ΔH-TΔS) rather than the enthalpy (ΔH),

however data suggest that entropy of the various phases are similar.

Consequently, consideration of phase stability using enthalpies as

discussed above is valid. Because different methods of synthesizing titania

yield titania crystals with different size distribution, the phase evolution of the

various titania polymorphs in turn is dependent on the synthesis route.

6

1.2.2 Synthesis routes

Various methods for synthesizing nanostructured or nanoporous titania

have been reported in the literature. Here, only methods pertaining to

deposition of nanostructured titania films will be discussed as these are

relevant to the objective of the present work. Generally, these techniques

may be broadly classified as either chemical or physical methods. The main

chemical methods include anodization3,25, hydrolysis26,27, sol-gel15,28, and

spray pyrolysis29 while physical techniques include sputtering30,31,

supersonic cluster beam deposition32 and laser ablation33.

Anodization of Ti and its alloys to form porous nanostructured titania has

been used since the 1980’s to optimize adhesion of Ti alloy components in

joints. During anodization, the sample is immersed in an electrolyte solution

and oxidized. Recent effort includes the work of Varghese et. al.3 who

anodized Ti foils in dilute hydrofluoric acid to form nanotubes. One

parameter to control morphology of titania nanotubes was the voltage used.

It was found that with decreasing anodization voltage, the tube diameter,

wall thickness and length of the nanotubes decrease. At an anodization

voltage of 20 V, the average diameter of the tubes is 76 nm, wall thickness

is 27 nm and length is 400 nm while at 10 V the corresponding values are

22 nm, 13 nm and 200 nm, respectively. In general, porosity is increased by

increasing the voltage, electrolyte concentration and temperature. This

7

method has also been used to form nanostructured titania from Ti-6Al-4V

alloy. As formed titania nanotubes are amorphous but transformed to

anatase at about 300°C. Generally a dense amorphous TiO2 barrier layer a

few nanometers thick is formed between the porous nanostructured titania

and the unreacted Ti foil.

Another chemical route is hydrolysis. In this method, insoluble metal

hydroxides precipitate from aqueous solution, which are then subsequently

converted into its metal oxide by heat-assisted dehydration26,27. To form

nanostructured titania using this technique, titanium-containing precursors

such as an ethanol solution of tetrabutyl titanate (Ti(C4H9)4) are added to

deionized water resulting in the precipitation of titanium hydroxide gel which

is then annealed to form titania. Using this technique, coatings of

nanostructured titania with particle size ranging from a few nanometers to

hundreds of nanometers has been fabricated.

Sol-gel processing is a technique in which a sol is first formed followed

by formation of a gel15,28. Traditionally sol-gel synthesis uses either a

colloidal suspension or inorganic precursors as the starting material. The

latter approach is more often reported in the literature and is based on the

use of metal alkoxides which have the general formula M(OR)x. The

alkoxide can also be formed by reaction of a metal species such as a metal

hydroxide or metal halide with an alcohol. The metal alkoxide then

8

undergoes hydrolysis to form a colloidal sol, which further undergoes a

condensation reaction to form a gel. In essence the sol-gel technique

involves two types of reactions namely hydrolysis and condensation, as

described below:

Hydrolysis: -MOR + H2O → -MOH + ROH

Condensation: -MOR + ROM → -MOM- + ROH

OR-MOH + HOM → -MOM- + H2O

The usefulness of the sol-gel technique rests on the fact that parameters

that affect either of these reactions will impact the properties of the gel.

Hence by controlling these parameters, tremendous control of the properties

of the gel obtained is possible.

Another chemical technique to form nanostructured titania is spray

pyrolysis29. A major advantage of spray pyrolysis is that it allows the

synthesis of titania powders with a wide range of diameter, from a few to a

hundreds of nanometers, at relatively low cost. In spray pyrolysis, the vapor

of precursors, such as titanium tetrachloride (TiCl4), react with oxygen at

high temperature to form titania powder, usually in the form of aggregates.

The reaction may be adequately represented as:

TiCl4 + O2 → TiO2 + 2Cl2

9

The size of the titania particles is dictated by material and process

conditions. Since combustion of TiCl4 is exothermic, little fuel is needed to

sustain the flame and propagate the reaction, except that required to initiate

the reactor.

Among the physical techniques, cluster beam deposition is one of the

newest techniques developed. In this technique, clusters of titanium are

produced by a pulsed microplasma cluster source under high vacuum

conditions32. Titania is formed once these clusters are exposed to air, since

Ti is highly reactive to oxygen. Also, clusters of different sizes are separated

in the radial direction of the beam and nanostructured titania with various

cluster sizes can be obtained by intersecting substrates at different locations

of the beam. Clusters with largest mass are found nearest to the axis of the

beam while the smallest mass is farthest away from the beam axis. Using

this technique, nanostructured titania coatings with average particle size of

about 10 nm have been deposited.

Sputtering has been used in industrial settings to deposit titania for low

emissivity coatings and high reflection coatings30,31. Titania coatings were

formed by sputtering either Ti metal targets in an Ar plasma in the presence

of oxygen or titania targets in a pure Ar plasma. It was found that titania

formed by reactive sputtering adheres better to an organic (polyethylene

terepthalate) substrate due to formation of Ti-O-C bonds at the interface

10

between the titania layer and the substrate. Using a closed loop reactive

sputtering technique, coatings of nanostructured titania having average

grain sizes of about 10 nm have been deposited.

In laser ablation deposition of nanostructured titania, a laser source

ablates a Ti target in a chamber purged with a flow of oxygen33. The ablated

Ti species then reacts with oxygen to form titania nanoparticles that are

collected downstream. Average diameter of titania nanoparticles formed

increases with oxygen flow rate but is independent of the laser fluence.

Titania nanoparticles formed using this technique, however, tend to

coalesce on specific planes to form single crystals with dislocations at the

interface. Particles formed are about one order of magnitude larger than that

of the individual coalescing particles.

1.2.3 Physical properties

It is now widely accepted that the properties of nanostructure materials

are usually different than that of its bulk counterpart. In this section, only

properties with direct relevance to the present research will be briefly

discussed - namely electrical, chemical adsorption and wear resistant

properties.

11

As in other material systems, the electron transport properties of

nanostructured titania are different from its bulk counterpart. Because of

high surface to volume ratio in nanostructured titania, electron transport was

found to be significantly enhanced by surface effects. Rothschild et. al.34

recently showed that surface and grain boundary barriers, brought about by

electron trapping at interface states associated with chemisorbed oxygen

species, drastically reduced the conductance of nanocrystalline titania thin

films. By annealing in a reducing ambient however, the conductance can be

recovered.

The adsorption of chemical species such as oxygen, which may affect

the electrical properties as discussed above, in turn is affected by the size of

the titania nanocrystals. Recent work by Zhang et. al.35 showed that the

adsorption constant of nanostructured anatase for various organic acids

may show up to 70 fold increase when size of titania particles was

decreased from 16 to 6 nm. This observation is interesting in light of recent

work that showed enhanced cellular attachment on nanocrystal ceramics

compared to their microcrystal forms6. In physiological conditions, cellular

attachment on surfaces is always preceded by adsorption of proteins such

as fibronectin. Only after these proteins have been adsorbed will cellular

attachment and spreading occur36.

12

Another interesting property is the enhanced wear resistance of

composites containing nanostructured titania. Recent work has shown that

Au-titania nanocomposites produced using the sol-gel route exhibit excellent

wear resistance15. It was found that beyond about 5 mol% Au, the wear

resistance of the composite was significantly increased. In these

nanocomposites, Au was present as particles embedded in the titania

matrix. Wear resistance of the Au-titania composite is greater than that of Au

or its alloys which, at present, are widely used as contacts in many MEMS

devices.

1.3 Concluding remarks

The implementation of novel nanostructured materials promises

additional functionalities in microelectronic devices. Integration of

nanostructured TiO2 in microsystems is particularly important as titania is a

versatile material with applications in energy conversion, chemical sensing,

biocompatibility and drug delivery. However, there are a number of critical

issues that remain to be resolved before nanostructured TiO2 can be

implemented in devices for large volume manufacturing. First, the issue of

crack formation which is a common occurrence in nanostructured titania

structures produced using the commonly used sol-gel technique needs to be

13

resolved. Second, for successful implementation nanostructured titania must

be integrated using low-cost techniques that are compatible with current

microelectronics device manufacturing practices. In this context,

compatibility refers to both materials as well as process compatibility.

Techniques to integrate nanostructured titania that use existing process

tools and material sets would, of course, be ideal since no additional cost

would be incurred. It is the objective of the present research to realize such

a technique.

The following chapters describe a process developed that addresses the

above issues. In addition, investigations of integrated nanostructured titania

for gas sensing and biological cell attachment are presented. Also, using the

porous nanostructured titania as the ceramic component, patterned

micrometer scale interpenetrating metal-nanostructured titania

nanocomposites have been fabricated. Because the metal phase is

percolating throughout the structure, the composite is expected to have low

electrical resistance and possessing high wear resistance and hardness of

the titania phase.

14

1.4 References

1. R. Breslow, M. V. Tirrell, J. K. Barton, M. Barteau, C. R. Bertozzi, R. A. Brown, A. P. Gast, I. E. Grossmann, J. M. Meyer, R. W. Murray, P. J. Reider, W. R. Roush, M. L. Shuler, J. J. Siirola, G. M. Whitesides, P. G. Wolynes, — Beyond the Molecular Frontier: Challenges for Chemistry and Chemical Engineering, National Academies Press, Washington D. C, USA (2003)

2. A. Dowling, R, Clift, N. Grobert, D. Hutton, R. Oliver, O. O’neill, J.

Pethica, N. Pidgeon, J. Porritt, J. Ryan, A. Seaton, S. Tendler, M. Welland, R. Whatmore — Nanoscience and nanotechnologies: opportunities and uncertainties, The Royal Society and The Royal Academy of Engineering, London, UK (2004)

3. O. K. Varghese, D. Gong, M. Paulose, K. G. Ong, E. C. Dickey and C. A.

Grimes, Advanced Materials, 15, 624 (2003). 4. J. Kong, N. Franklin, C. Zhou, M. G. Chapline, S. Peng, K. Cho, H. Dai,

Science, 287, 622 (2000). 5. C. Y. Cui, Q. Wei, Q., H. Park, C. M. Lieber, Science, 293, 1289 (2001). 6. T. J. Webster, C. Ergun, R. H. Doremus, R. W. Siegel and R. Bizios,

Biomaterials, 22, 1327 (2001). 7. T. J. Webster, L. S. Schadler, R. W. Siegel and R. Bizios, Tissue

Engineering, 7, 291 (2001). 8. H-. H. Huang, S-. J. Pan, Y-. L. Lai, T-. H. Lee, C-. C. Chen and F-. H.

Lu, Scripta Materialia, 51, 1017 (2004). 9. M. Pardo and G. Sberveglieri, MRS Bulletin, 29, 703 (2004). 10. J. P Novak, E. S. Snow, E. J. Houser, D. Park, J. L. Stepnowski, R. A.

McGill, Applied Physics Letters, 83, 4026 (2003). 11. Q. Wan, Q. H. Li, Y. J. Chen, T. H. Wang, X. L. He, J. P. Li, C. L. Lin,

Applied Physics Letters, 84, 3654 (2004). 12. C. R. Martin, Science, 266, 1961(1994).

15

13. T. Paunesku, T. Rajh, G. Wiederrecht, J. Maser, S. Vogt, N. Stojićević, M. Protić, B. Lai, J. Oryhon, M. Thurnauer, G. Woloschak, Nature Materials, 2, 343 (2003).

14. L. Leoni and T. Desai, Advanced Drug Delivery Reviews, 56, 211 (2004). 15. W- M. Liu, Y-. X. Chen, G-. T. Kou, T. Xu and D. C. Sun, Wear, 254, 994

(2003). 16. M. C. Simmonds, A. Savan, E. Pfluger and H. V. Swygenhoven, Journal

of Vacuum Science and Technology, 19, 609 (2001). 17. C. Donnet and A. Erdemir, Tribology Letters, 17, 389 (2004). 18. F. Favier, E. C. Walter, M. P. Zach, T. Benter, R. M. Penner, Science,

293, 2227 (2001). 19. A. Kolmakov, Y. Zhang, G. Cheng, M. Moskovits, Advanced Materials,

15, 997 (2003). 20. Y. Wang, X. Jiang, Y. Xia, Journal of the American Chemical Society,

125, 16176 (2003). 21. A. Navrotsky and O. J. Kleppa, Journal of the American Ceramic

Society, 50, 626 (1967). 22. M. R. Ranade, A. Navrotsky. H. Z. Zhang, J. F. Banfield, S. H. Elder, A.

Zaban, P. H. Borse, S. K. Kulkarni, G. S. Doran and H. J. Whitfield, Proceedings of the National Academy of Sciences, 99, 6476 (2002).

23. T. Mitsuhashi and O. J. Kleppa, Journal of the American Ceramic

Society, 62, 356 (1979). 24. H. Z. Zhang and J. F. Banfield, Journal of Physical Chemistry B, 104,

3481 (2000). 25. J. P. Wightman and J. A. Skiles, SAMPE Journal, 24, 21 (1988). 26. H. Hirashima, H. Imai, M. Y. Miah, I. M. Bountseva, I. N. Beckman and V.

Balek, Journal of Non-crystalline Solids, 350, 266 (2004). 27. W. F. Zhang, M. S. Zhang, Z. Yin, Physica Status Solidi A-Applied

Research, 179, 319 (2000).

16

28. C. J. Brinker and G. W. Scherer, Sol-Gel Science: The Physics and

Chemistry of Sol-Gel Processing, Academic Press, San Diego (1990). 29. S. E. Pratsinis, Progress in Energy Combustion Science, 24, 197 (1998). 30. R. Dannenberg and P. Greene, Thin Solid Films, 360, 122 (2000). 31. H. Tang, K. Prasad, R. Sanjinés and F. Lévy, Sensor and Actuators B,

26-27, 71 (1995). 32. E. Barborini, I. N. Kholmanov, A. M. Conti, P. Piseri, S. Vinati, P. Milani,

C. Ducati, The European Physical Journal D, 24, 277 (2003). 33. H. Huang and X. Yao, Surface Coatings and Technology, 191, 54

(2005). 34. A. Rothschild, Y. Komem, A. Levakov, N. Ashkenasy and Y. Shapira,

Applied Physics Letters, 82, 574 (2003). 35. H. Z. Zhang, R. L. Penn, R. J. Hamers, J. Banfield, Journal of Physical

Chemistry B, 103, 4656 (1999). 36. B. D. Ratner, A. S. Hoffman, F. J. Schoen and J. E. Lemons,

Biomaterials Science: An Introduction to Materials in Medicine, Academic Press, San Diego (1996).

17

Chapter 2: Synthesis and characterization of

integrated and patterned nanostructured TiO2

2.1 Introduction

The excellent properties of nanostructured titania (NST) makes it the

material of choice in many applications. Porous NST has been used for

enhancing performance of implants1-3, gene delivery4, energy conversion5,

separation6, catalysis7 and gas sensing8-10. Many techniques have been

proposed for fabricating integrated and patterned micrometer scale NST

features. These techniques may be classified as either a reductive approach

in which features are etched from a continuous layer or the additive

approach where patterned features are directly deposited on the substrate.

Reductive approaches to forming patterned NST films have been

practiced for some time. A variety of techniques have been used to deposit

a continuous titania film11-14. The sol-gel technique, in particular, has

received tremendous attention as it renders molecular-level control, is

relatively low cost and allows incorporation of various metal dopants into

titania matrices15,16. These films are then patterned using techniques such

as reactive ion etching17 and embossing18,19 and laser trimming20. It is

noted that except for reactive ion etching, other techniques are not

18

compatible with high-volume semiconductor manufacturing processes.

Patterning of the TiO2 film could be done either before or after an annealing

step in which the film is heated at elevated temperature to convert the

amorphous titania gel into crystalline TiO2. In addition, precautions are

usually required to ensure crack formation21-23, does not occur and carbon,

from the organic precursors, is not incorporated24 in or on the NST features.

Additive approaches have been widely reported in the literature.

Because of their relatively low cost, screen printing23,25,26 and self-

assembled monolayers-assisted deposition of TiO2 have received particular

attention27-29. As the name implies, screen printing involves printing paste

containing titanium dioxide powders at desired locations on substrates using

a hard mask (the screen). Although this technique is low-cost, there are

several limitations to the process. First, the use of a hard mask means that

accurate alignment techniques are required to deposit the paste at desired

locations on the substrate. Second, the probability of adjacent

paste/features developing bridges increases as distance between adjacent

features decreases thus lowering process yield. Third, the dimension of

smallest features that can be deposited using a hard mask is about 100 μm,

which is larger than dimensions that can be deposited using

photolithography. For some microsystems applications these limitations

need to be resolved before screen printing can be implemented.

19

The microprinting of self-assembled monolayers (SAMs) using stamps

was first proposed by Whitesides et. al. 30. Selective deposition of TiO2

features occurs by interactions of functional group of the self-assembled

monolayers with TiO2 nuclei homogenously nucleated in solution. Because

surface coverage of SAMs is, in most cases, not high, this method results in

poor yield especially when large-area substrates are used. In addition,

relatively long time - up to a few hours - is required for deposition of TiO2

using this method28. Furthermore, the edge acuity of titania features

deposited using SAMs-assisted direct deposition is generally poor28.

2.2 Formation of nanostructured TiO2

Aqueous oxidation of Ti surfaces is an attractive technique for growing

NST. Using aqueous hydrogen peroxide (aq. H2O2) solution as an oxidant

Wu et. al. 31 reported the formation submicrometer porous titania layers from

thick sheets of Ti while Tengvall32-34 formed transparent bioactive titania gel

from Ti powder and unpatterned films. Similarly, Nishiguchi et. al.35 reported

the formation of a porous titania layer by reacting Ti with aqueous sodium

hydroxide. However, titania layers formed have high crack density and

delaminated extensively from the Ti substrate.

20

In this chapter, we present a technique that eliminates crack formation

and delamination in titania layers by oxidizing Ti thin films that have been

patterned below a threshold dimension. The issue of carbon incorporation

does not arise since no organic precursor is used. In addition, the patterning

technique developed allows the fabrication of miniaturized NST features and

formation of crystalline titania at relatively low temperatures, thus permitting

use of aluminium-based metallization. Hence, the technique described in

this chapter represents a practical route for fabricating and integrating NST

structures into Si-based microsystem devices.

2.3 Formation of nanostructured titania by aqueous

oxidation of pre-patterned Ti films

This research investigated the formation of NST on Ti bulk sheets and

thin films. Bulk Ti sheets (Goodfellow, 99.6% purity and 500 μm thick) were

first polished to a mirror finish with 0.3 μm colloidal silica and then rinsed

copiously with de-ionized (DI) water (>18.9 MΩ) with ultrasonic agitation. To

prepare Ti thin film samples, we used 2.5 cm square pieces of N-type

Si(100) (Mitsubishi Electronic Materials) as substrates. Si chips were

thermally oxidized at 1100 °C to grow 1 μm thick SiO2 layer (T-SiO2). These

chips were then cleaned with ultrasonic agitation for 5 min each in acetone,

21

2-propanol and de-ionized (DI) water (18.9 MΩ) and blown dry with nitrogen

prior to deposition of Ti films. A schematic flow diagram for formation of NST

on patterned Ti pad arrays is shown in Figure 2.1. Ti thin films were

patterned using either lift-off or selective masking process. In selective

masking, Ti film was electron beam evaporated on Si chips followed by SiO2

deposition. Silicon dioxide (PECVD-SiO2) was deposited using plasma-

enhanced chemical vapor deposition (PECVD) using silane (SiH4) and

nitrous oxide (N2O) precursors at 250 °C. Photo resist (PR) was deposited

on the SiO2 layer and patterned. The pattern on the PR layer was

transferred to the SiO2 layer by etching with CHF3 gas. After patterning, PR

was removed by soaking in acetone. In the lift-off technique, a PR layer was

deposited on Si chips and patterned. A Ti film was then evaporated. These

Si chips were then soaked in acetone for 24 hr, rinsed sequentially in 2-

propanol, DI water and blown dry. For depositing blanket Ti films, Si chips

were used as cleaned. The process pressure during evaporation of Ti films

was ~5.0 x 10-7 Torr. All Ti sources were of 99.995% purity or better and

cleaned by pre-evaporation to remove native oxide layer on surface.

22

Figure 2.1. Schematic of procedures for forming NST pad arrays. (a) Lift-off technique; (b) Selective masking technique.

2.4 Results and discussion

NST was formed by aging samples in aqueous H2O2 solution. Prior to

aging, Ti films were acid pickled in dilute hydrochloric acid for ~2 min to

remove native oxide layer, rinsed in DI water and then blown dry with

nitrogen. Aging was done in an oven at 80 ± 2 °C in air. Samples were

stored in a vacuum box prior to analysis. Crystal structure was analyzed by

23

X-ray diffraction (XRD) in Bragg-Brentano configuration using CuKα radiation

(1.5406Ǻ) (Phillips X’pert-MPD). Structural characterization was done using

an FEI dual beam focus ion beam (FIB) system equipped with Ga ion and

electron columns for high resolution machining and imaging, respectively.

Micromachining was done using a Ga ion current of 100 pA. Surface

chemical species were determined using a Kratos Axis Ultra X-ray

photoelectron spectroscopy (XPS) system. High resolution XPS scans were

obtained with monochromated Al Kα source (1486.6 eV) and 20 eV pass

energy with steps of 0.05 eV at a base pressure of 7.5 x 10-9 Torr. XPS

spectra collected were fitted to line shapes constructed from a linear

combination of Gaussian and Lorentzian profiles using a commercial

software (CasaXPS). Transmission electron microscopy (TEM) was done

using an FEI Sphera T20 machine operating at 200 kV. The following

sections present and discuss results of these investigations.

2.4.1 Morphological study using optical and scanning electron

microscopy

SEM micrographs of unpatterned bulk Ti sheets and Ti films after aging

in aq. H2O2 revealed formation of NST layers with high crack density -

Figures 2.2 (a) to (d). High resolution SEM shows that titania layers consist

of walls of pores having thicknesses and pore diameters ranging from 25 nm

– 50 nm and 50 nm – 200 nm, respectively. The high crack density results in

24

the formation of ‘grains’ about 5 μm – 7.5 μm average diameter. Cracks on

NST layers formed on thin films extend from the surface to the thermally

grown SiO2 layer and resulted in complete delamination NST layers

especially after prolonged oxidation times. The morphology of titania layers

formed on bulk Ti sheets and evaporated Ti films is similar. However,

delamination of NST layers formed on bulk Ti sheets is less extensive and

cracks are narrower. In addition, pores in the NST layer formed on bulk Ti

sheets are smaller.

Figure 2.2. Crack formation in NST layers formed on unpatterned Ti surfaces. SEM micrographs of unpatterned (a and b) bulk Ti sheets; and (c

and d) Ti thin films showing high crack density in NST formed on unpatterned Ti surfaces.

25

By using Ti films patterned below a threshold dimension, crack formation

on NST layers was eliminated. Figures 2.3 (a) to (e) are SEM micrographs

of NST layers formed from patterned Ti square pads of various dimensions

after aging for 2.5 hr at 80 °C in 10% aq. H2O2 solution - thickness of the Ti

layer is 2.0 μm. Cracking is most extensive on 100 μm pads and resulted in

the NST/unreacted Ti bilayer peeling off from the Si substrate – inset in

Figure 2.3 (a). Cracking is significantly reduced for 70 μm pads and for

arrays of 20 and 5 μm pads, crack formation is eliminated. In addition, gaps

developed between NST/unreacted Ti bilayer and the mask oxide. However,

NST pad arrays formed using lift-off technique have little adhesion to the

SiO2 substrate and delaminate easily during aging in aq. H2O2 solution.

Figure 2.3 (e) shows a 20 μm NST pad displaced from its original position.

In contrast, NST pads formed using selective masking have excellent

adhesion to the underlying SiO2 layer. Figure 2.3 (e) also shows that NST is

formed on sidewalls of pads formed using lift-off technique. In contrast, NST

is not observed on the sidewalls of pads formed using selective masking –

inset in Figure 2.3 (c) and (f). This observation suggests that gaps were

formed during the latter stages or possibly after aging.

Cracks were not formed on 20 μm and 5 μm pads even after annealing

at 300 °C for 8 hr, Figure 2.3 (f). However, the width of gaps between

NST/unreacted Ti bilayer and the mask oxide increased due to pad

26

shrinkage. For 20 μm pads, gap width is about 0.7 μm after drying in air but

increased to about 1.2 μm after annealing at 300 °C for 8 hr - compare

insets in Figure 2.3 (c) and (f). For arrays of 5 μm pads, width of gap before

and after annealing is estimated to be about 50 nm and 90 nm, respectively.

These observations suggest that compressive forces are created in NST

layers during aging as well as during oxidation. This hypothesis is further

supported by the curvature (concave upwards) of the peeled-off

titania/unreacted Ti bilayer - inset in Figure 2.3 (a).

Porosity of NST layers obtained could be due to morphology of the

intermediate gel layer formed during aging in aq. H2O2. Reaction of metallic

Ti with hydrogen peroxide had been investigated by Tengvall32-34 and was

shown to result in formation of a hydrated TiO2 gel layer. Recent studies by

Wu et. al. indicate that a submicron porous titania layer results when this gel

layer was annealed31. However, no high resolution microscopy images were

provided for comparison. The dark brown layer that appears after aging in

aq. H2O2 is a hydrated TiO2 gel layer. Gel layers observed by Tengvall were

yellowish. However, the color difference could be attributed to a difference in

concentration of aq. H2O2 solutions used. Focused ion beam milling was

used to investigate the structural properties of NST layers. Figures 2.4 (a

and b) are cross-section SEM micrographs obtained after milling NST layers

grown on 2.0 μm thick evaporated Ti pads after aging in 10% aq. H2O2 for

27

2.5 hrs at 80 °C. The NST/unreacted Ti interface is robust with no

delamination. In addition, the NST layers have uniform thickness with a

planar NST/unreacted Ti interface. Figures 2.4 (c) and (d) are cross-section

SEM micrographs of supported TiO2 membranes formed by oxidizing 0.35

μm thick evaporated Ti films by aging in 10 % aq. H2O2 for 3.5 hrs at 80 °C.

Although the Ti films were completely oxidized no cracks were observed on

the 20 μm pad arrays. In addition, gaps were not observed between NST

membranes and the mask oxide. This indicates that the thickness of Ti films

affects the extent of shrinkage of NST patterns.

28

Figu

re 2

.3. C

rack

elim

inat

ion

in N

ST

laye

rs fo

rmed

on

patte

rned

Ti f

ilms

belo

w a

thre

shol

d di

men

sion

. SE

M m

icro

grap

hs o

f NS

T la

yers

form

ed o

n (a

) 100

; (b)

70;

(c a

nd f)

20

and

(d) 5

μm

pad

s us

ing

sele

ctiv

e m

aski

ng; a

nd (e

) 20 μm

pad

s us

ing

lift-o

ff. C

rack

form

atio

n is

el

imin

ated

for 2

0 an

d 5 μm

pad

s. C

ompa

rison

of 2

0 μm

pad

arr

ay, (

c) b

efor

e an

d (f)

afte

r an

neal

ing

at 3

00 ºC

for 8

hr s

how

s an

incr

ease

of g

ap w

idth

afte

r ann

ealin

g (s

ee in

sets

).

29

Figure 2.4. Structure of NST layer interfaces. Cross-section SEM micrographs obtained after milling (a and b) NST layers formed after partial

oxidation of a 20 μm square pad Ti film, 2.0 μm thick; (c and d) NST membranes formed after complete oxidation of a 0.35 μm thick Ti film.

2.4.2 Formation and elimination of cracks in NST pads

Cracks in NST layers may start forming during oxidation, because of

stresses due to thermal mismatch between NST and residual titanium

layers, and/or during drying, as a result of stresses associated with

shrinkage. To elucidate the cause of crack formation, NST arrays were

observed while still in aq. H2O2 solution at room temperature, and again after

drying. Figure 2.5 (a) to (c) and (d) to (f) are optical images of 20 μm, 40 μm

30

and 50 μm pad arrays of NST before and after drying, respectively. In these

experiments, 0.35 μm thick Ti films were completely oxidized to form

supported TiO2 membranes. Prior to drying, cracks were observed only on

50 μm NST pads – it was found that 30% of pads were cracked (100 pads of

each size were studied). After drying, cracks were observed on all pads of

the 40 μm and 50 μm arrays - however, the 20 μm NST pad array remains

crack-free. These observations indicate that crack formation in NST is

primarily due to stresses associated with shrinkage during drying.

Crack elimination below a threshold dimension could be explained as a

result of load transfer between NST pads and the substrate at the edges.

Assuming NST pads are elastic membranes, load transfer at the edges is

described by the analysis of Freund and Suresh36. When NST pads are

dried, tensile stress is generated. However, points at the pad edges are

traction free as stress is relieved. However, for points at a distance away

from the edges stress is unrelaxed. The stress state approaches the

equibiaxial stress asymptotically with distance away from the edge. Hence

points away from the pad edge experience larger tensile force during drying.

As sizes of NST pads get larger a smaller volume of NST pads are nearer to

the edge. Consequently, a larger volume of the NST is experiencing a

higher tensile force which ultimately leads to crack formation when a

threshold NST pad size is exceeded.

31

Figu

re 2

.5. C

rack

form

atio

n in

NS

T pa

d ar

rays

dur

ing

oxid

atio

n an

d dr

ying

ste

ps. O

ptic

al

mic

rogr

aphs

of N

ST

arra

ys (a

-c) i

n aq

. H2O

2 afte

r oxi

datio

n; a

nd (d

-f) a

fter d

ryin

g. T

hese

imag

es

clea

rly s

how

that

cra

ck fo

rmat

ion

occu

rs p

rimar

ily d

urin

g th

e dr

ying

ste

p. T

he p

ads

in th

e ar

rays

ar

e 20

, 40

and

50 μ

m s

quar

es a

s in

dica

ted.

32

2.4.3 Surface chemical study using X-ray photoelectron spectroscopy

X-ray photoelectron spectroscopy of NST layers suggest that aging in

aq. H2O2 solution resulted in the formation of TiO2 species only. Figures 2.6

(a) and (b) are survey and high resolution XPS spectra, respectively, of NST

layer formed on evaporated Ti thin film. Similar results were obtained for

NST layers formed on bulk Ti sheets. All spectra are referenced to C1s peak

at 285.0 eV37. Assuming a Tougaard background, raw spectra were fitted

using Gaussian-Lorentzian components with appropriate constraints for

area, full-width-at-half-maximum and position parameters using a

commercial software (CasaXPS). From the analysis, binding energies for

the Ti 2p3/2 components were found to be 459.0 eV and 458.9 eV for titania

on bulk and evaporated Ti film, respectively. For NST on bulk and

evaporated thin film Ti, the Ti 2p1/2 components have a value of 464.8 eV.

These experimental values obtained are close to those reported in the

literature for TiO2 of 458.9 eV and 464.6 eV for Ti 2p3/2 and Ti 2p1/2

components, respectively38,39.

33

Figure 2.6. XPS spectra of TiO2 species after drying. (a) Survey and (b) high resolution XPS spectra of NST layer formed on evaporated thin Ti film.

2.4.4 Phase evolution study using X-ray diffraction

XRD studies show that amorphous TiO2 and nanocrystals of anatase

TiO2 polymorph were formed after aging and the amorphous phase

transforms to anatase upon annealing. Figure 2.7 shows XRD spectra of

evaporated Ti film, as-aged and annealed TiO2 layer formed from

evaporated blanket Ti films. Spectra of the sample aged in 10 % aq. H2O2

solution for 2.5 hr at 80 °C exhibit Ti peaks, which correspond to unreacted

Ti in aged films and three broad peaks at 2θ values of 25.20°, 47.97° and

62.68° which can be assigned to anatase 101, 200 and 204 planes. The

broadness and low intensity of these peaks suggest that as-formed titania

34

layer consists of anatase nanocrystals in a largely amorphous titania matrix.

Upon annealing at 300 °C for 8 hr, these peaks sharpened significantly and

increased in intensity. The peak sharpening and intensity increase upon

annealing are due to transformation of the amorphous phase to anatase

which also agrees with the appearance of additional peaks at 2θ values of

53.91°, 54.95° and 75.04°. These latter peaks correspond to 105, 211

and 215 reflections of anatase. All anatase peaks in the spectrum of

annealed TiO2 layer match perfectly to corresponding ones in spectrum

collected from reference TiO2 anatase powder (Alfa Aesar, 99.6 %). No

peaks from other TiO2 polymorphs are observed in spectra of annealed

samples. Hence, XRD and XPS data indicate that only nanocrystals of

anatase TiO2 and an amorphous titania phase are formed after aging.

Subsequent annealing transforms the amorphous phase to anatase.

Formation of single phase nanostructured anatase from amorphous TiO2 by

annealing in air as well as the coexistence of these phases had been

reported previously40,41.

35

Figure 2.7. XRD spectra of TiO2 species formed after aqueous oxidation. XRD spectra of (i) anatase reference sample, (ii) evaporated Ti film, (iii) as-

aged and (iv) annealed NST layer. ( g indicates anatase peak)

X-ray pole figure determination of bulk and thin film Ti surfaces indicates

orientation of anatase crystals formed during annealing is not influenced by

texture of the underlying Ti substrate. Samples in this experiment are

partially oxidized to prevent delamination during oxidation. Figures 2.8 (a)

and (b) are pole figures collected for Ti 0002 (2θ = 38.42 °) of bulk Ti sheet

and evaporated thin Ti film, respectively, before aging. These pole figures

agree with other studies and are typical of hot rolled Ti sheets and deposited

thin Ti film, respectively42,43. Figures 2.8 (c) and (d) are corresponding

scans for anatase 101 (2θ = 25.20°), collected after aging in aq. H2O2

solution and subsequent annealing at 300 °C for 8 hr. The similarity of

Figures 2.8 (c) and (d) and the narrow intensity range of these spectra

36

suggest that anatase crystals formed during annealing have a random

orientation for both bulk and thin film Ti and texture of the underlying Ti

substrate does not influence orientation of these anatase crystals. This is in

agreement with prior reports which showed that crystallization of anatase

during annealing of amorphous TiO2 layer takes place by recrystallization of

small anatase particles and solid-state aggregation of amorphous

particles40. These mechanisms would be expected to produce randomly

oriented anatase crystals.

Figure 2.8. XRD pole figure of parent Ti and TiO2 formed after aqueous oxidation. Pole figures shown are for Ti 0002 of (a) Ti sheet and (b)

evaporated thin Ti film, both before aging; and of anatase 101 collected after aqueous oxidation and annealing of (c) bulk and (d) thin Ti film

samples.

37

2.4.5 Structural study using Transmission Electron Microscopy

In this section we discuss in detail structural characterization of NST

using TEM. Cross-sectional TEM samples were prepared using a ‘lift-out

technique’, which involves micro-machining a thin slice of the NST layer

using a focus ion beam (FIB). The use of this technique for TEM preparation

of fragile and porous samples is relatively new44,45 and is described in the

following. Figure 2.9 shows important steps in the preparation of cross-

sectional TEM samples. First a membrane was made by milling a trench on

either side it. The membrane was thinned to 100 nm and then suspended by

milling its bottom and sides. A notch is then made at each end of the

supporting beam. Subsequently, the membrane was lifted out using a glass

needle attached to a micromanipulator with the help of microscopes and

placed on a carbon coated Cu grid. Three TEM samples were made from

different NST patterns to ensure that TEM observations were representative

of the whole specimen.

38

Figu

re 2

.9. S

teps

in c

ross

-sec

tion

sam

ple

prep

arat

ion

for T

EM

usi

ng F

IB m

illing

. (a)

Sam

ple

befo

re

mill

ing;

(b) t

renc

h m

illin

g on

eith

er s

ide

of m

embr

ane;

(c) m

embr

ane

afte

r fur

ther

thin

ning

; (d)

mem

bran

e su

spen

ded

by m

illing

the

side

s an

d bo

ttom

; (e)

one

end

of t

he m

embr

ane

befo

re n

otch

ing

and

(f) th

e ot

her e

nd a

fter n

otch

ing.

39

Figu

re 2

.10.

Typ

ical

cro

ss-s

ectio

nal T

EM

imag

es a

nd d

iffra

ctio

n pa

ttern

s. (a

) Low

mag

nific

atio

n im

age;

(b) a

nd (c

) hig

h m

agni

ficat

ion

imag

es o

f tw

o sa

mpl

es ta

ken

from

diff

eren

t NS

T pa

ttern

s; (d

), (e

) an

d(f)

are

SA

ED

patte

rns

corr

espo

ndin

gto

regi

ons

deno

ted

by(1

),(2

)and

(3)r

espe

ctiv

ely

in(c

).

40

Figures 2.10 (a) to (c) are TEM images of two samples from different

NST pads showing typical cross-sections. The difference in contrast

indicates that exposed Ti film is completely oxidized. Pores traversing the

entire thickness of NST layer are observed. The diameter of these pores is

about 150 nm. Also, the NST layer is considerably thinner than that of the

parent Ti film. From the TEM images a 350 nm thick NST layer was formed

from a 500 nm thick Ti film, in agreement with SEM imaging. It is postulated

that this reduction in thickness results from shrinkage of the hydrated titania

gel layer and dissolution due to finite solubility of Ti in the aqueous H2O2

solution. The NST/ T-SiO2 interface is sharp with no cracks detected. This

suggests that NST pads formed are strongly adhering to the T-SiO2

substrate. Since thickness of Ti films deposited using microelectronics

process tools can be controlled accurately to a few tens of nanometers, and

by taking into account shrinkage of titania gels, ultra-thin and porous NST

features could be integrated into devices using this technique.

Cross-sectional TEM studies also reveal features that are not apparent in

the SEM investigations. Figures 2.10 (b) and (c) show that lateral oxidation

of Ti under the P-SiO2 mask layer occurred. In both cases, oxidation

progressed ~580 nm into Ti films from the edge of the P-SiO2 mask. The

lateral oxidation rate under the experimental conditions used is estimated to

be ~2.8 nm/min. To investigate the crystal structure of the NST/Ti interface,

41

selected area electron diffraction (SAED) studies were carried out in regions

labeled (1), (2) and (3). The corresponding diffraction patterns are shown in

Figures 2.10 (d), (e) and (f), respectively. Electron diffraction pattern

obtained from region (1) suggests the presence of polycrystalline anatase.

The diffraction pattern acquired from region (2) is similar to that obtained

from (1). However, the intensity of ring patterns is noticeably reduced, which

suggests a higher proportion of amorphous phase in regions of NST nearer

to the unreacted Ti. The ring diffraction pattern indicates that titania formed

is crystalline and that nanocrystals of anatase are randomly oriented in the

walls of the pores. These observations are in agreement with XRD studies.

The diffraction patterns are indexed accordingly as labeled in Figures 2.10

(d) and (e). Energy dispersive X-ray spectroscopy studies at regions (1) and

(2) show that elemental compositions of these regions are similar. These

results confirm that TiO2 is formed in these regions after annealing. However

the proportion of the crystalline titania is lower in region (2). As expected,

the diffraction pattern of region (3) is typical for that of polycrystalline Ti

consisting of discrete rings with clustered elongated spots which suggests

small grains with multiple orientations46.

The presence of a completely amorphous titania phase had been

reported at the interface between Ti and porous titania formed by

anodization of thick Ti foils followed by annealing treatment. This amorphous

42

titania phase has been suggested to confer electrical isolation of the

semiconducting titania phases from Ti in some gas sensing devices47.

However a direct comparison with our results is not possible as we used a

different method to form the titania layer. Nevertheless, SAED results

indicate a fully amorphous phase is not present in our samples.

Figure 2.11. Crack formation near the Ti/NST interface. Cross-sectional image of sample (a) before and (b) after annealing.

Cross-sectional TEM also reveals cracks under the P-SiO2 mask layer

near the interface between the NST and unreacted Ti film. Such cracks were

observed on all three TEM samples studied. Two examples of crack

43

formation are shown in Figures 2.10 (b) and (c). Cracks are observed along

the P-SiO2/NST interface and continue near the NST/Ti interface before

being arrested ~150 nm from the NST/ T-SiO2 interface. It is noted that

cracks may be formed due to shrinkage during drying in air after oxidation or

during the annealing treatment. Figures 2.11 (a) and (b) show SEM images

of samples cross-sectioned by FIB milling before and after annealing. These

samples have been milled using a high beam current. Redeposition of

debris causes the pores of NST to be filled up. Cracks are observed in both

samples which indicate that crack formation occurs prior to thermal

annealing. It is postulated that cracks are formed due to shear stresses near

the NST/Ti interface as a result of shrinkage of the hydrated titania gel

during drying in air. However, no cracks were observed along the NST/T-

SiO2 interface. Also the NST under the P-SiO2 mask is dense with relatively

less porosity.

44

2.5 Conclusions

In conclusion, a method of forming integrated and patterned NST layers

for microsystems applications have been demonstrated. Using Ti thin films

patterned below a threshold area, crack formation on NST layers can be

eliminated. Crack formation in NST layers occurs primarily during drying and

is attributed to stresses associated with shrinkage. NST layers formed have

sponge-like morphology with pore diameter and wall thickness of about 50

nm - 200 nm and 25 nm - 50 nm, respectively. As-formed NST is largely

amorphous but transformed to anatase upon annealing at 300 °C. Pore size

of NST formed from bulk Ti foils is smaller that on Ti thin films. TiO2 grains in

NST on both Ti foils and Ti thin films have a random orientation.

45

2.6 References

1. F. Heidenau, F. Stenzel, G. Ziegler, Key Eng. Mater., 192-1, 87 (2000).

2. D. M. Brunette, P. Tengvall, M. Textor, P. Thomsen, Titanium in

medicine: materials science, surface science, engineering, biological responses and medical applications. Springer-Verlag, New York, (2001).

3. T. J. Webster, R. W. Siegel, R. Bizios, Biomaterials, 20, 1221 (1999).

4. T. Paunesku, T. Rajh, G. Wiederrecht, J. Maser, S. Vogt, N. Stojićević, M. Protić, B. Lai, J. Oryhon, M. Thurnauer, G. Woloschak, Nat. Mater., 2, 343 (2003).

5. M. Grätzel, Nature, 414, 338 (2001).

6. T. Sano, N. Iguchi, K. Lida, T. Sakamoto, M. Baba, H. Kawaura,

Applied Physics Letters, 83, 4468 (2003).

7. M. P. Harold, C. Lee, A. J. Burggraaf, K. Keizer, V. T. Zaspalis, R. S. A. Delange, Mater. Res. Soc. Bull., 19, 34 (1994).

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Nelli, C. Perego, L. Sangaletti, Adv. Mater., 8, 334 (1996).

9. O. K. Varghese, D. W. Gong, M. Paulose, K. G. Ong, E. C. Dickey, C. A. Grimes, Adv. Mater., 15, 624 (2003).

10. P. I. Gouma, M. J. Mills, K. H. Sandhage, J. Am. Ceram. Soc., 83,

1007 (2000).

11. G. P. Burns, Journal of Applied Physics, 65, 2095 (1989).

12. C-. C. Ting, S-. Y. Chen and D-. M. Liu, Journal of Applied Physics, 88, 4628 (2000).

13. S. H. Oh, D. J. Kim, S. H. Hahn and E. J. Kim, Materials Letters, 57,

4151 (2003).

14. H. Tang, K. Prasad, R. Sanjinés and F. Lévy, Sensors and Actuators B, 26-27, 71 (1995).

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15. C. J. Brinker and G. W. Scherer, Sol-Gel Science: The Physics and

Chemistry of Sol-Gel Processing, Academic Press, San Diego (1990).

16. J. Livage, Current Opinion in Solid State and Materials Science, 2, 32

(1997).

17. A. S. Holmes, R. R. A. Syms, M. Li and M. Green, Applied Optics, 32, 4916 (1993).

18. H. Krug, N. Merl and H. Schmidt, Journal of Non-Crystalline Solids,

147, 447 (1992).

19. R. L. Roncone, L. A. Wellerbrophy, L. Weisenbach and B. J. J. Zelinski, Journal of Non-Crystalline Solids, 128, 111 (1991).

20. T. Brylewski and K. Przybylski, Applied Superconductivity, 1, 737

(1993).

21. H. Kozuka, Journal of The Ceramic Society of Japan, 111, 624 (2003).

22. K. Kajihara, K. Nakanishi, K. Tanaka, K. Hirao and N. Soga, Journal

of The American Ceramic Society, 81, 2670 (1998).

23. M. C. Carotta, M. Ferroni, V. Guidi, G. Martinelli, Advanced Materials, 11, 943, (1999).

24. E. Halary, G. Benvenuti, F. Wagner ad P. Hoffmann, Applied Surface

Science, 154-155, 146 (2000).

25. L. Gao, Q. Li, Z. Song and J. Wang, Sensors and Actuators B, 71, 179 (2000).

26. C. J. Barbe, F. Arendse, P. Comte, M. Jirousek, F. Lenzmann, V.

Shklover and M. Gratzel, Journal of The American Ceramic Society, 80, 3157 (1997).

27. R. J. Collins, H. Shin, M. R. Deguire, A. H. Heuer and C. N. Sukenik,

Applied Physics Letters, 69, 860 (1996).

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28. K. Koumoto, S. Seo, T. Sugiyama and W. S. Seo, Chemistry of Materials, 11, 2305 (1999).

29. M. Bartz, A. Terfort, W. Knoll, W. Tremel, Chemistry – A European

Journal, 6, 4149 (2000).

30. A. Kumar, H. A. Biebuyck and G. M. Whitesides, Langmuir, 10, 1498 (1994).

31. J. M. Wu, S. Hayakawa, K. Tsuru, A. Osaka, Scripta Materialia, 46,

101 (2002).

32. P. Tengvall, “Titanium-hydrogen peroxide interaction with reference to biomaterial applications”, PhD thesis, Department of Physics and Measurement Technology, Linköping University, Sweden (1989).

33. P. Tengvall, I. Lundstrom, L. Sjoqvist, H. Elwing, L. M. Bjurstein,

Biomaterials, 10, 166 (1989).

34. P. Tengvall, H. Elwing, I. Lundstrom, Journal of Colloid and Interface Science, 130, 405 (1989).

35. S. Nishiguchi, S. Fujibayashi, H. M. Kim, T. Kokubo, T, Nakamura,

Journal of Biomedical Materials Research A, 67A, 26 (2003).

36. S. Suresh, L. B. Fruend, Thin Film Materials: Stress, Defect Formation and Surface Evolution, Cambridge University Press, Cambridge (2003), p. 222.

37. J. F. Moulder, W. F. Stickle, P. E. Sobol, K. D. Bomben, Handbook of

X Ray Photoelectron Spectroscopy: A Reference Book of Standard Spectra for Identification and Interpretation of XPS Data, Physical Electronics Inc., Minnesota (1995).

38. T. Choudhury, S. O. Saied, J. L. Sullivan, A. M. Abbott, Journal of

Physics D - Applied Physics, 22, 1185 (1989).

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40. H. Z. Zhang, M. Finnegan, J. F. Banfield, Nano Letters, 1, 81 (2001).

48

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Scripta Materialia, 48, 23 (2003).

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Publishing Company, Massachusetts (1959).

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49

Chapter 3: Kinetics of reaction between Ti films and

aqueous hydrogen peroxide

3.1 Introduction

It is widely known that titanium has good corrosion resistance. This

property is attributed to the presence of a compact thin film of TiO2. In air, a

thin TiO2 layer readily forms on a pristine Ti surface due to the large

corresponding reduction in the free energy, ΔG=-203.8 kCal/mol1. The

overall reaction can be represented as follows:

Ti + O2 ⇒ TiO2 ………. Eqn. 3.1

In an aqueous environment, oxidation of a pristine Ti surface has been

explained on the basis that the Ti/TiO2 potential is more negative than the

potential of the hydrogen electrode2. The reaction between Ti surfaces with

aqueous H2O2 however, is more complicated and involves the formation

intermediate species such as Ti-peroxo compounds which takes the form of

gel-like structures3-5.

The oxidation of Ti surfaces in aqueous H2O2 solution had, over many

years, been studied by numerous research groups. This interest lies in the

fact that H2O2 released in biological systems significantly reduced the

50

corrosion resistance of titanium implants and results in the accumulation of

titanium ions in adjacent tissues. In contrast, titanium implants in H2O2-free

aqueous environment, in-vitro, have high corrosion resistance and,

correspondingly, release lower amounts of Ti ions3. From a technological

stand point, these studies have contributed to the development of better

performance implants.

Despite these studies, gaps remain in our understanding of Ti surface

oxidation in aqueous solution. It is now generally accepted that two layers of

surface oxide are formed when a Ti surface is exposed to an aqueous

environment. The oxide layer adjacent to the metal is dense and acts as a

barrier to further corrosion while the second layer, adjacent to the aqueous

solution, is porous and does not impede the corrosion process. Addition of

H2O2 results in a significant release of Ti ions into the solution and growth of

the porous titania layer. Healy and Ducheyne6,7 suggested that thickness of

these two layers is about the same. These researchers found that oxide

growth decreased significantly after the layer reached a thickness of ~7.5

nm and suggested that oxidation kinetics can be explained on the basis of a

limiting oxide thickness in which electric-field assisted transport of metal ions

into the oxide is the rate limiting step. Using results obtained from

electrochemical impedance spectroscopy techniques, Pan et. al.8,9

calculated that the dense-TiO2 layer is much thinner than the porous layer.

51

However, Pan et. al.8,9 was not able to determine the rate limiting step

controlling growth of the porous titania layer but alluded to the study of

Tengvall. Tengvall3 had proposed that the rate limiting step controlling

growth is dissolution of Ti ions into the solution. Recently Bearinger et.

al.10,11 used atomic force microscopy (AFM) coupled with electrochemical

techniques to elucidate the mechanism of oxide evolution. These

researchers concluded that growth of the titanium oxide layer occurs at the

metal/dense-TiO2 interface. Models proposed by Tengvall3 as well as

Ducheyne and Healy6,7, however, suggest growth of the total oxide layer

occurs at the oxide-solution interface – see Figure 3.1.

Another gap in the understanding of aqueous Ti oxidation is the effect of

the underlying Ti substrate. No systematic study has been carried out on the

effect of microstructure of Ti substrates on oxidation kinetics. It is widely

known that grain boundaries are regions of high energy and, therefore,

preferentially corrode. In addition, grain boundary diffusion is faster than

bulk diffusion. Hence, grain size of the Ti substrate would be expected to

have a large effect on the rate of oxidation. In addition, for thin Ti films,

kinetics of the oxidation process may be affected by thickness of the

underlying Ti layer. Healy and Ducheyne6,7 performed their experiments on

sputtered Ti thin films while Pan et. al.8,9 and Bearinger et. al.10,11 used bulk

52

Ti rods and discs, respectively. However, no prior work has investigated the

effect, if any, of the underlying Ti microstructure.

Figure 3.1. Schematic, based on current understanding, of processes taking place during reaction of Ti surfaces with aqueous hydrogen peroxide

solution. (a) In Bearinger’s model growth occurs at the interface between dense and hydrated Ti oxide layers due to mass transport of H2O2

molecules across the oxide layer. (b) In the model proposed by Tengvall as well as Healy and Ducheyne, the oxide growth occur at the hydrated Ti oxide layer/solution interface due to mass transport of Ti ions across the

hydrated Ti oxide layer. ( represents flux of Ti from the substrate) TiJ

53

The aim of this chapter is two-fold: First, to elucidate the mechanism

controlling growth of the porous titanium oxide layer and, second, to study

the effect of Ti film microstructure on oxidation kinetics. Since the overall

goal of this research is to integrate NST into microsystems, focus is placed

on Ti thin films rather than bulk Ti.

3.2 Experimental procedure

Oxidation kinetics of Ti films was investigated by monitoring change in

resistance of patterned Ti films deposited on glass substrates. In this

chapter, all Ti films have been deposited using electron beam evaporation.

The effect of thickness and average grain size of Ti films on the oxidation

kinetics was studied. Other parameters investigated include temperature

and concentration of hydrogen peroxide solution.

A schematic of the set-up used to monitor Ti oxidation kinetics is shown

in Figure 3.2 (a). Patterns of Ti lines were deposited on glass substrates

using the lift-off technique. Glass substrates were degreased by soaking

consecutively in acetone, isopropyl alcohol and deionized (DI) water, with

ultrasonic agitation for 5 min each. These substrates were then baked at

120 °C for 10 min. Photoresist was spun, exposed with patterns and

developed. A thin Ti film was deposited on substrates using electron-beam

54

evaporation. To reduce oxygen incorporation, Ti source was evaporated for

about 5 min prior to actual deposition on glass substrates. During this time,

Ti was evaporated on a shutter that lies between the glass substrate and the

source. After 5 min, the shutter was removed and Ti was evaporated on the

substrate. Pressure in the chamber was about 1 - 2 x 10-7 Torr during

deposition. Thickness of Ti films deposited was estimated using a quartz

crystal monitor. Chamber pressure at the end of a deposition run was lower

than that at start. After deposition, substrates were soaked in acetone

overnight. For oxidation experiments, the samples were connected as

shown in Figure 3.2 (a). Oxidation was done at various temperatures. The

apparatus was heated using a hot plate and temperature maintained to

within ±3 °C of that required. Samples were dipped into the reaction bath

only when the temperature stabilized. Resistance of Ti film was recorded

every 10 or 15 seconds intervals. For each condition, three samples were

oxidized and the average calculated.

55

Figure 3.2. Schematic of (a) apparatus for oxidation kinetics experiments and (b) Ti line on sample.

Phase transformation and chemical reactions in solids often result in

changes in measurable physical properties. Experimental quantification of

these changes enables a reliable and accurate estimate of the extent the

transformation has occurred or the progress of reaction. The use of

electrical resistance as a parameter to follow the extent of chemical

reactions involving thin films had been investigated by many researchers in

the microelectronics industry12-15. Tu et. al. used this technique to study

crystallization induced by thermal annealing of amorphous co-evaporated

transition metal-silicon13,14 thin films while Howard et. al studied kinetics of

56

reaction of transition metal-aluminum thin films15. As mentioned previously,

oxidation of Ti results in the consumption of conducting Ti and formation of a

two layered oxide that is, electrically, highly insulating. Hence by measuring

the electrical resistance of the residual Ti film, kinetics of the oxidation

reaction can be monitored. The leakage current across the solution is

expected to be small since deionized water was used. Nevertheless, the

resistance of DI water is expected to decrease during oxidation because of

dissolution of the hydrated titania gel layer.

Following the analysis of Tu and Howard et. al.12-15, a measure of

progress of the reaction called the extent of reaction, Xt, can be defined as:

( ) ( )( ) ( )0

0

FilmFilm

FilmFilmt RR

RtRX−∞−

= ………. Eqn. 3.2

In Eqn. 3.2, and ( )0FilmR ( )∞FilmR are resistance of films at start and the

end of the reaction while ( )tRFilm is resistance of films at time t. It follows

from Eqn. 3.2 that Xt at start and end of the reaction is 0 and 1, respectively.

Kinetics of the oxidation reaction can be followed by plotting Xt against time.

57


3.3.1 Characterization of Ti films and calibration of apparatus

The measurement of resistance of the residual Ti film to monitor extent

of Ti oxidation in an aqueous environment has not been reported previously.

Therefore, careful calibration of the oxidation apparatus was carried out to

ensure resistance measured was due to current through the Ti film and not

from leakage current through the DI water. To calibrate the apparatus, Ti

films of various thickness were evaporated at 2.5 Å/s. Ti films used for

calibration were identical to those used to study the effect of film thickness

on oxidation kinetics. In the following, results of characterization of Ti films

used in the calibration are presented.

Characterization of Ti films for calibration

During evaporation, thickness of Ti films deposited was measured using

a quartz crystal monitor. The thickness of evaporated Ti film samples was

further verified using x-ray reflectivity experiments after evaporation. Incident

angle in reflectivity experiments was fixed at 3 degree. Reflectivity spectra

are shown in Figure 3.3 (a). The thickness, Δ , of a Ti film sample was

calculated using the formula:

( )( )ii

ii nnθθ

λsinsin2 1

1

−−

=Δ+

+ ………. Eqn. 3.3

58

where λ, and in iθ are the wavelength of radiation used (1.54056 Å for

CuKα radiation), the ith order of a reflection and the corresponding diffracted

angle at the peak of the ith order, respectively. Thicknesses of films

determined using X-ray reflectivity are close to those recorded using quartz

crystal monitor, Figure 3.3 (b). Digression between the two sets of thickness

values is greater as thickness increases. On average, film thickness

determined using X-ray reflectivity is 95% of that recorded using a quartz

crystal monitor. In the following, nominal film thickness recorded using the

quartz crystal monitor is used.

Figure 3.3. Plots of (a) reflectivity spectra of various films and (b) correlation between thickness determined using X-ray reflectivity and

quartz crystal monitor techniques.

59

Average grain size was determined using X-ray broadening. Figure 3.4

(a) shows x-ray diffraction spectra taken in Bragg-Brentano geometry of Ti

films of various thicknesses. Ti films are highly textured in the (0002) plane,

similar to a previous report22. To determine average grain size, reflections

from (0002) planes were used. Average grain size was determined using

Scherer’s relation22:

θλ

cosBCd = ………. Eqn. 3.4

where C, λ, B and θ are a constant with value of 0.9, the wavelength of

radiation used (1.54056 Å for CuKα radiation), true full width at half

maximum intensity and the diffracted angle, respectively. The technique

above gives average grain size perpendicular to the surface22. To obtain the

true full width at half maximum intensity B (in radians), Warren’s method

was employed22. In this method, B is determined from:

222SM BBB −= ………. Eqn. 3.5

where and are the measured full widths at half maximum

intensity of the sample and a standard, respectively. The standard used was

a bulk Ti sheet obtained from a commercial supplier that has been annealed

for a few hours at elevated temperatures and has an average grain size of a

few micrometers.

MB SB

60

Figure 3.4. Plots of (a) x-ray diffraction spectra taken in Bragg-Brentano geometry of Ti films and (b) Variation of average grain size with

thickness. (Evaporation rate 2.5 A/s)

Figure 3.4 (b) shows variation of average grain size with thickness.

Average grain size increases with thickness from 25 to 100 nm. For 100,

150 and 200 nm thick films, the average grain size is ~25 nm.

Calibration of apparatus

Calibration was done by measuring resistance of films, while immersed

in DI water, at various temperatures. The corresponding variations of

resistivity with thickness and temperature were then calculated. From the

61

variation of resistivity with temperature, the temperature coefficient of

resistivity (TCR) was calculated. Resistance values of samples in air and in

DI water were found to be similar, which suggest current leakage through DI

water is negligible. In the following, calibration results obtained from

samples immersed in DI water are presented and compared to those in the

literature.

The resistance measured using the apparatus shown in Figure 3.2 (a)

include resistance of the film, , and the contact resistances, . The

total resistance measured, R, can be written as:

FilmR RContacts

FilmContacts RRR += ………. Eqn. 3.6

By definition:

ALRFilmρ

= ………. Eqn. 3.7

where ρ, L and A are the resistivity, length and cross-sectional area

perpendicular to current flow, respectively. To estimate , R was

measured at a few L values. This was done by placing one terminal at A and

the other at D, and measuring R. The terminal at A was then moved to B

and R measured, and so on to C. A plot of R versus L was then made. As

L→ 0, from Eqn. 3.6 → 0. Hence, at L=0,

ContactsR

R RRFilm Contacts= . of 4 ContactsR

62

samples was measured and found to be similar. The average value is ~38 Ω

and is used as an estimation of the contact resistance in all samples.

Resistivity of Ti films deposited was calculated from measured in DI

water using Eqn. 3.8, which takes into account the effect of thermal

expansion

FilmR

16:

( ) (( RTl

wtTK

2931293

−+⎟⎠⎞

⎜⎝⎛= αρ )) ………. Eqn. 3.8

In Eqn. 3.8, w, l and t are the width, length and thickness of Ti film

patterns, respectively, and R is the resistance measured. The linear thermal

expansion coefficient, α, is taken as 9.7 μ°C-1 17. Figures 3.5 (a) and (b)

show variation of resistivity with thickness and temperature respectively.

Resistivity of Ti films deposited decreases with thickness. In addition, values

determined in the present research are comparable to those obtained from

similarly prepared evaporated Ti thin films18,19.

63

Figure 3.5. Variation of resistivity with (a) thickness and (b) temperature of Ti films. (Evaporation rate 2.5 A/s)

The reduction in resistivity with thickness in metal films has been

investigated by many researchers16,19-21. Mayadas and co-workers modeled

resistivity in polycrystalline films as a combination of 3 types of scattering –

(1) isotropic scattering due to point defects, (2) scattering due to external

surfaces and (3) scattering due to grain boundaries18,19 – and suggested the

last contribution to be dominant. Singh and Surplice observed that gettering

of gaseous impurities such as oxygen in thinner films results in increased

resistivity19. Igasaki and Mitsuhashi16 suggested that reduction in resistivity

with thickness increase is due to grain boundary scattering.

64

Resistivities of Ti films increase linearly with temperature and are almost

parallel for the range of temperature investigated. The rate of resistivity

increase with temperature, ⎟⎠⎞

⎜⎝⎛

dTdρ , ranges from 0.214 to 0.246 μΩcmK-1 and

agrees with values reported in prior studies16,23,24, see Figure 3.6. The

similarity between resistivity and ⎟⎠⎞

⎜⎝⎛

dTdρ values obtained in the present

research and others in the literature can be regarded as an indication of the

experimental validity and accuracy of resistance values measured from Ti

films immersed in DI water.

Figure 3.6. Comparison of TCR obtained from various studies. Data of Igasaki and Mitsuhashi, and Wasilewski were obtained from 300 nm

films and bulk Ti, respectively.

65

3.3.2 Effect of film thickness

Oxidation of Ti films of various thicknesses occurs through nucleation

and growth of the oxide phase. Depending on initial thickness of the Ti film,

oxidation behavior of films during growth can be divided into two sub-stages.

Plots of and XFilmR

R

t with time immersed in aqueous hydrogen peroxide

solution at 80 °C are shown in Figures 3.7 (a) and (b), respectively. Figure

3.7 (c) is a schematic of and XFilm t plots showing different oxidation

behavior for thick and thin films. In the first stage, resistance increases very

gradually while in the second, the rate of resistance increase is significantly

greater. Bearinger et. al.10,11 reported that oxidation of Ti surfaces occurs by

nucleation of titanium oxide domes followed by subsequent growth and

coalescence of domes to form a continuous porous layer. It is postulated

that the two stages observed in the present study correspond to nucleation

and growth processes observed by Bearinger et. al.10,11. Resistance of films

at the end of a reaction, ( )∞FilmR , decreased with increasing thickness. This

is because more Ti ions will dissolve into DI water at the end of a reaction

for thicker Ti films. However, the decrease is low; ( )∞FilmR of 25 and 200 nm

thick films are 1.7 and 1.0 MΩ, respectively.

66

Figure 3.7. Variation of (a) resistance and (b) extent of reaction, Xt, with time immersed in H2O2 solution, and (c) schematic of Xt for thin and thick

films.

During nucleation, little Ti is consumed (or, alternatively, hydrated

titanium oxide is formed) and the rate of change of Xt, dt

dX t , increases very

gradually. For 25 to 50 nm films, nucleation time differs by 110 s. In contrast,

duration of nucleation for 100, 150 and 200 nm is comparable – 362, 344

and 377 s, respectively. Grain size of 100, 150 and 200 nm films is similar,

Figure 3.4 (b), but decreased significantly as thickness decreased from 100

to 25 nm. These observations suggest that nucleation time increases with

grain size.

During the growth stage, variation of Xt with time for 25 and 50 nm Ti

films differs significantly from those of thicker ones. For 25 nm and 50 nm

films, Xt increases linearly with time until the Ti film is fully consumed. For

67

films with thickness 100 nm and greater, variation of Xt is more complicated.

Early in the growth stage before ttrans, Xt increases at a constant rate,

( )dt

transt ttdX <

( )

, which is dependent on grain size - in a similar manner as the

case of duration of nucleation. A plot of grain size and dt

transt ttdX <

( )

against

film thickness is given in Figure 3.8. From this figure, it is clear that structure

of Ti films affect oxidation kinetics. Films with smaller grains have greater

dttranst ttdX < . The effect of grain size on reaction kinetics will be investigated

in greater detail in Section 3.3.3. This dependence on grain size suggest

that dt

t ( )dt

transtdX of 25 and 50 nm films and ttdX < of thicker films are

controlled by a reaction at the interface between the dense-TiO2 to the

hydrated titanium oxide layer. Interface controlled reactions generally are

linear with time25. Dependence on grain size can be rationalized as follows.

Assuming, Ti species diffuse out through the dense-TiO2 layer for reaction to

proceed, flux of Ti species is greater for smaller grains due to grain

boundary diffusion. The increased flux subsequently results in increased

growth rate.

68

Figure 3.8. Variation of grain size and ( )

dtttdX transt < with thickness.

(Evaporation rate 2.5 A/s)

For 100, 150 and 200 nm films, the growth stage can be divided into sub-

stages I and II, see Figure 3.7 (c). Upon reaching a critical value, , at

the subsequent rate of increase of X

transtX

transt t, ( )

dtttdX transt >

trans

t

, decreases

appreciably. for 100, 150 and 200 nm films are 74, 65 and 46 %,

respectively, and corresponds to consumption of 74, 97 and 92 nm thick Ti.

These values as well as , time at which transition in rates took place, are

listed in Table 3.1. For 100, 150 and 200 nm films,

tX

trans

( )dt

ttdX transt > values are

69

similar – ~0.19 s-1. For these films, plots of Xt (and ) for t are

parallel. The similarity of

FilmR t> trans

( )dt

transt ttdX > values for thick films suggests a

common growth mechanism is occurring in these samples. According to the

models of Tengvall3, and Ducheyne and Healy6,7, Ti ions must diffuse

across the hydrated titanium oxide layer to grow. Alternatively, the model of

Bearinger et. al.10,11 requires that oxidants diffuse through the hydrated

titanium oxide layer for growth to take place. One possible mechanism

controlling growth in sub-stage II is mass transport across the growing

hydrated titanium oxide layer. As the layer increase in thickness, a critical

thickness is reached at time t after which mass transport of reactants

across the oxide layer is controlling growth rather than reaction at the

interface. Hence, rate change in sub-stages I and II is probably due to a

corresponding change in mechanism controlling growth.

trans

Table 3.1. Oxidation parameters for Ti films of various thicknesses.

Film thickness

(nm)

Duration of

nucleation (s)

transt (s)

transtX

(%)

(dt

ttdX trant <

(s-1)

( )dt

ttdX transt >

(s-1)

Thickness consumed

at transition

(nm) 25 110 None None 0.61 None None 50 220 None None 0.44 None None

100 362 538 74 0.42 0.18 74 150 344 548 65 0.32 0.18 97 200 377 515 46 0.31 0.20 106

70

Transition from sub-stages I and II occur with increasing thickness of Ti

consumed. Assuming that all Ti consumed goes into forming a hydrated Ti

oxide gel layer, this implies that thickness of the gel layer at which transition

from sub-stages I to II occurs increases with parent Ti thickness. One

possible explanation is porosity of the hydrated Ti oxide gel layer. It is

speculated that hydrated Ti oxide gel layer formed from 200 nm Ti film has a

more open porous structure resulting in diffusion controlled growth at greater

oxide thickness. To prove this hypothesis unambiguously, in-situ probing of

the hydrated Ti oxide gel layer is required. Nevertheless an estimation of

compactness of the hydrated Ti oxide gel layer can be obtained from its

structure after annealing. Figure 3.9 (a) and (b) show SEM images of NST

formed on 150 and 500 nm films after annealing. Clearly, NST formed from

the 500 nm films is more porous. No discernable difference can be observed

between corresponding images of 100 and 200 nm Ti films.

Figure 3.9. NST from (a) 150 and (b) 500 nm Ti films.

71

The nucleation and early growth of hydrated titanium oxide layer can be

described using the kinetic law in the theory of transformation kinetics25.

According to the kinetic law, the percentage of volume transformed (Xt) at

any time t ( ) at time t is given by( )

t

tX t25:

( ) )exp(100100 nt kttX −−= ………. Eqn. 3.9

where k and n are constants. The exponent n is a parameter related to the

mode of transformation13,14,25. Application of this law assumes growth is

linear with time and is valid during the early stage of diffusion controlled

growth. Bearinger et. al.10,11 reported that hydrated titanium oxide domes

grow linearly with time during nucleation. During the growth stage, rate of

increase of Xt is constant indicating that hydrated titanium oxide layer from

25 and 50 nm Ti films grows linearly, see Figure 3.7 (b). In contrast, growth

from 100, 150 and 200 nm films is parabolic overall and involves diffusion

after . Hence, the kinetic law in Eqn. 3.9 was applied only to oxidation

data of 25 and 50 nm films in Figure 3.5 (b). A good fit of the experimental

data to Eqn 3.9 is shown in Figure 3.10. The n values extracted for 25 and

50 nm films are 3.23 and 3.96, respectively. According to the theory of

transformation kinetics, for transformations constrained to two dimensions,

as the case in this study, these values suggest increasing nucleation rate.

trans

72

Figure 3.10. Fitting of kinetic law to experimental data for 25 and 50 nm films

In conclusion, the effects of Ti film thickness on oxidation kinetics can be

summarized as follows. Oxidation of Ti films of various thicknesses occurs

through nucleation and growth stages. dt

dX t is low during nucleation but

increased significantly during growth. For 25 and 50 nm thin films, dt

tdX is

constant during the growth stage. For thicker films, however, the growth

stage can be subdivided into sub-stages I and II. In sub-stage I, growth

occurs due to reaction at the hydrated titanium oxide/dense-TiO2 interface

73

( )dt

transt ttdXwith a constant rate < . Growth in sub-stage I is affected by grain

size of the underlying Ti film. After time t (or in sub-stage II), growth

occurs at a constant but significantly reduced rate

trans

( )dt

transt ttdX > and is

controlled by mass transport through the hydrated titanium oxide layer. Rate

of oxide growth in sub-stage II is independent of grain size but is affected by

porosity of the hydrated titania gel layer.

3.3.3 Effect of grain size

Grain size was found to have a significant effect on oxidation kinetics. To

study the effect of grain size, 50 nm thick Ti films were evaporated at

various deposition rates. Similarly, thickness was verified using X-ray

reflectivity and was found to be within 95% of the nominal values. Films are

highly textured in the (0002) plane as shown in Figure 3.11 (a). Average

grain size was calculated from (0002) peaks using the Scherer’s relation,

Eqn. 3.4, and was found to increase with deposition rate, Figure 3.11 (b).

This trend could be explained on the basis of adatom mobility. As

evaporation rates increase, the mobility of Ti adatoms arriving on the

surface correspondingly increases26,27. More mobile Ti adatoms are able to

migrate farther on a surface and consequently result in increased grain size.

74

Figure 3.11. Plots of (a) X-ray diffraction spectra in Bragg-Brentano geometry for 50 nm films deposited at various rates and (b) average grain

size calculated using Scherer’s relation.

As grain size increases, the rate of change of Xt decreases. Plot of Xt

with time immersed in aqueous hydrogen peroxide solution at 80 °C for the

various films is shown in Figure 3.12 (a). Distinct nucleation and growth

stages are observed. Duration of nucleation increases with grain size.

During growth, the oxidation of Ti proceeds with a constant dt

dX t until

completion. With increasing grain size, dt

dX t was found to decrease, Figure

75

3.12 (b). Rate of change of Xt, dttdX , of samples deposited at 0.2 A/s were

large and not enough data points could be recorded. Results are

summarized in Table 3.2.

Figure 3.12. Variation of (a) extent of reaction, Xt, with time immersed in

H2O2 solution and (b) grain size and dt

dX t with deposition rate.

Table 3.2. Oxidation parameters for 50 nm Ti films of various evaporation rates.

Evaporation

Rate (Å/s)

Grain Size (nm)

Duration of nucleation (s)

dtdX t

(s-1)

0.5 12.5 138 2.1 1.5 21.8 207 0.82 2.5 23.7 220 0.44 3.5 26.1 346 0.56 4.5 28.4 453 0.54

76

Results above indicate that grain size has a significant effect on kinetics

of reaction between hydrogen peroxide solution with 50 nm Ti films. The

effect of grain size could be the result of enhanced grain boundary diffusion

as discussed previously. For hydrated titanium oxide domes to nucleate, Ti

ions must be present on the surface of the dense-TiO2 layer. For the case of

films with small grain size, a higher concentration of Ti ions from the

substrate, due to enhanced grain boundary diffusion. The higher

concentration of Ti ions can affect the nucleation stage in two ways. First,

the nucleation density can be increased and second, the rate of growth of

nuclei will be increased. In both cases, duration of nucleation stage will be

increased.

Similarly, as grain size increases dt

dX t was found to increase. From the

preceding discussion in Section 3.3.2, it was deduced that growth rate for 50

nm films is controlled by reaction at the interface between the hydrated

titanium oxide layer and the dense-TiO2 layer. This assertion is further

supported here. Flux of Ti ions from the film would be expected to be higher

for smaller grain size as a result of grain boundary diffusion. Consequently,

the rate of reaction at the interface would be accelerated, resulting in

increased dt

tdX . This is in agreement with results from Section 3.3.3.

77

3.3.4 Effect of temperature

The effect of temperature can be modeled using the kinetic theory of

transformation. This theory is valid only during linear growth, which in the

present study, is strictly applicable to 25 and 50 nm thick films. Therefore, to

study the effect of temperature, 50 nm thick Ti samples evaporated at 2.5

Å/s were used. Thickness of samples was verified using X-ray reflectivity.

Plots of Xt with time immersed in aqueous hydrogen peroxide solution at

various temperatures are shown in Figure 3.13. As expected, the time to

reach completion of the reaction in shortened. With increasing temperature

duration of nucleation decreases while rate during growth, dt

tdX was found

to increase. Table 3.3. summarizes the effect of temperature on reaction

kinetics.

78

Figure 3.13. Variation of extent of reaction, Xt, with time immersed in H2O2 solution at various temperatures.

Table 3.3. Parameters for 50 nm Ti films oxidized at various temperatures.

Temperature

(°C) Duration of nucleation

(s) dtdX t (s-1)

51.5 795 0.25 60 600 0.29 70 375 0.29 80 220 0.44 90 177 0.68

79

The experimental data can be modeled by the kinetic law, Eqn. 3.9.

Fitting of the law to experimental data is shown in Figure 3.14 (a). n values

extracted range from about 4 to 7 which are in agreement with reported

values in the literature13,14,25. To obtain the temperature dependence of the

oxidation rate, a time τ can be defined at a particular reaction temperature at

which Xt(τ)=constant. Assuming an Arrhenius dependence, the relation

between τ and temperature T is given by:

⎟⎠⎞

⎜⎝⎛=

kTEaexp0ττ ………. Eqn. 3.10

In Eqn. 3.10, 0τ is a constant, k is the Boltzman’s constant, and Ea is

the apparent activation energy for the reaction. We follow the analysis of

Weiss et. al. and used ( ) 5.0=τtX to extract the apparent activation

energies14. Figure 13.4 (b) plots time, 5.0τ , (in logarithmic scale) at which

( ) 5.0=τTX against T1 - the exponent then gives

kTEa . The apparent

activation energy, , for the transformation was calculated to be 0.37 eV. aE

80

Figure 3.14. (a) Fitting of kinetic law to experimental data and (b) plot of 5.0τ

against T1 for 50 nm films oxidized at various temperatures.

The pathways for reaction between Ti(IV) ions with H2O2 are under

active investigation in the literature. Samuni suggested that the mechanism

involves reaction between more than one H2O2 molecules with one Ti (IV)

ion4. Recent theoretical work by Sever and Root postulated that energy

required to bind one H2O2 molecule to a Ti(IV) ion is about 0.17 eV5. As

discussed in Section 3.3.2, for Ti films of thickness less than 50 nm, kinetics

of oxidation is controlled by reaction of H2O2 molecule with Ti(IV) ions at the

hydrated titanium oxide/dense-TiO2 interface. The ratio of apparent

activation energy extracted to the energy required to bind one H2O2

81

molecule to one Ti(IV) ion is 2.2. Hence, it is very plausible that the

mechanism for the interface-reaction controlled kinetics during the growth

stage of 25 and 50 nm films as well as in sub-stage I of 100 – 200 nm films

involves reaction between two H2O2 molecules with one Ti (IV) ion. This

reaction will also occur in sub-stage II during oxidation of 100, 150 and 200

nm films. However, the growth rate in this stage will be controlled by mass

transport of species across the hydrated Ti oxide layer.

3.3.5 Effect of hydrogen peroxide concentration

The oxidation rate of Ti films increases with hydrogen peroxide

concentration as shown in Figure 3.15. For 30, 20 and 15 % (by volume)

hydrogen peroxide solutions, the reaction rate is large and only a few points

were able to be recorded. Nonetheless, it is clear from these plots that the

rate of reaction, dt

tdX, increases with hydrogen peroxide concentration.

Similarly the duration of nucleation is reduced with increasing hydrogen

peroxide concentration. Table 3.4. summarizes the effect of hydrogen

peroxide concentration on reaction kinetics.

82

Figure 3.15. Extent of reaction, Xt, with time immersed in H2O2 solution at 80 °C.

Table 3.4. Parameters for 50 nm Ti films oxidized at various H2O2 concentrations.

Hydrogen peroxide

concentration (volume %)

Duration of nucleation (s) dt

dX t (s-1)

10 220 0.44 15 97 2.57 20 61 3.21 30 42 5.54

83

The effect of hydrogen peroxide concentration suggests that the species

diffusing across the hydrated titanium oxide layer is hydrogen peroxide

molecules. This suggestion is borne out from Fick’s First Law which states

that:

xC

DJ OHOHOH ∂

∂−= 22

2222………. Eqn. 3.11

In Eqn. 3.11, , and denotes flux, diffusion coefficient

and concentration of hydrogen peroxide molecules, respectively. After

nucleation and after the oxide domes have coalesced to form a continuous

hydrated titanium oxide layer, a concentration gradient is set up,

22OHJ22OHD

22OHC

xOH

∂22

C∂,

across the hydrated titanium oxide layer. Assuming diffusion in the aqueous

solution is large, which is a reasonable in a solution, then concentration of

hydrogen peroxide molecules at the hydrated titanium oxide layer/solution

interface is equal to that in the bulk. Hence, the flux of hydrogen peroxide

molecules diffusing across the oxide layer increases with its bulk

concentration, in accordance with Eqn. 3.11. This explains the increase in

oxidation rate with increasing hydrogen peroxide concentration, Figure

13.15. This explanation also agree with the observation of growth occurring

at the hydrated titanium oxide layer/dense-TiO2 layer interface rather than at

84

the solution hydrated titanium oxide layer/solution interface as postulated by

Bearinger et. al. 10,11.

3.4 Phenomenological model of Ti oxidation in aqueous

hydrogen peroxide

From the above discussions, a phenomenological model of Ti oxidation

in aqueous hydrogen peroxide can be proposed, see Figure 3.16. Oxidation

of Ti surfaces occurs in two distinct stages, namely nucleation and growth.

Nucleation involves formation of hydrated titanium oxide domes on the

dense-TiO2 layer. These domes grow laterally and finally coalesce to form a

thin continuous hydrated titanium oxide layer. During the nucleation stage,

rate of consumption of the underlying Ti films is low. Growth of domes

occurs at the interface between the dense-TiO2 layer and the hydrated

titanium oxide layer by reaction between hydrogen peroxide molecules with

Ti ions.

Kinetics during the growth stage of thin (≤100 nm) Ti films is different

from that of thick ones. For thin films, growth occurs at a constant rate, dt

tdX ,

until the Ti substrate is fully consumed. In contrast, the growth of thick films

occurs in two distinct stages – sub-stages I and II. During sub-stage I, rate

85

( )dt

transtof growth, ttdX <

( )

, of the hydrated titanium oxide layer is controlled by

reaction between hydrogen peroxide molecules with Ti ions at the interface

between the dense and hydrated titanium oxide layers. Growth rate in sub-

stage I is sensitive to the structure of the underlying Ti substrate. As

reaction proceeds and thickness of the hydrated titanium oxide layer

increases, a transition thickness is reached at time t , at which mass

transport of hydrogen peroxide molecules across the hydrated titanium

oxide layer, rather than reaction at the interface, controls growth of the

hydrated titanium oxide layer. At , the rate of oxide growth,

trans

transtdt

transt ttdX > ,

decreases significantly. ( )

dttranst ttdX < is insensitive to parent Ti

microstructure but is affected by compactness of the hydrated titanium oxide

layer. A smaller transition thickness was observed for more compact of

hydrated titanium oxide layer formed from thinner Ti films.

86

Figure 3.16. (a) to (c) Schematics of processes occurring during oxidation of Ti thin films in aqueous hydrogen peroxide solution. (d)

Variations of extent of reaction, Xt, for thin and thick films; processes occurring at various stages of oxidation are labeled accordingly. ( represents Ti flux from the substrate due to grain boundary diffusion)

GBTiJ

87

3.5 Conclusions

In conclusion, the reaction kinetics of Ti thin films with hydrogen peroxide

solution has been investigated experimentally. Extent of reaction, Xt, was

monitored by measuring the resistance of the residual Ti film. It was found

that oxidation of Ti films occurs by nucleation and growth mechanism. Both

grain size and thickness of films affects the oxidation kinetics. Grain size

was found to affect both nucleation and growth stages. Films with finer grain

size have shorter nucleation period and higher growth rate.

During oxide growth, oxidation rate of films (25 and 50 nm) occurs at a

constant rate until the Ti films are completely consumed. For thicker films

(100, 150 and 200 nm films), growth rate decreases after a certain thickness

of porous titania has been formed. The oxide thickness at which the rate

reduction occurs is dependent on porosity of the oxide layer. Change in

oxidation rate is attributed to a change in the mechanism controlling growth

of the oxide layer. During the initial period of growth (sub-stage I) growth of

oxide is controlled by reaction of Ti species with hydrogen peroxide

molecules. During the later stage (sub-stage II), diffusion of hydrogen

peroxide molecules through the oxide layer is the rate controlling

mechanism. A phenomenological model of Ti oxidation in aqueous hydrogen

peroxide solution that takes into account the effect of grain size and

thickness of Ti films is proposed which explains the above observations and

is consistent with recent reports in the literature.

88

3.6 References

1. L. L. Sheir, Corrosion, Volume 1, Newnes-Butterworth, London (2000).

2. K. J. Vetter, Elektrochemische Kinetik, Springer, Berlin (1961), as cited by A. M. Shams el din, A. A. Hammoud, Thin Solid Films, 167, 269 (1988).

3. P. Tengvall, “Titanium-hydrogen peroxide interaction with reference

to biomaterial applications”, PhD thesis, Department of Physics and Measurement Technology, Linköping University, Sweden (1989).

4. A. Samuni, Journal of Physical Chemistry, 76, 634 (1972).

5. R. S. Sever, T. W. Root, Journal of Physical Chemistry B, 107, 4090

(2003).

6. K. E. Healy, P. Duchyene, Biomaterials, 13, 553 (1992).

7. K. E. Healy, P. Duchyene, Journal of Colloid and Interface Science, 150, 404 (1992).

8. J. Pan, D. Thierry, C. Leygraf, Journal of Biomedical Materials

Research, 107, 4090 (2003). 9. J. Pan, D. Thierry, C. Leygraf, Electrochimica Acta, 41, 1143 (1996).

10. J. P. Bearinger, C. A. Orme, J. L. Gilbert, Journal of Biomedical

Materials Research Part A, 67A, 702 (2004).

11. J. P. Bearinger, C. A. Orme, J. L. Gilbert, Surface Science, 490, 371 (2001).

12. S. P. Murarka, Metallization: Theory and Practice for VLSI and ULSI,

Butterworth-Heinemann, Boston (2001).

13. F. Nava, T. Tien, K. N. Tu, Journal of Applied Physics, 57, 2018 (1985).

14. B. Z. Weiss, K. N. Tu, D. A. Smith, Journal of Applied Physics, 59,

415 (1986).

89

15. J. K. Howard, R. F. Lever, D. A. Smith, P. S. Ho, Journal of Vacuum Science and Technology, 13, 68 (1976).

16. Y. Igasaki, H. Mitsuhashi, Thin Solid Films, 51, 33 (1978). 17. E. S. Greiner and W. C. Ellis, Transactions of the American Institute

of Mining and Metallurgical Engineers, 180, 657 (1949).

18. F. Huber, IEEE Transaction on Component Parts, CP-11, 38 (1964)

19. B. Singh and N. A. Surplice, Thin Solid Films, 10, 243 (1972).

20. A. F. Mayadas, M. Shatzkes, Physical Review B, 1, 1382 (1970).

21. A. F. Mayadas, R. Feder, R. Rosenberg, Journal of Vacuum Science and Technology, 6, 690 (1969).

22. B. D. Cullity, Elements of X-ray Diffraction, Addison-Wesley, Reading

(1959).

23. R. J. Wasilewski, Transactions of the American Institute of Mining and Metallurgical Engineers, 224, 5 (1962).

24. M. E. Day, M. Delfino, J. A. Fair, W. Tsai, Thin Solid Films, 254, 285

(1995).

25. J. W. Christian, The Theory of Transformations in Metals and Alloys, Pergamon Press, Oxford (1965).

26. C. R. M. Grovenor, H. T. G. Hentzell, D. A. Smith, Acta Metallurgica,

32, 773 (1984).

27. R. Messier, A. P. Giri, R. A. Roy, Journal of Vacuum Science and Technology A, 2, 500 (1984).

90

Chapter 4: Gas sensing using nanostructured TiO2

elements

4.1 Introduction

An electronic nose is, by definition, an intelligent chemical sensor array

system for odor classification. The concept of an electronic nose comes

about due to the increasing demand for electronic analogues of the human

olfactory system that are able to rapidly process, at relatively low-cost, odor

information for various applications ranging from food processing to

environmental monitoring and protection1. Similarly, electronic tongues are

counterparts of electronic noses that operate in liquid. In general, electronic

noses and tongues consist of sampling systems (for collecting samples),

arrays of chemical sensing elements (for detecting chemical species),

electronic circuitry (for transferring information collected by the sensing

elements and for heating sensing elements to the operating temperature)

and the data processing system. Among these components, chemical

sensing elements are undoubtedly the most critical as these largely dictate

the sensitivity and selectivity of the device — and which is the subject of this

chapter.

91

Chemical sensing elements may be categorized by the type of materials

used as either inorganic crystalline materials, organic materials or

biologically-inspired materials. Inorganic crystalline materials include

semiconductors such as metal oxides, silicon or carbon nanotubes.

Inorganic materials are particularly attractive as they hold the promise of

being easily integrated into microelectronic devices. Once material

integration has been achieved, it is then possible to leverage on the

enormous microelectronics fabrication infrastructure to produce electronic

noses and tongues in large volume at low-cost. Consequently, there has

been extensive effort towards integrating inorganic nanostructured materials

in devices. The mechanism of operation of a sensing element is the

transduction of interactions occurring between chemical species with its

surface to a measurable change in at least one property of the element. In

most cases, the measured property is resistance, capacitance or

temperature. Since nanostructured materials have large surface to volume

ratio, surface events result in a large change of a material property.

One approach to integrating nanostructured materials into devices is the

use of discrete nanostructures such as tubes or wires. Recent studies have

focused on single-walled carbon nanotubes2, silicon nanowires3, palladium

mesowire arrays4, metal oxide nanowires5,6 and polymeric nanowires7,8. As

a first step towards realization of electronic noses and tongues, the chemical

92

sensing properties of these materials have been characterized. However,

the challenge of locating and manipulating these nanostructures, fabricated

by “bottom-up” approaches, onto desired locations on chips impedes their

utilization in “real world” devices. To address this challenge, many

imaginative methods have been proposed such as electric9 and magnetic10

field induced or fluidic-assisted alignment together with surface patterning

techniques11-13, the fabrication of silicon nanowires from a macroscopic

silicon layer on an insulator14, self-organization or template synthesis of

nanostructures, and so on. Another approach involves the development of

processes to tailor, at the nanoscale, morphology of metal oxides, such as

titania15, zinc oxide16 and tin oxide17.

In this vein, we describe a new approach for fabricating micrometer-scale

sponge-like structures consisting of interconnected nanoscale walls or wires.

Configured as active elements of conductometric sensing devices, these

structures possess the ultra-high sensitivity of nanostructures while

benefiting from the enormous capabilities of micrometer-scale fabrication

that makes them amenable to integration into practical devices. We

demonstrate this approach by fabricating a gas sensor utilizing arrays of

NST pads as sensing elements on both Si and Kapton®. The latter is a

commercially available organic flexible substrate which is relatively low cost.

93

4.2 Gas sensing using nanostructured metal oxides

The use of metal oxides in electronic chemical sensing devices dates

back to the late 1950’s when Kiukkola and Wagner18 as well as Weissbart

and Ruka19 used ZrO2 as solid electrolytes. Since then, there have been

tremendous developments in metal oxide based sensors20. Titanium dioxide

(TiO2) has been widely studied for applications as sensing elements21-24. It

is attractive since its electrical conductivity and surface properties can be

modified by judicious use of surface activation and dopants25-28. TiO2 can

also be made porous and the pore structure controlled, allowing

functionalization with biomolecules29. In recent reports nanostructured

titania (NST) was used as a sensing element in ultra-sensitive sensors30,31.

Unlike conventional sensors using discrete nanotubes and nanowires, NST

has a sponge-like morphology with interconnected walls that act as

percolating channels for electron (hole) transport. This makes NST

particularly suitable as sensing elements. Finally, the approach described

has a high process yield and is compatible with current microchip

manufacturing practices and specifically photo-lithography which can place

an NST element at any predetermined locations on a chip.

94

4.3 Integration of nanostructured TiO2 as sensing elements

on silicon

Figure 4.1. Schematic process flow for fabrication of patterned integrated NST arrays.

Compared to other methods of forming ns-titania, the technique

developed for integrating NST is compatible with current microsystems

device manufacturing practices. As a demonstration of this compatibility, we

have fabricated gas sensor arrays with 20 μm ns-titania pads as sensing

elements. The detailed description for forming NST has been given in

Chapter 2. Here, only a brief description of the process to integrate NST as

95

sensing elements is given. The fabrication process of NST pad arrays is

summarized in Figure 4.1. A 1 μm thick SiO2 (T- SiO2) layer was thermally

grown on a 4 inch N-type Si(100) wafer at 1100 °C. A 500 nm thick Ti film

was then deposited using electron-beam evaporation. Subsequently, a 1 μm

thick SiO2 (P-SiO2) layer was deposited at 250 °C using plasma-enhanced

chemical vapor deposition (PECVD) using silane (SiH4) and nitrous oxide

(N2O) precursors. The patterns on the P-SiO2 layer were then etched using

CHF3 gas to expose Ti surfaces. These exposed Ti surfaces were then

oxidized by aging in aqueous 10 % hydrogen peroxide solution at 80 °C for

2.5 hrs. The samples were annealed at 300 °C for 8 hrs. Subsequently

Ti(10nm)/Pt(250nm) electrodes were then deposited using electron beam

evaporation.

Figures 4.2 (a) and (b) show the layout of a prototype gas sensor utilizing

NST pads as sensing elements consisting of a 5 by 5 array of square NST

pads, each 20 μm wide. In principle, each of these pads could be

individually addressed. For simplicity we have metallized only a few,

selected pads of the 5 by 5 array with 10 μm wide metal lines. Moreover,

further scaling down of the device is feasible with currently available

microelectronics process tools.

96

Figure 4.2. Optical micrographs of integrated NST arrays as sensing

elements in prototype gas sensors with 2 different configurations; (a) single pad metallized and (b) 3 pads metallized in series.

Although the NST pads are porous, the evaporated Ti/Pt electrodes have

good step coverage on the masking oxide, are continuous on the NST

elements and give satisfactory electrical contacts. Figures 4.3 (a) to (d) are

SEM micrographs of Ti/Pt electrodes on NST pads deposited before and

after NST annealing. Although Ti/Pt metallization has good step coverage

on the sidewalls of the PECVD masking oxide, as shown in Figure 4.3,

97

shrinkage of NST during annealing results in discontinuity of the

metallization lines at the corner of the masking oxide face with the NST

surface. However, this problem was not observed when Ti/Pt metallization

was deposited after NST pads have been annealed.

Figure 4.3. Continuity of Ti/Pt metallization layers deposited (a and b) before

and (c and d) after NST annealing.

98

The surface of Ti/Pt electrode layer is rough due to the underlying porous

NST morphology but is nevertheless continuous over the NST pad — see

Figures 4.4 (a) to (c). Electrode continuity is critical as it reduces electrical

contact resistance.

Grain size of Pt crystals on the NST layer is significantly larger than

those deposited on the P-SiO2 mask oxide. The average grain size of Pt

crystals on the NST and P-SiO2 mask oxide are ~35 nm and ~20 nm,

respectively. This smaller Pt grain size on NST pads could be the result of a

difference in surface density of Pt nuclei over the different materials during

evaporation. On P-SiO2 mask oxide, a larger nuclei surface density would be

expected. This is because, on the NST, nuclei can only be formed on the

surface of the pore walls. Because of the smaller nuclei density on the NST,

Pt crystals are able grow to over a wider distance before coalescing with an

adjacent grain, resulting in larger Pt grains on NST pads.

Ti/Pt electrodes deposited on the NST sensing elements do not

penetrate through the entire thickness of the NST layer as revealed by

cross-sectional SEM and TEM analysis as shown in Figures 4.5 (a) and (b).

The procedure for TEM sample preparation has been described in detail in

Chapter 2. In brief, TEM samples were micromachined using a focus ion

beam system using a Ga ion beam current of 100 pA. An FEI Sphera T20

machine operating at 200 kV was used for Transmission Electron

99

Microscopy analysis. Superficial penetration of metallization layers into the

NST pads is expected because electron beam deposition is a line of sight

process. Hence the metal layers deposited could only penetrate as deep as

the pores on the top surface of the NST. Hence the volume of NST under

the electrode is porous and remains unfilled. Similar observations of partial

penetration of metals evaporated on porous Si have been reported. In

contrast, complete infiltration of metal contacts has been obtained using

electroless deposition — this is the subject of a later chapter.

Figure 4.4. SEM micrographs of showing continuity of Ti/Pt electrodes

evaporated on single NST pads. Regions labeled in (a) represents location from which micrographs (b) to (d) were obtained.

100

Figure 4.5. Superficial penetration of Ti/Pt metallization into NST as

revealed by cross-sectional (a) SEM and (b) TEM analysis.

4.4 Fabrication of nanostructured TiO2 on Kapton®

Low temperature synthesis of NST was demonstrated by fabricating gas

sensing devices on Kapton® 300 HN foils (Kapton® foils were gifts from

DuPont). There are a few reasons why Kapton® was chosen. First, Kapton®

is suitable for use as a gas sensing substrate. Kapton® films are able to

101

withstand relatively high temperatures of up to 400 °C for at least 2 hrs. In

addition, Kapton® films are relatively inert to chemicals used during

processing, have low moisture absorption and high dielectric strength.

Secondly, Kapton® foils are widely used in industry and hence suitable for

large scale implementation.

Figure 4.6. Integration of NST on Kapton®: (a) Arrays of Ti window after SiO2 etch; (b) A complete prototype gas sensing device.

The processing steps described for integration of NST into devices on Si

are directly applicable to Kapton® substrates. However, as Kapton® films

are flexible, one difficulty encountered is the bending of substrates after

evaporation when Ti films are deposited. This issue however, could be

reduced by using sputtering where stresses in films are more readily

controlled. Stress in evaporated Ti is tensile while SiO2 films deposited by

102

plasma enhanced chemical vapor deposition are under compressive stress.

It was found that for 500 nm Ti films deposited on Kapton® 300 HN foils, an

overlying 100 nm thick SiO2 layer would produce a flat specimen. Figure 4.6

(a) shows an array of 20 μm square pads after etching of the SiO2 layer. As

expected, no delamination of Ti from the underlying Kapton® substrate was

observed. A complete gas sensing device with evaporated metallization is

shown in Figure 4.6 (b). A tip of a pen is included for scale.


Sensing experiments were carried out in a vacuum chamber

equipped with microprobe contacts for current-voltage (I-V) measurements.

The gas sensing behavior of NST to oxygen and hydrogen gases were

studied. The partial pressure of the gas to be detected in the test chamber

was controlled by means of a pulsed needle valve and confirmed using an

ion gauge. A halogen lamp heater was used to raise the sensor assembly to

the desired temperature, which was measured using a K-type thermocouple.

Prior to sensing experiments, the device was exposed to UV for about 5 min

an elevated temperature in house vacuum and then flushed with nitrogen

gas for about 10 min. This cleaning cycle is done for about 45 min every

time a new device is put under test or after long periods of exposure to air.

103

4.5.1 Oxygen sensing

Figures 4.7 (a) and (b) show I-V characteristics of an NST pad in vacuum

and under oxygen exposure at various temperatures and O2 partial

pressures ranging from 0.3 to 0.65 mTorr. The linear I-V characteristics

obtained indicate that contacts between the Ti/Pt electrodes and the NST

pads were Ohmic at the temperatures utilized. The conductance of the NST

pads is very sensitive to the presence of oxygen and changes monotonically

with oxygen partial pressure. In the presence of O2, conductance of the pad

is approximately an order of magnitude less than that in vacuum. This high

sensitivity to O2 is also illustrated in Figure 4.7 (b) where O2 pressure

variations in the sub-mTorr range were easily distinguishable. Assuming O2

to be entrained in an unreactive gas such as nitrogen, these mTorr pressure

variations at 523 K correspond to ppm detectivity by an NST pad. By using

narrower and longer NST pads the detectivity could be enhanced to the ppb

level. Reducing the 20 μm square pads used in this prototype device is

feasible with current process tools.

104

Figure 4.7. I-V characteristics of an NST pad in vacuum and under oxygen exposure at various temperatures and O2 partial pressures ranging from 0.3

to 0.65 mTorr.

105

Figure 4.8. Sensing characteristics an NST pad element: (a) Response to oxygen cycles; (b) sensitivity to oxygen at various temperatures.

106

The response time and sensitivity of the sensor for oxygen detection was

obtained by measuring resistance changes when oxygen is introduced

cyclically in the chamber — results of these studies are shown in Figures 4.8

(a) and (b). Upon exposure to oxygen gas at 0.8 mTorr at 473 K the

resistance increased from 10 to 90% of its maximum value in 48 s. This

value, which will be referred to as the response time is comparable to those

of other oxygen sensing devices utilizing porous undoped-titania as a

sensing material. Sharma et. al. have fabricated porous undoped titania

films using various techniques. The resulting titania films have pore size

distribution ranging from ~0.5 to 1.5 μm and 0.75 to 2.5 μm35,36. However,

the response times they reported were obtained at higher temperatures. To

compare these response times with the values obtained in this study we

assumed that response times depend on temperature in an Arrhenius

fashion. For the titania films with ~0.5 to 1.5 μm and 0.75 to 2.5 μm pores,

the response times at 473 K were calculated to be 49 and 20 s, respectively

– see Figure 4.9. The sensitivity of the sensor is defined as the ratio of the

maximum resistance in the presence of oxygen to the resistance in the

absence of oxygen. The best sensitivity of the NST-based sensor, which

was measured to be 60 at 473 K, is superior to those of undoped-titania

based sensors reported in the literature. Sharma et. al. and Gao et. al.

reported sensitivity values of 1.43 and ~30 at their higher operating

temperatures35-37. The sensitivities expected for those systems at 473 K

107

would be much lower. Hence the oxygen sensitivity of sensors fabricated in

this study is at least twice that reported by Sharma et. al.. Another

advantage of the present approach to integrating NST into microsystems

devices is the low level of drift in the electrical properties of the NST sensor

with time, despite the low processing temperatures involved. This could be

due to the high crystallinity of the NST pads formed as indicated by electron

and X-ray diffraction data.

Figure 4.9. Comparison of response time for oxygen gas at 273 K of NST-based sensor to those of other devices based on porous undoped titania.

108

Undoped TiO2 shows n-type semiconducting behavior due to the

presence of shallow intrinsic donors22,38. Below 1273 K, the dominant donor

species are singly and doubly ionized oxygen vacancies, with associated

defect levels located just below the conduction band22,38. For TiO2-based

sensors, the resistance change upon O2 exposure is due to the interplay

between the ionic chemisorption of oxygen at the TiO2 surface and the

density of carriers in the material’s conduction band. Upon adsorption, an O2

molecule migrates to an oxygen vacancy site39, where it interacts with

surface vacancies eventually forming chemisorbed species such as O2- and

O- (depending on temperature). The process involves electron transfer

across the titania surface depleting the electron density of the bulk TiO2. The

net reduction in conduction is due to annihilation of the donor states. At

elevated temperatures the oxygen-induced vacancy annihilation process

could involve subsurface and even bulk donors35,36. Hence, for enhanced

sensing performance, materials with structures that allow permeation of

oxygen through a porous sensing element with a large surface area and thin

channels are desirable. For conventional TiO2 dense film-type sensors, the

compact nature of the particles means that only a thin layer of the film close

to the surface can effectively be used for sensing. NST on the other hand

has thin pore wall thickness of 25 – 75 nm, which means that a significantly

greater TiO2 surface area can be used for sensing as the open porous

109

structure allows enhanced oxygen gas permeation. As compared with

previously reported porous TiO2 films35, the pore size of NST used in this

study, ranging from 50 nm to 200 nm, is ~10 times lower and hence is

expected to have reduced oxygen gas permeation. However, the response

times at 523 K are similar. In addition, the NST-based O2 sensor exhibits

superior sensitivity. These observations suggest that enhanced oxygen

sensing performance of NST is due to a higher level of surface reactivity,

possibly due to presence of a greater number of surface defects.

Another possible mechanism that enhances the sensitivity of the sensors

is the sponge-like structure of NST. As discussed, for chemical sensing to

take place the presence of chemical species on the NST surface must be

transduced into an electronic signal. In the case of NST sensors, the

sponge-like morphology significantly facilitates electron percolation through

the TiO2 pad resulting in an increase in current. In the case of sensors

employing pads consisting of agglomerations of TiO2 nanoparticles, electron

percolation may not occur as well. This explanation is plausible as

analogous observations had been made in the case of solar cells using TiO2

electrodes40. In that case, it was found that current harvested from solar

cells was increased when TiO2 electrodes were treated with TiCl4. This

increase was due to the formation of interparticle neck growth that facilitates

electron conduction.

110

4.5.2 Hydrogen sensing

NST-based sensor was also used to detect hydrogen gas. The I-V

characteristics obtained at room temperature in the absence and presence

of hydrogen gas are shown in Figure 4.10 (a). The concentration of

hydrogen used was partial pressure 0.8 mTorr which is equivalent to ~1

ppm. In the presence of hydrogen, current is increased since hydrogen is a

reducing gas, as expected. The current-time response of the device to

hydrogen at a partial pressure of 0.8 mTorr at 300 °C is shown in Figure

4.10 (b) which indicates poisoning of the NST element. Upon introduction of

hydrogen, the resistance measured decreased to a ‘minimum’ value and,

when the hydrogen valve is switched off, increased to a ‘maximum’ value

and saturated. Upon further introduction of hydrogen, the resistance drops

to a ‘minimum’ which is lower than the ‘minimum’ value of the preceding

cycle. Similarly, when hydrogen gas flow was switched off, resistance

increased and saturated again at a ‘maximum’ value which is lower than the

‘maximum’ value of the preceding cycle.

Response time and sensitivity of the sensor towards hydrogen were

calculated from Figure 4.10 (b). Varghese et. al.41 and Patel et. al.42 have

recently reported hydrogen sensing studies using undoped titania. In these

studies, response time was defined as the time required for the sensor to

111

reach 90% of the maximum resistance and sensitivity is defined by the

formula:

gs

gs

RRR

S−

= 0 …….Eqn. 4.1

where and are the resistances in the absence and presence of

hydrogen gas, respectively. These definitions are different from that those

discussed in the section on oxygen sensing. However, for the sake of

making comparisons to the work of Varghese et. al.

0R gsR

41 and Patel et. al.42

(instead of sensitivity, this parameter is called % resistance change in the

work of Patel et. al.42) the definitions proposed by these researchers are

used in this section. From results shown in Figure 4.10 (b), it was found that

the response time and sensitivity of the NST based sensor to hydrogen are

78 s and 3.7%, respectively.

Using titania nanotubes as sensing elements, Varghese et. al.41

studied sensing down to 100 ppm hydrogen ambients. However, it was

found that sensitivity has a logarithmic relation with hydrogen concentration

and it was calculated from their data that the sensitivity for 1 ppm hydrogen

at 290 °C is 2 which is comparable to the present work. The corresponding

value reported by Patel et. al.42 at 300°C is 1. Varghese et. al.41 reported a

response time for ~350 s for 100 ppm hydrogen ambient – response time for

112

1 ppm hydrogen ambient was not reported. Patel et. al.42, did not report of

response time values in their work.

The mechanism of transduction of hydrogen adsorption on NST

surface to a change in electrical resistance is still an open question. There

are three possible mechanisms all of which can occur, independently of

each other. First, hydrogen molecules adsorb to the NST surface and

dissociate to atomic hydrogen which are subsequently chemisorbed. The

hydrogen atoms then donate an electron to the conduction band of the

titania causing the observed reduction the resistance. The protons can be

trapped as surface states or diffuse into the NST lattice. Second, an effect

known as spill over may occur. In spill over, molecular hydrogen is adsorbed

on the surface of the Pt electrodes and dissociates into atomic hydrogen.

These hydrogen atoms diffuse onto the NST surface and contribute to the

electrical conduction as mentioned above. Third, hydrogen molecules are

adsorbed onto the NST surface and dissociates into atomic hydrogen.

These atomic hydrogen then reacts with oxygen species such as O- or O-2

ions already adsorbed on the surface to form H2O. This reaction removes

electron traps from the surface causing a drop in the resistance observed.

Hydrogen sensing measurements were done in oxygen-free

environments, hence the increase in conductivity cannot be due to removal

of oxygen O- or O-2 ions by hydrogen. Also, the extremely fast response of

the NST sensor to changes in hydrogen concentration suggests that surface

113

effects are dominant. Upon dissociation of molecular hydrogen, hydrogen

atoms are chemisorbed on the titania surface. Chemisorbed atomic

hydrogen act as a surface state and charge transfer occurs from hydrogen

to the conduction band of titania resulting in an increase in conductivity. The

poisoning of titania by hydrogen, as indicated by the continuous drops in

‘minimum’ and ‘maximum’ values of different cycles could be explained by

trapping of hydrogen in bulk states. Chemisorbed hydrogen atoms may

diffuse into the bulk titania and form donor type OH-defects that join two

oxygen atoms43. These defects are not removed when the hydrogen gas

flow is switched off which results in the lowering of the ‘minimum’ resistance

from that in the preceding cycle.

114

Figure 4.10. Hydrogen sensing: (a) I-V characteristics of an NST pad in vacuum and under hydrogen exposure (8x10-4 Torr) at room temperature and their differential; (b) Response to hydrogen cycles at 300 °C (8x10-4

Torr).

115

4.6 Conclusions

In conclusion, we have described a method which is compatible with

current microelectronics manufacturing practices for forming an array of

crack-free patterned and integrated NST (NST) pads for sensor-on-a-chip

and other micro- and nanodevice applications. Using patterned Ti thin films

with lateral dimensions below a certain threshold, crack formation in NS-

titania layers can be eliminated. NST has sponge-like morphology with pore

diameter and wall thickness in the ranges of ~25 – 200 nm and 25 - 75 nm,

respectively. Using these integrated NST pad arrays as resistive

components, we successfully fabricated a prototype oxygen sensor capable

of detecting oxygen and differences in oxygen concentration at ppm levels.

The sensor operates at lower temperatures, has a fast response time and

superior sensitivity relative to oxygen sensors based on undoped-titania.

116

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119

Chapter 5: Fabrication of patterned micrometer scale

interpenetrating Au–TiO2 nanocomposites

5.1 Introduction

Nanocomposites are, by definition, materials in which at least one of the

phases has constituents less than 100 nm in size1. Metal-ceramic

nanocomposites thin films presently attract tremendous attention due to their

fascinating properties and potential applications in electronics,

optoelectronics, magnetics and catalysis2. For MOS-type storage

applications, metal-ceramic nanocomposites consisting of either Ag, Au, Pt

or Ni nanoclusters in ultra-thin SiO2 or HfO2 layers hold advantages over

those employing semiconductor nanoclusters 3-5. In optoelectronics, metal-

ceramic nanocomposites with large third order optical non-linearities are

promising candidates for use in all-optical switches and optical

computation6-10. Also, nanocomposites of Co and FePt particles dispersed

in a non-magnetic ceramic phase such as SiO2, Al2O3, B2O3 and Si3N4 are

being intensively investigated for high density magnetic storage

applications11-13. In catalysis, nanocomposites of catalytically active metal

nanoparticles of either Ni or Pd, for example, in either TiO2, ZrO2 or MgO

phases are being studied for applications such as decomposition of

120

methanol to produce hydrogen and dry reforming of methane to produce

syngas14-16.

Various methods have been proposed for fabricating metal-ceramic

nanocomposites. These include evaporation10,14, sputtering9,11,12, ion

implantation7,8, deposition-precipitation15,17 and sol-gel processing6,13,16,18-

21. The sol-gel process has received particular attention as it renders

molecular-level control, allows the incorporation of various metal dopants

into ceramic matrices and is relatively low cost. Although conventional sol-

gel processing produces a continuous gel film, in applications where

patterned features are required, additional process steps such as reactive

ion etching, laser trimming, embossing or through-mask UV irradiation of

chelate modified gels and subsequent dissolution are used22-28.

Alternatively, self-assembled monolayers have been used as masks to

direct formation of gel layers from precursor solutions29-31. However, all the

aforementioned methods produce dispersed, non-percolating metal

nanoclusters embedded in a ceramic phase.

121

5.2 Interpenetrating network nanocomposites

An interpenetrating network composite is one in which each constituent

phase is continuous and interpenetrating throughout the microstructure32.

Previous reports of interpenetrating network nanocomposites are largely

limited to organic-organic33 and organic-inorganic material systems34. Prior

work in metal-ceramic material systems involve micrometer-scale phases

and hence, by definition, may not be classified as nanocomposites35. Au-

TiO2 interpenetrating network nanocomposites are expected to have the

functionalities of its constituents such as the wear resistance of TiO2

together with the high electrical conductivity of Au32. Thin films of Au and its

alloy have been widely used as contact material in microswitch devices

because of their low-resistivity and oxidation resistance36. However, they

suffer from wear and stiction which shorten device lifetimes37. In contrast,

titania films formed using sol-gel processes have excellent wear resistant

properties19. Presently, there is no prior work on the fabrication of patterned

micrometer scale metal-TiO2 interpenetrating network nanocomposites. In

this chapter, we demonstrate a technique to form patterned features of an

interpenetrating network nanocomposite of Au and nanostructured TiO2

(NST). NST features were first fabricated by reacting pre-patterned Ti with

aqueous hydrogen peroxide, as described previously. Using electroless

122

deposition, Au was then infiltrated into the NST features. Although Au was

used in this study, the process can be generalized to other metals that can

be deposited using electroless deposition.

5.3 Fabrication of patterned interpenetrating Au–TiO2

network nanocomposites

Fabrication of interpenetrating Au-NST network nanocomposites involves

two main process steps. Firstly, porous NST patterns are formed and,

secondly, NST patterns are infiltrated with Au using electroless deposition. A

schematic of the process is shown in Figure 5.1. NST pad arrays with

porous sponge-like morphology were formed by reacting pre-patterned

arrays of exposed Ti surfaces with aqueous hydrogen peroxide (aq. H2O2)

solution followed by low temperature annealing. The process steps for

fabricating the NST has been discussed in detail in Chapter 2, here only a

brief description is given. To prepare Ti thin film samples, 2.5 cm square

pieces of either N-type Si(100) or glass pieces were used as substrates. Si

pieces were first thermally oxidized at 1100 °C to grow 1 μm thick SiO2

layer. For glass substrates, a 1 μm thick SiO2 layer was deposited by

plasma-enhanced chemical vapor deposition (PECVD). Substrates first

cleaned for 5 min each in acetone, 2-propanol and de-ionized (DI) water

123

(18.9 MΩ) and blown dry with nitrogen before Ti was deposited. Ti film was

deposited on substrates using electron beam evaporation. Photo resist (PR)

was then spun on the silicon dioxide layer and patterned. The pattern on the

PR layer was transferred to the silicon dioxide layer by etching with CHF3

gas. The exposed patterns of Ti surfaces were oxidized by aging in 10% (by

volume) aq. H2O2 at 80 °C. This is followed by annealing at 300 °C for a few

hours. Aging and annealing steps were done in ambient air.

Au was infiltrated into pores of the NST using selective electroless

deposition. Commercially available electroless gold plating solution, based

on alkaline gold cyanide complex with disodium ethylenediaminetetraacetate

(Na2EDTA) as chelating agent, was purchased in a ready to use condition.

Chelating agents are species that form multiple bonds to, and hence

sequesters, a metal ion. In this study, Na2EDTA is used to trap metal ions

that would otherwise initiate secondary reactions that reduce the lifetime of

the Au plating bath. Substrates were coated with Au by immersion in a

heated bath of the plating solution. Prior to immersion, temperature of the

solution was maintained and stabilized at 60 ± 2 °C for about 20 min.

Temperature of the sample was maintained at the set temperature using

automated electronic feedback control. Substrates were immersed using

Teflon holders to eliminate contamination of the electroless plating bath. For

each plating run, fresh plating solution was used to ensure similar Au

124

deposition rate was obtained for all runs. After Au deposition, samples were

soaked and rinsed thoroughly in deionized water and subsequently blown

dry with nitrogen.

Figure 5.1. Schematic process flow for forming patterned Au-NST nanocomposite. (T-SiO2 and P-SiO2 denote thermally grown SiO2 and

SiO2 deposited by PECVD, respectively.)

125

Various techniques were used to characterize the NST and Au-NST

nanocomposites fabricated. X-ray diffraction using CuKα radiation (1.5406Ǻ)

in Bragg-Brentano configuration was used to determine crystal structure.

Structural characterization was done using an FEI dual beam focus ion

beam system equipped with Ga ion and electron columns for

micromachining and imaging, respectively. To micromachine samples, a Ga

ion current of 100 pA was used. Surface chemical species was determined

using a Kratos Axis Ultra X-ray Photoelectron Spectroscopy (XPS) system.

High resolution and survey scans of various surfaces were obtained. XPS

scans were obtained using monochromated Al Kα source (1486.6 eV) and

20 eV pass energy with steps of 0.05 eV at a base pressure of 7.5 x 10-9

Torr. XPS spectra collected were fitted to line shapes constructed from a

linear combination of Gaussian and Lorentzian profiles using commercial

software (CasaXPS). Atomic force microscopy (AFM) was used to

investigate the surface morphology of parent Ti, NST and Au-NST

nanocomposite. A Digital Instruments (Veeco) Dimension D3000

microscope operating in tapping mode in air was used in the AFM studies.

126


As discussed previously, reaction of metallic Ti with hydrogen peroxide

had been investigated previously and was shown to result in formation of a

hydrated titania gel layer38,39. Because of the high crack-density in TiO2

layers formed using this technique, they are generally not suitable for device

applications40,41. We have recently developed a technique that allows

integration of NST into microsystems and eliminates crack formation. This

method involves pre-patterning Ti surfaces prior to reaction with hydrogen

peroxide41. After aging in hydrogen peroxide and subsequent annealing,

NST patterns have a nanoporous sponge-like morphology with

interconnecting pore walls; see Figures 5.2 (a) to (c). The pore walls have

thickness ranging from about 75 – 125 nm while pore diameters range from

50 nm – 200 nm. Unannealed NST patterns are hydrated titania gel layer

consisting of peroxo compounds38. In some regions, walls of pores are long

and narrow enough to be described as wires. The porous structure of NST

could be due to morphology of the intermediate gel layer formed during

aging in aq. H2O2. Similar porous microstructures were observed when

unpatterned bulk Ti foil was reacted with aq. H2O2 41. Figure 5.2 (c) is a

cross-sectional SEM image of NST obtained after focus ion beam milling. To

reduce material redepositing on the cross-sectioned surface, ion beam

127

milling was done at 100 pA. The cross-sectional image reveals that the 350

nm thick Ti layer had been fully oxidized and the porous structure of the

NST layer formed extends to the underlying SiO2 surface. NST patterns

formed have excellent adhesion to the underlying SiO2.

Representative SEM images of Au-NST nanocomposites formed after

electroless deposition of Au on NST patterns are shown in Figures 5.3 (a) to

(c). Electroless plating is the result of a chemical reaction in which a metal

layer is formed by reduction of a metallic ion by a reducing agent in an

aqueous solution. Protrusions 25 to 50 nm high decorated the surface of the

Au-NST nanocomposites. These protrusions were formed when Au plates

on titania wall protruding from the NST layer. After infiltration, pores of the

NST patterns were completely filled with Au. Figure 5.3 (c) shows that Au

has deposited in the pores throughout the NST feature and not just on the

surface. In addition, no significant void formation was detected in the Au-

NST nanocomposites. This suggests that during the electroless plating

process, there is no significant difference in the rate of Au deposition in the

NST layer than that on the top surface of the NST which could be due to the

relatively large pore size of the NST after annealing.

128

Figure 5.2. SEM images of NST: (a-c) before Au infiltration, (a) top view of a 20 μm pad, (b) higher magnification image (c) cross-sectional images

after micromachining (tilt 30 degree).

129

Figure 5.3. SEM images of Au-NST interpenetrating network nanocomposite: (a) top view of a 20 μm pad, (b) higher magnification image

(c) cross-sectional images after micromachining (tilt 52 degree).

130

The phase evolution of NST after has been discussed in Chapter 241.

Here we provide only a brief description. We found that after aging in aq.

H2O2, NST consists, largely, of amorphous TiO2 and nanocrystals of

anatase and that the amorphous TiO2 phase transforms to anatase upon

annealing as discussed in previously. XRD studies of patterned NST pad

arrays indicate that Au was deposited on NST pads only and not on the SiO2

mask. Such selective deposition was observed on amorphous and

crystalline NST; corresponding to the as-aged and annealed conditions,

respectively. Figure 5.4 shows XRD spectra of amorphous NST after plating,

demonstrates the selective deposition of Au after various immersion times in

the electroless plating bath. A comparison of XRD spectra of the sample

after Au plating indicates the presence of three peaks at 2θ values of 44.18,

64.46 and 81.76 degree on the NST pad array. These peaks, however, are

not present in spectra collected from the SiO2 mask region. These

observations strongly indicate that Au has deposited only in pores of the

NST pads and not on the surface of SiO2 mask. The three peaks observed

can be assigned to the 200, 220 and 222 of Au. A prior study of

electroless Au deposition on Ti films, which has an amorphous native oxide,

has reported the presence of these three Au reflections42. It is noted that the

patterned NST sample has also been partially oxidized leaving a layer of

residual Ti below the NST layer. Ti films on the patterned and blanket

samples used for XRD studies were deposited in different evaporation

131

machines, which resulted in different textures as reflected in the as-

deposited XRD spectra. Prior studies indicate that Ti surface texture has

little effect on the morphology of NST formed41.

Figure 5.4. XRD of samples illustrating selective Au deposition on unannealed NST(i) As deposited Ti; (ii) Plate 40 min, SiO2 mask; (iii) Plate 30 min, NST array; (iv) Plate 50 min, NST array. (Peaks for Ti and Au are

denoted by and ∆, respectively).

132

Results of XPS studies are shown in Figures 5.5 and 5.6. All spectra

were referenced to C 1s peak at 285.0 eV. Figure 5.5 (a) is a high resolution

scan of the Ti 2p peaks of unpatterned NST film formed on PECVD SiO2 on

Si substrate after annealing at 8 hr for 300 °C. Similar results were obtained

from NST formed on a PECVD SiO2 coated glass substrate. Assuming a

Tougaard, background, fitting to the raw spectra was done using Gaussian-

Lorentzian components using a commercial XPS analysis software

(CasaXPS). From the analysis, binding energies of Ti 2p3/2 and Ti 2p1/2 were

found to be 458.9 and 464.8 eV, respectively. These values are close to

those reported in the literature for TiO2 of 458.9 and 464.6 eV for Ti 2p3/2

and Ti 2p1/2, respectively, and confirm the formation of TiO2 after

annealing43,44.

Figure 5.5 (b) shows XPS spectra for binding energies from 55 to 110 eV

of unpatterned NST film before and after 5 min Au plating. Before plating

only Si 2p signal with a peak at 101.3 eV was detected. After plating,

however, three additional peaks were detected. Assignment of these peaks

was similarly done using CasaXPS. Two peaks at binding energies of 84.60

and 88.16 eV are assigned to Au 4f7/2 and Au 4f5/2. These experimental

values of Au 4f7/2 and Au 4f3/2 obtained are consistently higher, by 0.38 and

0.22 eV respectively, but close to the corresponding values reported in the

literature; which range from 83.70 to 84.25 eV for Au 4f7/2 and 87.71 to 87.94

133

eV for Au 4f5/2 of elemental Au 45-47. Hence, it is concluded that

electrodeposited Au is in the elemental form, in agreement with XRD results.

The peak at 63.13 eV in Figure 5 (b) is assigned to Na 2s; another peak,

which is assigned to Na 1s, at 1071.58 eV was also observed after plating.

A prior study has reported entrapment of sodium in a metal layer deposited

from a plating bath that similarly used Na2EDTA as a chelating agent48. It is

known that Na when dissolved in Au, may exist as Au(Na) alloy and/or

AunNa (where n = 5, 2, 1) intermetallic compounds49. Solomun had

determined Na 1s binding energies of adsorbed sodium atoms, in the

presence of coadsorbed iodine atoms, on Au(100) surfaces. In the absence

of iodine and when coverage of Na atoms on the Au surface is 0.09 of a

monolayer, the Na 1s binding energy was determined to be about 1071.05

eV, which is close to that obtained in the present study50 and suggests that

a Au(Na) alloy has been formed. Unfortunately there are no reports in the

literature of Au 4f7/2 and Au 4f5/2 binding energies from AunNa intermetallic

compounds for comparison. However, the presence of AunNa compounds

after plating, if any, is not extensive because no XRD reflections detected

can be assigned to these compounds. From the preceding discussion it may

be concluded that Au has been deposited into the pores of NST pads and

Na, at a dilute concentration, was incorporated into the Au layer during

plating.

134

Figure 5.5. XPS results of: (a) Ti2p high resolution line scans of NST before Au plating; (b) line scans before and after Au plating for binding

energies from 55 to 110 eV.

135

Results of area-mode XPS studies confirm that Au deposits selectively

on annealed NST. Figures 5.6 (a) and (b) show, respectively, Si 2p and Au

4f signals of one region of an NST pad array after electroless Au plating.

The results of the area-mode XPS are presented in gray-scale in which

areas with higher concentration of a chemical species appear brighter. In

addition, it is noted that XPS is a technique sensitive to chemical species on

the first few nanometers from the surface only51. From Figure 5.6 (a), the

strong Si 2p signal implies that very little, if any, Au has deposited on the

SiO2 mask oxide surface. This is confirmed in Figure 5.6 (b) where Au 4f

signal is observed only in areas that correspond to those of NST pads which

strongly indicate that Au has been deposited only on the NST pads and not

on the SiO2 mask. Hence, XPS results are in agreement with those of XRD

studies and indicate that Au has been deposited selectively only on NST

features.

136

Figure 5.6. XPS area-mode scans, after Au plating, of (a) Si 2p and (b) Au 4f signals.

137

Surface morphologies of the parent Ti thin film, NST and Au-NST

nanocomposite formed were studied using atomic force microscopy (AFM).

Figures 5.7 (a) to (c) show typical images of these surfaces. The surface of

the parent Ti film is rough, consisting of distinctly faceted Ti nanocrystals

separated by gaps. Ti crystals are platelet-like with thickness and longest

width of 40 ± 10 nm and 124 ± 31nm, respectively. (All stated dimensions

are an average and one standard deviation of 30 measurements.) In some

regions of the Ti surface, gaps are large enough to be described as troughs

with average diameter of about 101 ± 37 nm. AFM images of surfaces of

NST formed after oxidation and subsequent annealing steps show a

distinctly different morphology from that of the as deposited Ti. The NST

surface is rough with ridge-like protrusions that surround holes. The ridge-

like protrusions correspond to titania walls and consists of spherical

protrusions about 93 ± 32 nm in diameter. Average diameter of holes was

measured to be about 126 ± 33 nm. After Au plating however, the surface of

the Au-NST nanocomposite formed consists of globular grains with average

diameter of about 52 ± 12 nm. This results in a smoother surface, as Au was

deposited in the troughs and on the walls of the NST layer

138

Figure 5.7. AFM images of (a) as-deposited Ti film; (b) as-formed NST and (c) Au-NST nanocomposite formed after Au plating. (Images are 1 x 1

μm squares)

139

5.5 Integration of Au–TiO2 nanocomposites as contacts in

devices

In this section we demonstrate the fabrication of electrical contacts made

from interpenetrating Au-NST network nanocomposites. The device, in

which the contacts were integrated, is a microswitch that operates by the

closure and opening of a gap between two electrodes which are Au-NST

contacts. A schematic illustrating the operation of a switch is shown in

Figure 5.8. The gap between the electrodes is closed when a cantilever

deflects down and makes contact with the two electrodes. In this state the

circuit consisting of the electrodes is closed. When the cantilever is not

deflected, the gap between the electrode is open. Hence by deflecting the

cantilever up and down using a drive signal, the circuit consisting of the

electrodes can be switched on and off. At present the contacts in

microswitch devices are made from Au or its alloys. However, these

materials present reliability issues of wear and stiction36,37. In this section,

we propose and demonstrate a process that integrates Au-NST network

nanocomposites as contact materials.

The process steps for fabricating the electrical contacts are similar to

those describe in the Figure 5.1. Here only the salient steps are described in

detail. Because of shrinkage, thickness of NST features after annealing is

less than that of the masking oxide PECVD. To make the contacts a Ti layer

140

is first sputtered and subsequently patterned by etching in fluorine

chemistry. Of course, the Ti pattern can be formed using the lift-off

technique just as well. Au patterns were then evaporated on the Ti layer

using a lift-off technique to protect area of where NST is not desired. The

device is then soaked in hydrogen peroxide solution to form NST and

subsequently annealed at 300 °C. As mentioned earlier, incorporation of

impurities such Na may increase the resistance of the Au-NST composite.

To alleviate this problem, a layer of gold is evaporated on the Au-NST

composite. Figure 5.9 shows the integrated Au-NST composite contacts.

Figure 5.8. Schematic of a micro-switch device structure and operation: (a) Top view of device; Device in (b) on and (c) off states. (Micro-switch

fabrication work, except Au-NST integration, was carried out by Chang Song Ding)

141

Figure 5.9. Integration of Au-NST nanocomposite contacts in devices: NST pad (a) before Au-infiltration and (c) after Au-infiltration followed by

flash Au evaporation. (Magnification: 200X)

142

5.6 Conclusions

In this chapter we have demonstrated a facile technique to form

integrated and patterned micrometer-scale features of interpenetrating Au-

nanostructured TiO2 (NST) network nanocomposites. First, NST pads were

formed by aging Ti surfaces in aqueous hydrogen peroxide (aq. H2O2)

solution. As aged NST is largely amorphous but transforms to single phase

anatase upon annealing at 300 °C. Second, pores of the NST are then filled

with Au using electroless deposition. Cross-sectional SEM images show that

complete infiltration of NST pores was achieved with little void formation. X-

ray diffraction and x-ray photoelectron spectroscopy studies indicate that Au

was deposited selectively on NST pads and not on surfaces of the SiO2

mask. It must be noted that although Au was used in this study, it is

postulated other metals that can be deposited using electroless deposition

would work just as well. Hence, this technique represents a general

technique for integrating interpenetrating metal-NST network

nanocomposites features in microsystems.

143

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147

Chapter 6: Attachment of mouse fibroblasts on

nanostructured TiO2

6.1 Introduction

At present there are tremendous efforts towards using microelectronics

devices and microsystem (MEMS) devices for biological applications1-3. The

cellular milieu/electronic material interface may need to be engineered and

optimized to suit these applications. Various approaches have been

investigated for engineering the cellular milieu/electronic material interface.

A few research groups have demonstrated that grafting of organic moieties

such as peptides, proteins and sugars on surfaces allows modulation of cell

attachment on surfaces4-9. Use of plasma polymerized tetraglyme coating

(similar to polyethylene glycol or PEG) on the other hand results in

biocompatible surfaces with non-fouling properties10. Areas on a device

where cells are not desired can be arbitrarily defined by patterning and

removing the tetraglyme layer. In addition to modifying the chemical

moieties on the surface, other techniques include engineering the surface

topography. In this way it was found that transformed as well as primary glial

cells preferentially attached to silicon pillar arrays rather than smooth silicon

surfaces11. Similarly, etching of Si results in a biocompatible porous Si that

148

sustains growth of calcium phosphates and viability of primary rat

hepatocytes12,13.

The excellent in-vivo biocompatibility of titania is well known and use of

bulk Ti implants for dental and bone prosthetic applications is well

documented14,15. Recent work indicated that nanophase titania, with grain

sizes from ~20-40 nm enhanced osteoblast cell functions in vitro16. Hence,

NST seems to be an attractive material for optimizing the cellular

milieu/electronic material interface. The objective of this chapter is to

investigate the attachment of mouse fibroblasts cells on blanket

(unpatterned) NST films as well as patterned NST features. As a

comparison, two other types of substrates were studied – Si substrates

coated with either plasma-deposited SiO2 (PECVD SiO2) or sputter-

deposited titania. At present, PECVD SiO2 is the most common passivation

layer for BioMEMS devices.

6.2 Attachment of cells on surfaces

Attachment of cells on various materials was investigated by performing

cell seeding experiments. The area of a sample covered with cells, after

duration of cell seeding was then determined using image analysis.

Morphology of cells was studied using optical microscopy, laser scanning

confocal microscopy and scanning electron microscopy.

149

Materials preparation

PECVD SiO2 and smooth titanium oxide films were coated on separate

electronic grade N-type silicon wafers (Mitsubishi Electronic Materials).

Plasma enhanced chemical vapor deposition were used to deposit 200 nm

thick PECVD SiO2. A coating of ~150 nm thick titanium oxide films was

deposited by reactive sputtering Ti in an Ar/O2 plasma. These wafers were

then cut into 1 cm square tabs using an automated saw. Tabs were then

soaked, with ultrasonic agitation, sequentially in acetone, isopropyl alcohol

and deionised water (18 MΩcm) for 5 min each. The tabs were then blown

dry with nitrogen and heated at 120 °C on a hotplate for 10 min to remove

moisture. Porous NST used were formed from 500 nm Ti films that was

aged for about 2.5 hr in hydrogen peroxide solution and subsequently

annealed at 800 °C for 8 hr. Prior to annealing the hydrated titania gel layer

was soaked in DI water for about 30 min.

Cell seeding

Mouse fibroblasts cells from a cell line were maintained in Dulbecco’s

modified Eagle’s medium (DMEM) supplemented with 10 % (v/v) fetal

bovine serum (FBS) and 1 % penicillin/streptomycin antibiotics in 5 % CO2

ambient at 37 °C. Cell cultures were passaged every 3 days. Cell seeding

experiments were carried out in 12 well plates with one tab per well. Cells

were dislodged by incubating in 5 ml of 0.05 % trypsin / 0.5 mM EDTA

150

solution. After incubation, trypsin was inhibited by adding 5 ml of serum

containing media. The cell suspension was concentrated by centrifuging at

1500 revolution per minutes (rpm) for 5 min to form a pellet of cells. The

cells were then resuspended in warm DMEM. Concentration of cells in the

suspension was measured using a hemacytometer.

Cells were plated at a density of ~105 cells per well. Each well contains

one tab. The cells were then incubated at 37°C in 5 % CO2 / 95 % air

environment. After the desired incubation time, cells were rinsed twice with

warm 1 X phosphate buffered saline (PBS) and fixed with 4%

paraformaldehyde in PBS. Tabs were then observed under optical

microscopy and micrographs taken. The percentage of area covered by

cells in a micrograph was calculated using Scion Image - a free software

from the National Institutes of Health. Micrographs taken of bare Si tabs as

well as PECVD SiO2 and TiO2 coated tabs have good brightness and

contrast. However, micrographs of NST-coated tabs were dark due to

scattering. For the latter, micrographs taken using scanning electron

microscopy were used to determine percentage of area covered. To

calculate the percentage of area covered by cells, at least five readings

were used.

For laser scanning confocal microscopy studies, cells were stained using

Texas Red-dye obtained from a commercial vendor. Preparation of cells for

scanning electron microscopy studies involves dehydrating cells using a

151

critical point dryer. Ethanol was used as the intermediate solvent. Prior to

drying, cells were soaked in phosphate buffered saline-ethanol solutions at

increasing concentration of ethanol to prevent shrinkage of cells. To reduce

charging during SEM, tabs were coated with ~5 nm Ti layer. Samples were

observed in SEM mode of the FEI dual beam focus ion beam system.


6.3.1 Seeding of fibroblast on various materials

Fibroblast cells exhibit enhanced attachment to NST surfaces after short

seeding times. Figure 6.1 shows the average area of various surfaces

covered by fibroblast cells after cell seeding. The area covered by cells for

NST surfaces is consistently greater than that of PECVD SiO2 for culture

times up to 18 hrs. After 24 hr culture time, the area coverage for all

surfaces is ~45 %. There is no significant difference between those for

sputtered and NST surfaces. To determine the level of significance in

differences between average values of area coverage, the t-probability value

in student’s t-test is used. It is generally accepted that for a comparison

between two data sets a t-probability value less than 0.05 indicates a

significant difference in the means. Only t-probability values for comparisons

between PECVD SiO2 and NST are given in Figure 6.1. The t-probability

values for cell seeding times of 3 to 18 hours are less than 0.05. In contrast,

152

for 24 hr cell seeding time, t-probability value is 0.9912. These values

indicated that there is significant difference in means of corresponding data

sets of PECVD SiO2 and NST at seeding times up to 18 hr.

Figure 6.1. Trend of average area covered by cell of various surfaces with seeding time. Error bars indicate one standard deviation. t-

probability values are for comparisons between area coverage on PECVD SiO2 and NST.

In addition to differences in area coverage, cells attached on PECVD and

NST surface exhibit different morphology. Figure 6.2 (a) and (b) show

confocal scanning microscopy images of fibroblasts after 6 hrs seeding time.

Cells seeded on PECVD SiO2 have spherical globular morphology with little

153

contact to the surface. In contrast cells seeded on NST exhibit a flat

morphology with cells spread out on the surface. It is now widely accepted

that cells with little preference to adhesion on a substrate adopt spherical a

configuration17,18. The flat spread out morphology of fibroblasts on titania

surfaces then indicates enhanced adhesion relative to PECVD SiO2

surfaces. After 24 hr seeding time, cells seeded on PECVD SiO2 surfaces

also exhibit the flat spread out morphology.

Figure 6.2. Confocal microscopy images of fibroblast seeded for 6hr on (a) PECVD SiO2 and (b) NST. (Courtesy of Blaine Butler)

Further evidence for adhesion is the formation of numerous processes

shown in Figure 6.3 (a) to (c). These processes are about 75 nm in diameter

and can be up to ~7 µm long. It was expected that these processes would

154

penetrate these pores as the pore diameter range in size from about 50 to

250 nm. However, instead of going through pores, tips of processes were

observed to expand into lamella-like structures (~400 nm diameter) that

span the opening of large pores. Figure 6.3 (c) shows a single lamella-like

structure at the tip of a process. Recent studies have reported that

processes of human osteoblast-like cells are able to penetrate porous

alumina with 200 nm average pore diameter19.

155

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.3. S

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156

6.3.2 Morphology of fibroblasts on patterned NST

SEM micrographs of L-cells cultured for 3 days on arrays of patterned

NST pad are shown in Figure. 6.4. Generally, cells have flat and spread out

morphology similar to those on blanket layers. However, a surprising

observation was the influence on cell shape by arrays of patterned NST

pads. It was found that fibroblast cells may attach to and take the shape of

NST pads, Figures 6.4 (b) and (c) are SEM micrographs of a cell that has

taken the shape of a 20 µm pad. It may be observed that the cell is flattened

with lamellapodia wrapping the periphery of the pad – inset is a tilted view.

Figure 6.4 (d) shows a cell that has completely taken the shape of a square

pad. Such a strong interaction between cells and NST pads was not

expected.

Shape of single cells on synthetic surfaces had been shown to be

modulated by organic cues on surfaces5,20,21. Attachment of cultured cells to

synthetic surfaces is mediated by adsorption of proteins from serum in the

culture medium5,22-24. Hence by directing adsorption of proteins to spatially

defined regions using organic cues that had been deposited on surfaces

prior to cell seeding, it had been demonstrated that cell shape could be

controlled5,20,21. In the present study, it is speculated that because of the

large surface area of the NST, protein adsorption on pads is significantly

greater than on smooth silicon dioxide mask. In addition, it is likely that

157

proteins adsorbed on surface of NST do not lose any their activity.

Increased protein adsorption then enhances cell attachment on titania pad

which results in the observed cell patterning. However, because protein

adsorption is not specific to the NST pads, cells also attached and spread

on the silicon dioxide mask. Another factor that may influence the

attachment of cells on surfaces is surface morphology. It is possible that the

topography of pads enhance the attachment of cells.

Figure 6.4. Scanning electron microscopy images of fibroblasts seeded on patterned NST.

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6.4 Conclusions

In conclusion, we have investigated the attachment of mouse fibroblast

cells on PECVD SiO2, NST and sputtered titania. It was found up to about

18 hrs, significantly more attachment of fibroblast cells occur on NST than

on PECVD SiO2 as indicated by area coverage of cells on these surfaces.

However after 24 hr, the area coverage of cells on all surfaces is similar.

Overall, the area coverage of cell on NST and sputtered titania surfaces is

similar at all times. This enhanced initial attachment of cells is suggested to

be due to enhanced adsorption of protein on titania surfaces. It is speculated

that adsorbed proteins on titania surfaces do not lose their activity.

Morphology of cells on titania surfaces are flat and spread out. When

surfaces consisting of patterned NST pad arrays with PECVD SiO2 mask,

some fibroblast cells took the shape of pads.

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6.5 References

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3. L. Leoni, T. A. Desai, Advanced Drug Delivery Reviews, 56, 211 (2004).

4. J. H. Lee, J. W. Jung, I. K. Kang, H. B. Lee, Biomaterials, 15, 705 (1994).

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8. H. Hu, Y. Ni, V. Montana, R. C. Haddon, V. Parpura, Nano Letters, 4, 508 (2004).

9. P. H. Weigel, R. L. Schnaar, M. S. Kuhlenschmidt, E. Schmell, R. T. Lee, Y. C. Lee, S. Roseman, Journal of Biological Chemistry, 354, 10830 (1979).

10. V. H. Pan, Y. Hanein, D. Leach-Scampavia, K. F. Böhringer, B. D. Ratner, D. D. Denton. Proceedings of the IEEE Conference on Micro Electro Mechanical Systems (MEMS), Interlaken, Switzerland, January 21-25, 435 (2001).

11. H. G. Craighead, C. D. James, A. M. P. Turner, Current Opinion in Solid State and Materials Science, 5, 177 (2001).

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14. T. Albrektsson, P. I. Brånemark, H. A. Hansson, B. Kasemo, K. Larsson, I. LundstrÖm, D. H. McQueen, R. Skalak, Annals of Biomedical Engineering, 11, 1 (1983).

15. D. M. Brunette, P. Tengvall, M. Textor, P. Thomsen, Titanium in medicine: materials science, surface science, engineering, biological responses and medical applications. Springer-Verlag, New York, (2001).

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17. L. V. Domnina, O. Y. Ivanova, L. B. Margolis, L. V. Olshevskaja, Y. A. Rovensky, J. M. Vasiliev, I. M. Gelfand, Proceedings of The National Academy of Sciences – USA, 69 248 (1972).

18. J. Folkman, A. Moscona, Nature 273, 345 (1978).

19. M. Karlsson, , E. Pålsgård, L. Wilshaw, L. Di Silvio, Biomaterials, 24, 3039 (2003).

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21. A. Kumar, A., N. L. Abbott, E. Kim, H. A. Biebuyck, G. M. Whitesides, Acc. Chem. Res., 28, 219 (1995).

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161

Chapter 7: Conclusions and future work

7.1 Conclusions

A novel process has been developed to integrate crack-free crystalline

nanostructured titania (NST) into microsystems. This process involves

reacting patterned Ti films with aqueous hydrogen peroxide (H2O2) solution

followed by annealing. NST formed is porous with walls of pores having

thicknesses and pore diameters ranging from 25 nm – 50 nm and 50 nm –

200 nm. Crack elimination is achieved by oxidizing Ti films, pre-patterned

below a threshold dimension, in aqueous hydrogen peroxide solution. NST

formed is amorphous but transformed to anatase after annealing at 300 °C

for a few hours.

Oxidation kinetics of Ti films occurs by nucleation and growth

mechanism. It was found that grain size and thickness of films affects the

oxidation kinetics. Grain size was found to affect both nucleation and growth

stages. Films with finer grain size have shorter nucleation period and higher

growth rate. For thin films (25 and 50 nm), growth occurs at a constant rate

until oxidation in complete. For thicker films (100, 150 and 200 nm films),

growth rate decreases after a certain thickness of porous titania has been

formed. This change in rate is attributed to a change in the mechanism

controlling growth of the oxide layer. During the initial period of growth (sub-

162

stage I) growth of oxide is controlled by reaction of Ti species with hydrogen

peroxide molecules. During the later stage (sub-stage II), diffusion of

hydrogen peroxide molecules through the oxide layer is the rate controlling

mechanism.

Functionality of the NST formed and compatibility of the process

developed were demonstrated by exploring three applications. First,

prototype gas sensors were fabricated on both Si and plastic substrates.

These sensors were able to sense hydrogen and oxygen gases at parts per

million levels. Second, integrated micrometer-scale interpenetrating Au-NST

network nanocomposites were fabricated. Although Au has been used in

this work, other metals amenable to electroless deposition would be

expected to work as well. Patterned mictometer sale Au-NST

nanocomposite features have been integrated into MEMS micro-switches as

contact materials. Third NST features have been used as porous cell

adhesion layer in devices. Initial attachment of fibroblast cells on NST in

greater than commonly used PECVD-deposited SIO2.

7.2 Future work

Although the NST formed has been characterized using a variety of

techniques in this dissertation study, further characterization work is

163

required. One aspect that has not been studied thoroughly is porosity. In

future work it is suggested that gas adsorption studies be done to

characterize porosity. This could be done in conjunction with small angle x-

ray scattering for smaller pore sizes. Furthermore, other methods of

controlling pore size, besides Ti film thickness, need to be investigated. In

addition, the mechanical properties of NST have not been studied. The

mechanical properties could be investigated by growing patterned NST on Ti

thin foils and performing stress strain test.

In the area of applications, methods need to be developed to dope the

NST. Doping or incorporating metallic species for example could be

beneficial for gas sensing applications. Functionalizing of NST surfaces with

chemical moieties has already been done in conjunction with the research

group of Professor Moskovits and has shown positive initial results.

However, more work is required in this area to show the functionality of

these grafted molecules. In the area of wear resistant contacts, work is

required to generate data on wear properties of nanocomposite at various

device operating regimes. In addition, the contact resistance of these

nanocomposites needs to be investigated. Further work is also required

towards implementing NST in BioMEMS application. One area that needs

further study is adsorption of protein on NST. In particular, work is required

to characterize kinetics of proteins adsorption.

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