relation between toughness and molecular coupling at cross

University of Wollongong Thesis Collections

University of Wollongong Thesis Collection

University of Wollongong Year

Relation between toughness and

molecular coupling at cross-linked

polymer/solid interfaces

Michaela TymichovaUniversity of Wollongong

Tymichova, Michaela, Relation between toughness and molecular coupling at cross-linkedpolymer/solid interfaces, PhD thesis, School of Mechanical, Materials and Mechatronics,University of Wollongong, 2005. http://ro.uow.edu.au/theses/463

This paper is posted at Research Online.

http://ro.uow.edu.au/theses/463

RELATION BETWEEN TOUGHNESS AND MOLECULAR

COUPLING AT CROSS-LINKED POLYMER/SOLID INTERFACES

A thesis submitted in fulfillment of the

requirements for the award of the degree

Doctor of Philosophy

from

UNIVERSITY OF WOLLONGONG

by

MICHAELA TYMICHOVA

FACULTY OF ENGINEERING

2005

i

CERTIFICATION

I, Michaela Tymichova, declare that this thesis, submitted in fulfilment of the requirements

for the award of Doctor of Philosophy, in the Faculty of Engineering, University of

Wollongong, is wholly my own work unless otherwise referenced or acknowledged. The

document has not been submitted for qualifications at any other academic institution.

Michaela Tymichova

19 August 2005

iii

Table of Contents Page

Certification i

Table of Contents iii

List of Tables vi

List of Figures vii

Abstract xi

Acknowledgements xv

1. INTRODUCTION 1

1.1 Introduction 3

1.2 References 7

2. SYNTHESIS AND CHARACTERISATION OF POLYMERIC SYSTEMS 9

2.1 Functionalisation of Polystyrene Chains 11

2.1.1 “Living” Free Radical Polymerisation Techniques 11

2.1.1.1 Nitroxy-mediated Free Radical Polymerisation 12

2.1.1.2 Atom Transfer Radical Polymerisation 14

2.1.2 Chemical Modification of Polystyrene Chain Polymer Analysis 16

2.1.2.1 Bromination of Polystyrene 16

2.1.2.2 Conversion of Brominated Polystyrene 18

2.2 Polymer Analysis 20

2.2.1 Gel Permeation Chromatography 20

2.2.2 Nuclear Magnetic Resonance Spectroscopy 24

2.3 Results and Discussion 26

2.3.1 Copolymers prepared by NMP 26

2.3.2 Copolymers prepared by ATRP 26

2.3.3 Analysis of Brominated Polystyrene

by 13C-NMR Spectroscopy 27

2.3.4 Monitoring Process of Converting PS to PS-NH2

iv

by 1H-NMR Spectroscopy 30

2.4 References 35

3. POLYMER CHAINS AS COUPLING AGENTS 37

3.1 Coupling Agents 39

3.1.1 Commercially-available Coupling Agents 39

3.1.2 Mechanism of Coupling through Silane Coupling Agents 43

3.1.3 Nature of Bonding with Silane Coupling Agents 46

3.1.4 Self-Assembled Monolayers 47

3.1.5 Mixed Monolayers 49

3.2 Adsorption of Polymer Chains 51

3.2.1 Competitive Adsorption of Polymer Chains 53

3.2.2 Physisorption vs. Chemisorption 56

3.3 Tethered Polymer Chains 56

3.3.1 Grafting to vs. Grafting from Techniques 58

3.4 Techniques for Synthesis of Grafted Polymer Chains 59

3.4.1 Grafting of Polymer Brushes using NMP 60

3.4.2 Grafting of Polymer Brushes using ATRP 62

3.4.3 Grafting of Polystyrene chains and their Functionalisation 63

3.5 Analysis of Grafted Surfaces 65

3.5.1 Ellipsometry 65

3.5.2 Surface Energy by Contact Angle Measurement 68

3.6 Results and Discussion 71

3.6.1 NMP Results 71

3.6.2 ATRP Results 73

3.6.3 Bromination of PS-SiClMe2 and Deposition of Mixed

PSBr0.1-SiClMe2/PS-SiClMe2 monolayers 77

3.7 References 85

v

4. INTERFACIAL THOUGHNESS MEASUREMENTS 89

4.1 Mechanism of Adhesion 91

4.1.1 Donor-Acceptor Interactions 92

4.2 Interfacial Fracture Toughness 94

4.2.1 Modes of Fracture 95

4.3 Thermosets 97

4.4 System Studied 99

4.4.1 Mechanism of Cross-linking 99

4.4.2 Characterisation of Epoxy System 102

4.5 Fracture Mechanism of Thermosets 103

4.6 Interfacial Toughness Measurements 104

4.6.1 Asymmetric Double Cantilever Test 107

4.7 Tailoring the Interfacial Toughness using

Polymeric Coupling Agents 108

4.8 ADCT Experiments and Results 111

4.9 References 121

5. DISCUSSION 123

5.1 Discussion 125

5.2 References 135

6. CONCLUSION 137

6.1 Conclusion 139

6.2 Suggestions for Future Work 141

6.3 References 143

APPENDICES 145

vi

List of Tables

Page

II-1: Copolymerisation conditions 15

II-2: Examples of packaging materials for high resolution GPC 23

III-1: Examples of silanes with different non-hydrolysable groups 42

III-2: Surface Tension of Test Liquids 70

III-3: PS and PS/PHEMA brushes 71

III-4: PS brush on Si wafer prepared by NMP 73

III-5: Grafting of PS chains 74

III-6: PS brush on Si wafer 74

III-7: Thickness of PHEMA-TMS/PS brush 76

III-8: Mixed PSBr(0.1)SiClMe2 and PS-SiClMe2 layers (Exp. 1) 78




III-12: Change in thickness of the polymer layer after Heck’s reaction (Exp. 1) 82




IV-1: Cure regime of DGEBA/DMP system 102

IV-2: Gc and sample parameters 111

vii

List of Figures

Page

Figure 2-1: PS/PS(NH2), PS/PHEMA and PS/PGMA copolymers. 11

Figure 2-2: 1-phenyl-1-(2’,2’,6’,6’-tetramethyl-1’-piperidinyloxy)ethane. 12

Figure 2-3: NMP mechanism. 13

Figure 2-4: Bromination of polystyrene. 17

Figure 2-5: Formation of C-N bond using Heck reaction. 18

Figure 2-6: Introduction of -NH2 groups along the PS chains. 19

Figure 2-7: Schematic example of a chromatogram. Xi is an amount

of material eluted, Mi represents different fractions. 21

Figure 2-8: An example of a calibration curve [www.sdk.co.jp]. 22

Figure 2-9a: 13C-NMR spectrum of pure PS-SiClMe2. 27

Figure 2-9b: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2. 28

Figure 2-9c: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2. 28

Figure 2-9d: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2. 29

Figure 2-10: Conversion of PS to PS(NH2) . 30

Figure 2-11a: 1H-NMR spectrum of PS (MW=200K). 31

Figure 2-11b: 1H-NMR spectrum of PS(Br)x (x = 0.1). 31

Figure 2-11c: 1H-NMR spectrum of PS-N(Si(Me3)2). 32

Figure 2-11d: 1H-NMR spectrum of PS-NH2. 32

Figure 2-12a: 1H-NMR spectrum of PS-NH2. 33

Figure 2-12b: 1H-NMR spectrum of PS-NH2 with D2O added. 33

Figure 3-1: Schematics of coupling mechanism using coupling agents. 40

Figure 3-2: Functionalities of silane hydrolysable groups . 41

Figure 3-3: Examples of short and long chain silane molecules. 42

Figure 3-4: Hydrolysis of a silane coupling agent with three methoxy

end-functional groups. 44

Figure 3-5: Formation of oligomers and their condensation to the substrate. 44

viii

Figure 3-6: a) H-bonding to the substrate, b) consequent covalent bond

formation upon curing. 44

Figure 3-7: Bonding of γ-APS to silicon substrate. 46

Figure 3-8: Schematics of the silanation reaction. 48

Figure 3-9: Conformations of adsorbed polymer chains. 51

Figure 3-10: General adsorption isotherm of a polymer with a high affinity

for a substrate. 52

Figure 3-11: Pancake, mushroom and brush conformations. 57

Figure 3-12: Chemisorption of surface active initiators to Si substrate. 60

Figure 3-13: a) deposition of mixed initiators; b) polymer brush. 61

Figure 3-14: a) Deposition of γ-APS and n-BTMS on Si wafer

b) ATRP agent attachment. 62

Figure 3-15: Chemically adsorbed PSBr and PS chains on Si wafer. 63

Figure 3-16: Heck reaction using LiN(SiMe3)2. 63

Figure 3-17: Catalysts for Heck reaction: a) Pd(dba)2 , b) P(t-Bu)3. 64

Figure 3-18: Conversion of the PSBr to the silane protected PS. 64

Figure 3-19: Conversion of the silane terminated PS to PS-NH2. 64

Figure 3-20: Single-film model. ϕ is an angle of incident and reflected beam,

n and k are the real and imaginary parts of a refractive index

of a substrate and the film, t is the film thickness. 65

Figure 3-21: Schematics of an analysing section of AutoEl-II Ellipsometer. 66

Figure 3-22: Double-film model. ϕ is an angle of incident and reflected beam,

n and k are the real and imaginary parts of a refractive index

of a substrate and the film, t is the film thickness. 67

Figure 3-23: a) So > 0, b) So < 0. Θ is an equilibrium contact angle. 69

Figure 3-24: γ-APS thickness vs. γ-APS concentration deposited on Si wafer. 75

Figure 3-25: Bonding of Si to epoxy resin via PS-NH2 chains. 84

Figure 4-1: Adsorption of basic PMMA onto acidic silica from basic, neutral

and acidic solvents. 93

Figure 4-2: Modes of loading: a) mode I, b) mode II, c) mode III . 96

ix

Figure 4-3: Schematics of a) stable and b) stick/slip mode of crack propagation. 98

Figure 4-4: Structures of epoxy and diamino curing agents. 100

Figure 4-5: Mechanism of Cross-linking between di-epoxy resin

and diamino hardener system. 101

Figure 4-6: Tg of various cure regimes of DGEBA/DMP measured by DMA. 103

Figure 4-7: Schematics of the “sandwich” structure for testing

PS/glass interface. 105

Figure 4-8: Schematics of the “sandwich” structure for testing

PS/silicon interface. 106

Figure 4-9: Asymmetric double cantilever beam specimen; Gc = f (E, h, ∆2, a4). 108

Figure 4-10: Photographs of an ACDT specimen. 109

Figure 4-11: Monolayer thickness of the modified Si surfaces (Exp 1). 112

Figure 4-12: Surface energy of the modified Si surfaces (Exp 1). 112

Figure 4-13: Interfacial fracture toughness measurements using ADCT (Exp 1). 113









xi

Abstract

The relationship between the interfacial fracture toughness (Gc) and molecular

coupling between epoxy polymers and silicon wafers was studied using the asymmetric

double cantilever test (ADCT). In order to investigate the molecular coupling, the coupling

molecules had to be applied along the polymer/substrate interface in various concentrations.

The grafting from technique by means of the “living” free radical polymerisation

techniques, namely nitroxy-mediated “living” free polymerisation (NMP) and atom transfer

radical polymerisation (ATRP), were used to prepare suitable coupling molecules.

Unfortunately, these techniques did not produce satisfactory results, and the new route of

grafting to was investigated. This method involved chemical attachment of brominated

polystyrene chains (PS) to the silicon substrate. The bromine functional groups of the

brominated polystyrene (PS(Br)) were then further converted to amino (-NH2) groups using

the Heck reaction, producing PS(NH2).

The conversion method was first tested using the free (unattached) polystyrene

(MW = 200K) which was brominated (molar fraction of brominated units, x = 0.1). The

conversion from PS to PS(Br) and then to PS(NH2) was monitored by proton and carbon

nuclear magnetic resonance techniques (1H-NMR and 13C-NMR) and gel permeation

chromatography.

For the grafting to experiments, monochlorosilane end-functionalised polystyrene

(PS-SiClMe2) (MW = 8000) was used. Various ratios of PS and PS(Br) ranging from 0 to

20% PS(Br) were deposited onto silicon substrates. Applying the Heck reaction, Br groups

were converted to NH2, producing surfaces with different concentrations of amino groups

which were expected to react with the epoxy polymer.

ADCT was adapted to investigate the dependence of the interfacial fracture

toughness on the degree of coupling. The dependence of Gc on the concentration of the

coupling molecules was not directly confirmed. This was attributed mainly to the

xii

challenges in deposition of the polymer chains and the final conversion to PS(NH2).

However, the differences in Gc values between experiments suggested that Gc’s measured

by ADCT reflect the differences in surface properties of the polymer modified surfaces, and

therefore this technique is suitable for interfacial adhesion measurements between epoxy

polymers and solid substrate.

xiii

I dedicate this thesis to all people who dare to be different, to all who yearn for

achieving something that seems impossible but they do it anyway. Sometimes, the results

are not what we expected or were hoping for. It does not matter. The important thing is that

the seeds were planted.

xv

Acknowledgements

I’d like to take the opportunity to express my gratitude to many people who helped

me on my journey. Many thanks go to my supervisors prof. Hugh Brown and Dr. Chris

Lukey. To Hugh, for believing in my skills and capabilities to entrust this project to me. To

Chris, for the endless discussions, and also for his kindness and being there for me

whenever I was close to giving up. I would also like to express my appreciation to the

following people: prof. A. Whittaker (Centre for Magnetic Resonance, QUT), Dr. W. Lie

and Ms. S. Chapman (University of Wollongong) for the NMR analysis; prof. S. Pyne

(University of Wollongong) for his expert advice in chemical synthesis and analysis; Dr. C.

Hawker (University of California Santa Barbara, USA) and Dr. M. Whittaker (PolymerRat)

for their assistance with NMP and ATRP techniques; and Prof. S. X. Dou and Dr. K.

Konstantinov (University of Wollongong) for providing a space for the air sensitive

chemical procedures.

I am very grateful to Robert Oslanec, my dearest friend and life companion, for

sharing his experience but also for sometimes not saying anything and letting me figure it

out for myself. Special gratitude goes to Lorelle Pollard for her cheerful face and caring.

And at last but not at least, I’d like to thank all my colleagues, Sandra Cram, Nathan Jones,

Damien Jinks, Donna Capararo, Wanda Melfo, Daniel McCubbery, Haider K. Habboby,

Siu W. Wai, Nicolas L. Baut, Dominic Phelan, Kirsten Carpenter, and Mark Reed for their

sharing and comradeship.

This research was funded by ARC grant (Adhesion Fundamental Award).

CHAPTER 1

Introduction

3

1.1 Introduction

Cross-linked glassy polymers (unsaturated polyesters, epoxy resins, phenolic

resins, etc.) are materials widely used as adhesives, coatings, composites and

encapsulants. Some examples of epoxy resins are bisphenol-A-glycidyl ether oligomers,

novolacs, aliphatic glycidyl ethers and cycloaliphatic epoxies. Upon curing, these

polymers form insoluble 3-D networks. Various curing agents are used, including

aliphatic amines, aliphatic amine polyamides, aromatic amines, cyclic polyacid

anhydrides, dicyandiamides and others. In this study, a monomeric bisphenol-A-

glycidyl ether (DGEBA) and diamino methylpentane cross-linker (DMP) are used (see

Chapter 4). Epoxy resins have excellent resistance to many chemicals and high

temperatures. Because of these properties, they are widely used as adhesives for metals,

ceramics, glass, concrete and wood. They are also used in high performance coatings,

water resistant paints, solvent free paints, electrical insulators and fibreglass

composites. Interfacial adhesion is an important property in all these applications, so

improving the adhesion properties of these polymers would greatly increase their

applicability in various situations.

One of the important applications of interfacial adhesion between epoxy resins

and inorganic materials is in the manufacture of printed circuit boards (PCB). PCB

technology is based on glass reinforced epoxy laminates. The processing involves

mechanical and chemical treatments (drilling, degreasing, drying, copper plating,

rinsing, stripping, etching, immersion in molten tin, cleaning and humidity testing)1.

When water or moisture penetrates into the interfacial regions it causes delamination

and damages the boards. It has been shown, that adding an appropriate silane, as a

replacement for chrome treatments, creates bonds that withstand boiling water

conditions and therefore improves the interfacial adhesion strength between the glass

fibres and epoxy1.

To improve the adhesion strength between various materials, adhesion

promoters, also called coupling agents, are often used. Organofunctional silanes are one

4

of the most versatile coupling agents, used to bond organic and inorganic materials.

They are hybrids of silica and organic materials (see Chapter 3). There are a large

number of different silanes available, and they have to be carefully chosen for a

particular application. The best interfacial bond strength is usually achieved by

choosing a silane with the organofunctional group that has maximum reactivity with a

particular resin during cure.

Finding the optimal types of silanes and methods of their application to the

various surfaces has been well researched1-10. However, it is not known how much

coupling agent is required in order to achieve the desired adhesion strength. This is

because of the lack of understanding of the mechanism of interfacial adhesion between

resin/inorganic substrate on a molecular level. To the author’s best knowledge, there is

little understanding about the amount of a chemical reaction (concentration of coupling

molecules along the interface) needed to achieve a particular adhesion strength between

a solid substrate and a cross-linked polymer. This applies to other epoxy/inorganic

material applications also, including structural laminates, adhesives and coatings.

The aim of this study was to investigate the mechanism of coupling, and the

relationship between interfacial toughness and the areal density of covalently bonded

coupling molecules. In thermoplastics (polystyrene, polymethylmethacrylate, etc.),

coupling across an interface is mainly due to chain entanglements on both sides of the

interface. In cross-linked systems, strong adhesion is more likely to be achieved by a

chemical reaction between the thermoset and the coupling molecules. In order to gain a

quantitative understanding of the effects of the coupling agents on the interfacial

strength between a solid inorganic substrate and a cross-linked polymer (DGEBA), a

system had to be designed where only covalent bonds would contribute to the adhesion

strength.

A considerable amount of work was dedicated to finding macromolecules of a

suitable, well-defined size and structure, which would serve as a chemical bridge

between the substrate and the polymer. The first logical choice would be to use silane

5

coupling agents because of their ability to react strongly both with the inorganic

substrate and the cross-linked polymer. However, because of their single organic

functionality and undesirable behaviour during deposition (bonding of the functional

group to the surface - see Chapter 3), non-polymeric silanes were not suitable for this

study. Multiple organic functionality on each chain gives a better chance for each

molecule to react with the polymer. This requirement is an essential part of the

quantitative study of interfacial toughness.

Several avenues were explored in order to prepare suitable coupling molecules

with multiple functionalities. Initially, it was thought that dendrimers or hyperbranched

polymers, extensively studied by Tomalia et. al.11,12, might be used. The advantage of

these macromolecules is that their size, shape, surface chemistry and topology can be

controlled during the synthesis. The main disadvantage is that the preparation of these

molecules in a controlled manner is quite demanding.

Another avenue was to grow polymer chains with particular functional groups

directly from the substrate, referred to as “grafting from” method. This can be achieved

by controlled “living” free radical polymerisation techniques (CRP). These synthetic

routes have the advantage of producing polymers with well-defined molar mass and

structure. In the last several years, a technique known as nitroxy-mediated “living” free

radical polymerisation has been studied by Hawker to produce well-defined

macromolecules called polymer brushes13. Another CRP technique investigated in this

study was atom transfer radical polymerisation (ATRP). Compared to nitroxy-mediated

living free radical polymerisation, this technique allows a wider variety of monomers to

be co-polymerised.

An alternative way to prepare the multifunctional polymer coupling molecules is

a method of “grafting to”, where chemically modified polymer chains are chemically

attached to the substrate. A commercial silane-terminated polystyrene of very low

polydispersity was chemically modified to produce a polystyrene-based coupling agent

with multiple amino groups along the chain14-17. This procedure was tested using free

6

polymer chains. Each step of the chemical modification was monitored using 13C-NMR

and 1H-NMR techniques (Chapter 2). This last approach was the most successful in

preparing silicon oxide surfaces modified with various amounts of coupling agent with

amino functionalities, and therefore it was used for the adhesion studies of the

interfacial fracture toughness between epoxy (DGEBA) and silicon oxide.

It is known that polystyrene/silicon oxide and polystyrene/epoxy resin adhesion

strength is very weak18. By choosing this system and developing a coupling agent based

on polystyrene, the effect of hydrogen bonding, dipolar and acid-base interactions, that

otherwise contribute to the interfacial toughness, were minimised. Silane end-functional

polystyrene chains were chemically modified by attaching randomly distributed amino

groups. These molecules were used as coupling agents for studying the interfacial

toughness between a silicon substrate and epoxy resins. The multiple functionality of

the coupling molecules ensured that each molecule reacts with the epoxy resin during

the curing procedure.

In order to understand the effect of molecular coupling on adhesion toughness,

the degree of coupling across the interface has to be controlled and evaluated. The

amount of coupling can be calculated from the amount of coupling agent deposited onto

a substrate. In this work, ellipsometry was used to determine the layer thickness of the

coupling molecules. From the thickness of the layers, the chain areal density can be

calculated.

The relationship between interfacial toughness and molecular coupling in a

cross-linked system was studied by applying the methods of fracture mechanics. One

end of the silane end-functional aminopolystyrene chains was covalently attached to the

silicon wafer, and the other end is glued to a pre-cured epoxy beam using the same

epoxy resin (see Chapter 4). The Asymmetric double cantilever test (ADCT),

previously used by Smith and Brown18-21 to study interfacial toughness of non-

crosslinked polymers, was modified and applied to measure the interfacial fracture

toughness. In this study, some characteristics of the testing samples were altered in

7

order to overcome difficulties associated with brittleness of the silicon wafers, and to

ensure that the crack propagated along the silicon oxide/polymer interface.

1.2 References

1 E. P. Plueddemann, Silane Coupling Agents (Plenum Press, 1982). 2 B. Arkles; Vol. 7 (Petrarch Systems, Inc., 1977). 3 J. B. Brzoska, J. B. Azouz, and F. Rondelez, Langmuir 10, 4367-4373 (1994). 4 J. Duchet, J.-F. Gerard, J. P. Chapel, and B. Chabert, Composite Interfaces 8, 177-

187 (2001). 5 J. Duchet, J. P. Chapel, B. Chabert, and J. F. Gerard, Macromolecules 31, 8264-

8272 (1998). 6 J. Duchet, B. Chabert, J. P. Chapel, J. F. Gerard, J. M. Chovelon, and N. Jaffrezic

Renault, Langmuir 13, 2271-2278 (1997). 7 C. R. Kessel and S. Granick, Langmuir 7, 532-538 (1991). 8 A. J. Kinloch, Structural Adhesives: Developments in Resins and Primers (Elsevier

Applied Science Publishers, London). 9 M. E. McGovern, K. M. R. Kallury, and M. Thompson, Langmuir 10, 3607-3614

(1994). 10 P. Silberzan, L. Leger, D. Ausserre, and J. J. Benattar, Langmuir 7, 1647-1651

(1991). 11 D. A. Tomalia, Scientific American May, 62-66 (1995). 12 D. A. Tomalia, A. M. Naylor, and W. A. Goddard III, Angew. Chem. Int. Ed. Engl.

29, 138-175 (1990). 13 C. J. Hawker, Acc. Chem. Res. 30, 373-382 (1997). 14 R. P. Kambour and J. T. Bendler, Macromolecules 19, 2679-2682 (1986). 15 S. Lee, M. Jorgensen, and J. F. Hartwig, Organic Letters 3, 2729-2732 (2001). 16 V. C. Gibson and W. Reed, in Europian Patent Application (UK, 1998), p. 1-9. 17 K. Suzuki, K. Yamaguchi, A. Hirao, and S. Nakahama, Macromolecules 22, 2607-

2611 (1989).

8

18 J. W. Smith, E. J. Kramer, F. Xiao, C.-Y. Hui, W. Reichert, and H. R. Brown, J. of

Mat. Sci. 28, 4234-4244 (1993). 19 J. W. Smith, E. J. Kramer, and P. J. Mills, J. Polymer Sci.:Part B: Polymer Physics

32, 1731-1744 (1994). 20 H. R. Brown, Macromolecules 24, 2752-2756 (1991). 21 H. R. Brown, K. Char, V. R. Deline, and P. F. Green, Macromolecules 26, 4155-

4163 (1993).

CHAPTER 2

Synthesis and Characterisation

of Polymeric Systems

11

2.1 Functionalisation of Polystyrene Chains

Physical and chemical properties of polymer chains can be tailored by

introducing various functional groups along the chain. This can be achieved either by

chemical modification of an existing chain or by a synthetic route of copolymerisation.

Traditionally, copolymers are prepared by free radical or ionic polymerisation

techniques. When polymers with well-controlled properties are desired, “living” free-

radial polymerisation techniques can be applied.

2.1.1 “Living” Free Radical Polymerisation Techniques

“Living” free radical polymerisation techniques have been developed in the last

decade to synthesise molecules with well defined structure, physical and chemical

properties. Polymers with controlled molecular mass, desired functionality and low

polydispersity can be prepared without using the demanding reaction conditions of

anionic and cationic polymerisation techniques.

There are currently three types of “Living” free radical polymerisation

techniques: nitroxy-mediated free radical polymerisation (NMP), atom transfer radical

polymerisation (ATRP) and reversible addition fragmentation chain transfer (RAFT).

All of them follow the same principles, but use different types of the initiators and

therefore the mechanism varies. The NMP and ATRP are techniques used in the current

study, and are briefly described below. Styrene/4-aminostyrene (PS/PS(NH2)),

styrene/2-hydroxyethyl methacrylate (PS/PHEMA) and styrene/glycidyl methacrylate

(PS/PGMA) were copolymerised (Fig. 2-1) in order to create copolymers specifically

designed to react with the DGEBA/DMP system through their functional groups (-NH2,

-OH and –OCHCH2).

Figure 2-1: PS/PS(NH2), PS/PHEMA and PS/PGMA copolymers.

12

2.1.1.1 Nitroxy-mediated Free Radical Polymerisation

Nitroxy-mediated free radical polymerisation (NMP) uses TEMPO (2.2.6.6-

tetramethylpiperidinyloxy)-based unimolecular initiators (Fig. 2-2) with a thermally

unstable alkoxyamine (C-ON) bond to reversibly react with the growing polymer chain

end. The unimolecular initiator undergoes thermal fragmentation or homolysis to give a

stable nitroxide and a polymer radical. When the polymer radical reacts with the

monomer added to a solution, the result is the polymer radical with an increased degree

of polymerisation through monomer addition. The extended polymeric radical then

recombines with the nitroxide to form the inactive molecule, and the cycle repeats (Fig.

2-3)1.

Figure 2-2: 1-phenyl-1-(2’,2’,6’,6’-tetramethyl-1’-piperidinyloxy)ethane.

The nitroxide free radical reacts with the polymeric radical but it does not

initiate the growth of any other polymer chains. Because the concentration of these

radical chain ends is low there is a decreased possibility of side reactions (termination,

disproportionation and combination) compared to the traditional free radical

polymerisation techniques. This behaviour leads to the possibilities of preparing a wide

range of molecules. For example styrene based homopolymers (polystyrene and poly(4-

hydroxystyrene))2,3, random and block copolymers (PS/PMMA, PS/PHEMA)1, complex

macromolecular architectures4 (grafts, dendrimers, hyperbranched polymers) and other

unique macromolecular structures (dendritic-linear block copolymers) 5. All these novel

molecular structures are finding applications in surface science, adhesion, coating

technologies, microelectronics and biotechnology. A great advantage of this technique

is its simple, undemanding reaction conditions with no need for laborious reagent

purification.

13

Figure 2-3: NMP mechanism.

Controlled synthesis of random copolymers of polystyrene with poly(2-

hydroxyethyl) methacrylate (PS/PHEMA), poly(glycidoxy) methacrylate PS/PGMA

and aminopolystyrene (PS/PS(NH2)) were carried out according to the procedures

published by Hawker et. al.4. A mixture of the monomers (90% styrene, 10% other

monomer) and the unimolecular initiator (Fig. 2-2) were placed in a round bottom flask

sealed with a septum, placed in an oil bath and heated at 125oC for 24 hour under an

inert atmosphere. The resulting copolymers, which were either solid or very viscous,

were dissolved in CH2Cl2 and precipitated into a large amount of methanol, then dried

in an oven at 60oC under vacuum. This process of dissolving, precipitating and drying

was repeated 3 times to remove any unreacted monomers and solvent. The molecular

mass of the polymer is given by the molar ratio of monomers to the unimolecular

initiator. In the present work, 500:1 and 200:1 molar ratio of monomer to initiator were

used. The polymers were analysed by GPC and NMR.

14

2.1.1.2 Atom Transfer Radical Polymerisation

While NMP proved to be most successful for synthesis of styrene based

polymers, atom transfer radical polymerisation (ATRP) is also suitable for

copolymerisation of other monomers like methacrylates, methacrylamides and

acrylonitriles. In ATRP, the active species (radicals*) are generated by a reversible

redox process catalysed by a transition metal complex (Mtz−Y/Ligand). The oxidation

state of the metal complex changes by one and the halogen atom (X) is removed from

the dormant species Pn−X to give a carbon centred Pn* radical and the metal complex

X−Mtz+1−Y/Ligand6. This carbon-halogen cleavage is a reversible homolytic process.

Polymer chains grow by addition of monomers to the intermediate polymer radicals Pn*.

The mechanism of growing the polymer chains is shown below:

Pn−X + Mtz−Y/Ligand ↔ Pn* + X−Mtz+1−Y/ Ligand

Pn* + Monomer → Pn−Pm

Cu-based ATRP, which was used in the present study, is the most versatile

technique in the field of “living” polymerisation7. A wide range of polymers such as

acrylates, methacrylates, styrenes, acrylonitriles, acrylamides and vinylpyridines and

their copolymers have been synthesised using Cu-based catalyst and ligands with

amine, pyridine and imine structures6-10.

ATRP polymerization procedures were based on the work of Matyjaszewski7.

Styrene (St) was homopolymerized and co-polymerized with 2-(trimethyl(silyloxy)ethyl

methacrylate (HEMA-TMS) and glycidyl methacrylate (GMA). Ethyl-2-

bromoisobutyrate and CuBr were used as an initiator and a catalyst respectively. Two

different ligands were used, 2,2’-bipyridine (Bipy) and tris(2-

dimethylaminoethyl)amine (Me6-TREN). The advantage of Me6-TREN, compared to

bipyridine, is that it is homogeneous. A complex is created with CuBr which is soluble

in the polymerization solution. This is very important in the case of grafted copolymers

where the system is already heterogenous because of the solid substrate. All materials

were supplied by Aldrich and used without future purification unless otherwise

indicated.

15

Monomer(s), initiator, catalyst, ligand and solvent were mixed in appropriate

ratios and placed in the vials sealed with a septum. Where Bipy was used, the ratio of

monomers (M) : initiator (I) : catalyst (C) : ligand (L) was 200 : 1 : 1 : 2.5. When Me6-

TREN was used the ratio M : I : C : L was 200 : 1 : 1 : 1. The solution was degassed for

approximately half an hour with argon. The vials were then placed in an oil bath set

either to 90oC or 110oC. The polymerisation conditions are shown in Table II-1. After a

certain period of time, as shown in the Table II-1, the vials were removed from the oil

bath and cooled. The formed polymers were diluted in chloroform (1:1) and passed

through a small column filled with alumina to remove the catalyst. The polymers were

then precipitated in methanol and dried in an oven at room temperature under vacuum.

Table II-1: Copolymerisation conditions

Monomers Monomer

ratio

Ligand Solvent Temp Time

GMA/St 1/9 Me6-TREN toluene 110oC 4hrs

2/8 Bipy bulk 90oC ON*

HEMA-TMS/St 1/9 Me6-TREN toluene 110oC 9hrs

2/8 Bipy bulk 90oC ON*

* overnight

16

2.1.2 Chemical Modification of Polystyrene Chain

Many difficulties were associated with the preparation of styrene/aminostyrene

copolymer with an end-functional group from monomers using the methods described

in chapters 2.1.1.1 and 2.1.1.2. Due to the reactivity of aminostyrene monomer, the

amino group had to be protected by reacting with di-t-butylcarbonate (Boc2) prior to

copolymerisation (this procedure is not described in this thesis but can be found

elsewhere11). The side reactions of the Boc2 protecting group, incompatibility between

the monomers and inability to find a suitable solvent that was compatible with

monomers, homopolymers and copolymer caused further difficulties in purification,

handling and utilisation of the resultant mixture.

This led to a search for a method for the modification of existing end-

functionalised polystyrene chains. There is a well-established method that involves the

bromination of phenyl rings12. This relatively simple and quantitative method, described

below, allows the introduction of bromine functional groups along the chain that can be

further converted to the desired functionality.

2.1.2.1 Bromination of Polystyrene

Polystyrene (PS)/ brominated polystyrene (PS(Br)x) blends have been used in

many studies, including the miscibility and phase separation of polymer-polymer

blends12, polymer adsorption to solid substrates13 and surface segregation14. The Flory-

Huggins parameter that governs the interactions between polymers can be varied easily,

which makes the PS/ PS(Brx) blend an ideal system for such studies.

The bromination reaction is relatively easy and does not change the

polydispersity of a given polymer12, which makes it a perfect starting point for chemical

modification and/or functionalisation of PS chains. Bromination of phenyl rings is an

example of electrophilic aromatic substitution. Bartulin15 showed that under the

specified conditions in the absence of light the bromination occurs exclusively on the

aromatic ring at the para position (Fig. 2-4).

17

Figure 2-4: Bromination of polystyrene.

In the current project, the synthesis of PS(Br)x was carried out following the

procedure of Kambour and Bendler12. Polystyrene of MW = 200,000 g/mol and

polydispersity 1.06 was purchased from Polysciences, Inc. and used without further

treatment. Polymer was dissolved in nitrobenzene at a concentration of 7wt%. The

polymer solution was placed in a test tube equipped with a magnetic stirrer and closed

with a rubber septum. The tube was covered by aluminium foil to prevent the light-

catalysed free radical reactions that would produce backbone bromination. The desired

amount of bromine was added to the polymer solution and left under stirring for 24

hours at room temperature. Excess bromine was then neutralised by addition of a few

drops of 1-pentene. The quenched reaction mixture was then slowly poured into

methanol with high stirring. The precipitated polymer was filtered, washed with more

methanol, dried in a vacuum oven at 80oC and dissolved in toluene. This process was

repeated three times to remove any residual nitrobenzene and other contaminants.

The amount of bromine was calculated from the following equation:

xMMmEm Bro

PS

PSBr 2

= ,

where E is the excess factor for bromine stoichiometric amount (E = 1.20, i.e. 20%

excess), PSm is the mass of polystyrene, oPSM is the molar mass of polystyrene repeat

unit, 2BrM is the molar mass of Br2 molecule and x is the desired mole fraction of 4-

bromostyrene units in the chain.

18

As found in the previous work of Genzer14 and Oslanec13, the experimentally

determined bromination levels (by Quantitative elemental analysis) agreed well with the

degrees of bromination calculated from the formulae shown above.

2.1.2.2 Conversion of Brominated Polystyrene

The second step was to introduce amino groups along the PS chain. A method

published by Lee16 for converting aryl halides to the anilines using lithium

bis(trimethylsilyl)amide (LiN(SiMe3)2 or LiHMDS, Aldrich) and palladium-based

catalysts was adapted. This method is based on the Heck reaction, the palladium-

catalysed arylation and alkenylation of alkenes, discovered by Heck17 in 70's. This

reaction is frequently applied in the metal-catalysed C-C bond forming process. In this

study as well as in the study conducted by Lee at al.16, the Heck mechanism was used to

form C-N bond (Fig. 2-5).

Figure 2-5: Formation of C-N bond using Heck reaction.

19

Figure 2-6: Introduction of –NH2 groups along the PS chains.

This procedure was used to introduce amino groups along the polystyrene chain

(Fig. 2-6) as follows: Brominated polystyrene was dissolved in dry toluene and oxygen

was removed by bubbling nitrogen through the solution for 1/2 hour in a vial closed by

a septum. In an argon filled glove box, LiHMDS (20% solution in THF) and catalytic

amounts (0.05 mmol - 0.002mmol) of tert-butyl phosphine (P(t-Bu)3, 10wt% in hexane)

and bis(dibenzylideneacetone)palladium (Pd(dba)2, powder) were added to the polymer

solution (all chemicals supplied by Aldrich). The mixture was stirred for 24 hours.

Several drops of 1M methanolic HCl were then added and the mixture was stirred for

several minutes to convert the amino silane group to ammonium chloride (-NH3+Cl-).

Then the polymer was poured into a large excess of diethyl ether and isolated by

filtration. The solid was redispersed in THF/MeOH (9/1) solution and 1M methanolic

KOH was added18 to produce the free amino group (-NH2). The modified polymer was

precipitated into distilled water containing a few drops of triethylamine19, filtered and

dried in the oven at 60oC under the vacuum. The resulting aminopolystyrene was

analysed by GPC and NMR.

20

2.2 Polymer Analysis

Spectroscopic and chromatographic techniques are the main techniques for

polymer characterisations. In this study, Gel Permeation Chromatography (GPC) and

Nuclear Magnetic Resonance (NMR) techniques were used for structural and chemical

analyses. GPC was used to determine molar mass of modified polystyrene, to assure

that the chemical reactions involved caused no chain degradation. NMR was used to

monitor the product of each step during the chemical modification procedure.

2.2.1 Gel Permeation Chromatography

Gel Permeation Chromatography (GPC), also called size exclusion

chromatography (SEC), is a standard technique used for molar mass analysis of

macromolecules. The main advantages are that the analyses are relatively fast, only

small amounts of material are needed, and the equipment is fully automated. Another

important advantage of GPC over other methods used for molar mass analysis

(osmometry, light scattering, viscosity measurement etc.) is that it also gives molar

mass distributions. This is used for example in fractional GPC, where polydisperse

substances are separated into monodisperse fractions of different molar mass.

GPC is a special case of the liquid chromatography process, where the

stationary phase (a porous cross-linked gel) is packed in a chromatographic column and

swollen by a mobile phase (eluent). The gel consists of spherical beads with a very

small diameter and a narrow distribution size. The size of the pores should be of the

same order of magnitude as a size of particles to be separated. A dilute solution of the

polymer, dissolved in an appropriate solvent, is flushed through the stationary phase

where the molecules are separated according to their size. There must be no interactions

between the eluent and the stationary phase of the column. The separation is a diffusion

driven process. The accessibility of the pores by diffusion of different molecules is a

function of both molecular size and pore size. Retention and elution of particles in the

column is given by their hydrodynamic volume (equations 1 and 2), which can be

related to molar mass through the intrinsic viscosity of the material. Smaller molecules

penetrate deeper into the pores of the gel and therefore are eluted after the larger

21

molecules. Molecules whose average hydrodynamic radius is larger than that of the gel

pores cannot be separated.

Hydrodynamic volume Vh = (4/3)πRh3 , (1)

where Rh is

Hydrodynamic radius Rh = kbT/6πηD, (2) where kb = 13.81 x 10-24 J/K is Boltzmann's constant T [K] is temperature η is intrinsic viscosity D [m2/s] is diffusion coefficient

GPC is a relative method. Polymer standards with known molar mass and low

polydispersities are used to calibrate the instrument. The data are collected in the form

of an elution chromatogram (Fig. 2-7), which is converted into molar mass distribution

via a calibration curve. The profile of the elution curve is determined by the molar mass

distribution of the sample. The data from the chromatogram are analysed using

computer software to give number-average molar mass (Mn), weight-average molar

mass (Mw), Z-average molar mass (Mz) and polydispersity (PD).

Figure 2-7: Schematic example of a chromatogram. Xi is an amount of material eluted,

Mi represents different fractions.

Mi Time, Volume

X

Mn Mw

Mz

22

The precise correlation between elution volume and molecular size cannot be

calculated. Therefore, each column must be calibrated with polymer standards of known

molar mass (Fig. 2-8).

Figure 2-8: An example of a calibration curve [www.sdk.co.jp].

The operating range lies in the linear region that can be expressed as:

Log M = a + bVe ,

where a and b are system-specific constants.

If the calibration standards are not available, a universal calibration curve can be

used (Appendix A)20. This universal curve is based on the fact that the structure of the

polymers does not affect Ve, and that Ve is governed only by the product of the intrinsic

viscosity [η] and M for a given set of columns. Therefore:

[η]1 × M1 = [η]2 × M2 (Ve,1 = Ve,2)

It has been confirmed with different polymers that at a certain elution volume Ve, the

product ([η] x M) is constant21. The intrinsic viscosity [η] relates to molar mass M

through the empirical Mark-Houwink relationship:

[η] = KMa ,

where K and a are tabulated constants.

23

Knowledge of the two coefficients K and a is necessary in order to apply the universal

calibration method.

The GPC system consists of a chromatograph, a detector and a data acquisition

system. The liquid chromatograph is composed of a solvent reservoir with an in-line

filter, a high pressure pumping system with adjustable flow rate, an injector for sample

introduction, and a set of columns. A range of different packing materials is available.

Some examples are shown in Table II-2. The choice of columns is very important

because the gel type and its packing density determine the separation efficiency.

Usually, several columns with different pore diameter ranges are used to achieve a good

separation. The most common GPC detector is the differential refractometer that

provides a signal proportional to polymer concentration. The main advantage of this

type of detector is its non-specificity and quasi-universal application. However, it is

very sensitive to temperature variations so that the temperature has to be carefully

stabilised20. The data acquisition system converts the information from the

chromatogram to the average molar mass and molar mass distributions using the

calibration curve.

Table II-2: Examples of packaging materials for high resolution GPC 21

GPC analyses were performed using a SHIMADZU LC-10AT VP Liquid

Chromatograph equipped with 2 Waters polystyrene-packed Styragel columns (HR4

and HR2 4.6 x 300 mm) connected in series, and connected to a differential

refractometer mass detector. THF was used as the eluent at a flow rate 0.3mL/min, and

the instrument was calibrated using polystyrene standards from Polymer Laboratories.

24

Polymer solutions (1-2 mg/ml) in THF were prepared and filtered into GPC vials using

Millex 0.20 µm PTFE syringe filters.

2.2.2 Nuclear Magnetic Resonance Spectroscopy

NMR is a form of absorption spectroscopy that gives information about the

number of magnetically distinct atoms in a molecule. When a sample is exposed to a

magnetic field it can absorb electromagnetic radiation in the radio frequency region at

frequencies that are characteristic to the molecules in the sample.

Many nuclei have a spinning charge that generates a magnetic dipole along their

axis. 1H, 12C, 16O, 13C, 15N, 19F and 31P are examples of such nuclei. If such a nucleus is

placed in an external magnetic field, its spin aligns itself either along with or against the

applied magnetic field. In the magnetic field, the nucleus absorbs the radio waves of a

specific frequency, which causes it to flip from the lower to the higher energy state with

the opposite spin22. In other words, when the frequency of the external magnetic field

equals the frequency of the electric field generated by the spinning proton, then these

two fields can couple and the energy is absorbed causing a spin change. This is called

resonance. The energy absorbed equals the amount of energy between these two energy

states. A 1H NMR spectrum is a plot of the applied magnetic field strength as a function

of the intensity of the absorption by the individual protons in a molecule.

The position of the peaks also depends on the immediate electronic environment

of the absorbing proton. Chemically equivalent protons exhibit the same chemical shift

(δ). Protons with different chemical environment have different chemical shifts. This is

because the electrons in the molecule shield the proton slightly from the external

magnetic field. A reference signal of tetramethylsilane (TMS) is often used to determine

the chemical shifts in an analysed molecule. The chemical shift of the protons in TMS is

arbitrarily set to zero. TMS is used as a reference standard because it has 12 chemically

equivalent hydrogens which give a distinct peak.

The relative intensities (area under the peak) of signals are proportional to the

number of protons contributing to each signal. Each proton is also subjected to the

25

magnetic field generated by the adjacent protons. This phenomenon is referred to as

spin-spin coupling, and causes signal splitting. If a proton has n chemically non-

equivalent protons, its 1H NMR signal will be split into n + 1 peaks. Chemically

equivalent protons do not cause any splitting.

Another nucleus commonly used in NMR spectroscopy is 13C. Its natural

abundance is very low (only 1.1%) and its magnetic moment is only 1/4 that of the

proton, therefore the sensitivity of NMR detection is much lower that in the case of

proton NMR. Development of Fourier transform instruments made it possible to detect 13C and made 13C NMR spectroscopy a useful technique for determining the structure of

organic molecules and polymers.

Dissolving polymers in NMR solvents produces solutions of a high viscosity due

to the chain entanglements. The long-range motions of dissolved polymeric molecules

might be slow but the local segmental motions are usually rapid23. Therefore NMR is a

useful technique for probing the molecular structure of polymers and other

macromolecules.

In this study, NMR spectroscopy was used to monitor each step of the chemical

modification procedure of polystyrene (PS) to amino-polystyrene (PS(NH2)). 1H NMR

and 13C NMR spectra were collected using a Varian Mercury-300 MHz NMR

spectrometer. CDCl3 containing 0.03% of tetramethylsilane (TMS) as a reference signal

was used as the solvent. For 1H NMR, solution concentrations were 10-20 mg/ml, and

for 13C NMR the concentration was higher.

26

2.3 Results and Discussion

Three different paths to prepare copolymers with desired functionalities were

explored in this work - NMP, ATRP and the chemical modification of an existing

polymer chain. Even though all three methods were more less successful in preparing

the free copolymers, NMP and ATRP did not produce the satisfactory results when

applied to the technique of grafting these copolymers from the silicon surfaces.

Therefore, only a brief description of these results is given. On the other hand, the

method of chemical modification of the polystyrene chain was applied successfully in a

grafting to technique (Section 3.3) and therefore is discussed in more detail.

2.3.1 Copolymers prepared by NMP

The PS monomer, PS/PHEMA and PS/PS(NH2) copolymers prepared by NMP

(see Section 2.1.1.1) were analysed using GPC and NMR. GPC analysis, including

molecular weight determination, and NMR spectra of PS/PHEMA and PS/PS(NH2) can

be found in Appendix B. Copolymerisation of styrene/4-aminostyrene turned out to be

very complicated due to a limited solubility of 4-aminostyrene in styrene monomer.

Also, the molecular mass of the resulting copolymer was much lower then the

theoretical molecular mass calculated from the monomer/initiator ratio. Somewhat

better results were achieved by copolymerisation of BOC protected 4-aminostyrene

with styrene shown in Appendix C.

2.3.2 Copolymers prepared by ATRP

The PHEMA-TMS/PS and PGMA/PS copolymers were analysed using GPC

and NMR. For PGMA/PS copolymers, the molecular mass ranged from 4K to 54K with

polydispersities (PD) between 1.12 – 1.62. The molar mass of PHEMA/PS copolymers

varied from 10K to 37K with PD between 1.18 – 1.27. A summary of results and

reaction conditions can be found in Appendix D.

27

X = 0

2.3.3 Analysis of Brominated Polystyrene by 13C-NMR Spectroscopy

The bromination of polystyrene was monitored by 13C-NMR spectroscopy. The

starting material was polystyrene with dimethylchlorosilane end-functional group (PS-

SiClMe2, Polysciences) of molecular mass MW = 8000. Several polystyrene samples

with an increasing degree of bromination (x = 0.1, 0.3 and 0.5) were prepared following

the procedure described in Section 2.1. The 13C NMR chemical shifts of pure

polystyrene and partially brominated polystyrenes are shown in Figures 2-9(a-d).

Figure 2-9a shows 13C NMR (CDCl3) δ 145 (C1, carbon atom number 1), δ 128

(C2,3,5,6), δ 125 (C4), δ 42-46 (C7), δ 40 (C8). A peak around δ 119.2 (C4-Br) starts to

emerge with x = 0.1 (Fig. 2-9b). With a higher degree of bromination the intensity of

this peak increases (Fig. 2-9c,d). These results are in agreement with the findings of

Farrall and Fréchet24 in their study of halogenated polystyrenes. The samples were

analysed by Prof. Andrew Whittaker from CMR at the University of Queensland. The

13C solution-state NMR spectra were acquired on a Bruker DRX500 spectrometer,

operating at 500.13 MHz and 125.77 MHz at 1H and 13C, respectively. Samples were

dissolved in deuterated chloroform to a concentration of 5-10 wt.%. Spectra were

acquired using inverse gated decoupling, with a sweep width of 27.8 kHz and

acquisition size of 64k to make a total acquisition time of 1.18 s. The 90o pulse time was

9.2 µs. The recycle delay was 15 s, and for each sample 2048 scans were co-added to

improve the signal-to-noise ratio.

Figure 2-9a: 13C-NMR spectrum of pure PS-SiClMe2.

HC C

H21

2

3

45

6

78

x

28

X = 0.1

X = 0.3

Figure 2-9b: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2.

Figure 2-9c: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2.

HC C

H2

Br

HC C

H21

2

3

45

6

78

x y

29

X = 0.5

Figure 2-9d: 13C-NMR spectrum of partially brominated PS(Br)-SiClMe2.

The relaxation time of polymer chains in NMR analysis takes much longer than

for typical low molecular mass organic molecules, and this affects the final resolution.

In this study, a reasonable resolution was required in order to compare the various

degrees of bromination, therefore, short PS chains [MW=8000] were used.

In the following study of the conversion of PS(Br) to PS(NH2), 1H-NMR

spectroscopy was used to detect the presence of the amino groups. This was a

qualitative rather than quantitative study, therefore longer PS chains [MW=200,000]

were used. The use of the long chains has a great advantage in relation to the isolation

(precipitation) of the final product. Precipitation of short polymer chains is more

challenging, usually requiring special techniques such as decanting or centrifuging.

30

HC C

H2

Br NMe3Si

NH2

(I) (II) (III)

SiMe3

(IV)

n1

2

3

4

6

5

7 8

n n n

9

9

2.3.4 Monitoring Conversion of PS(Br) to PS(NH2) by 1H-NMR Spectroscopy

In the following experiment, the process of converting polystyrene to

aminostyrene, described in Section 2.1 was monitored by 1H-NMR spectroscopy. The

schematic of the conversion process is shown in Figure 2-10. The starting material was

a monodisperse polystyrene standard of MW = 200K, obtained from Polysciences Inc.

Figure 2-10: Conversion of PS to PS(NH2).

The 1H-NMR spectrum of pure PS (MW=200,000) is shown in Figure 2-11a.

Figure 2-11b shows the change in a shape of an aromatic polystyrene peak (δ = 6.05–

6.65 ppm) due to bromination in the para position (x = 0.1). In Figure 2-11c, the

broadening of the TMS peak (δ = 0 ppm) is due to the presence of the trimethylsilyl

amino group along the PS chain. Also, there is a new peak around δ = 3.4 ppm which

would suggest a partial conversion of the bis(trimethylsilyl) amine group (-N(SiMe3)2)

to the amino group during the precipitation of the polymer into methanol. The full

conversion of PS-N(SiMe3)2 to PS(NH2) is demonstrated in Figure 2-11d. The spectrum

shows the disappearance of the trimethylsilyl group (δ = 0 ppm) and the presence of the

amino group (δ = 3.4 ppm). The final PS(NH2) was only partially soluble when re-

dissolved in THF for further precipitation, but it dissolved completely in DMF. There

was no difference observed in the 1H-NMR spectra of these separate fractions.

31

Figures 2-12a,b show the effect of D2O addition, and confirm the presence of

the amino group. A few drops of D2O were added to an NMR tube containing a sample

of PS(NH2) dissolved in CDCl3, and the solution was shaken. In the presence of D2O,

the hydrogen atoms are exchanged for deuterium atoms, and as a result, the –NH2 peak

of δ = 3.4 ppm disappears from the 1H-NMR spectrum. This is an indirect method used

to confirm the presence of amino groups in the organic molecules25. The 1H-NMR

spectra were collected using Varian Mercury VX 300 MHz NMR spectrometer.

Figure 2-11a: 1H-NMR spectrum of PS (MW=200K).

Figure 2-11b: 1H-NMR spectrum of PS(Br)x (x = 0.1).

32

Figure 2-11c: 1H-NMR spectrum of PS-N(Si(Me3)2).

Figure 2-11d: 1H-NMR spectrum of PS(NH2).

33

Figure 2-12a: 1H-NMR spectrum of PS(NH2).

Figure 2-12b: 1H-NMR spectrum of PS(NH2) with D2O added.

34

Thus the conversion of polystyrene to PS(NH2) was confirmed by NMR

spectroscopy techniques. The 13C NMR and 1H NMR spectra show each step of the

conversion process. For polymers, the T1 relaxation time is very long compared to that

for smaller organic molecules. Therefore, a dead time set for each analysis needs to be

long in order to obtain quantitative analysis with a sufficient resolution. In the present

case, only qualitative analyses were required. The 13C NMR spectra of brominated

polystyrenes show a clear trend of increasing intensity of the 119.2 ppm peak with an

increasing degree of bromination along the polystyrene chain. The position of this peak

is in agreement with the findings of Farrall and Fréchet24 in their work on halogenated

polystyrenes. The authors showed that the peak δ = 119.8 ppm corresponds to the

hydrogen atom in a para position. Therefore the presence of this peak is a confirmation

that the bromination occurs in para position. The conversion of PS(Brx) to PS-

N(SiMe3)2), the disappearance of bis(trimethylsilyl) group in the next step, and the final

formation of the amino group are evident in the 1H NMR spectra.

35

2.4 References

1 C. J. Hawker, Acc. Chem. Res. 30, 373-382 (1997). 2 C. J. Hawker, G. G. Barclay, A. Orellana, J. Dao, and W. Devonport,

Macromolecules 29, 5245 (1996). 3 G. G. Barclay, C. J. Hawker, H. Ito, A. Orellana, P. R. L. Malenfant, and R. F.

Sinta, Macromolecules 31, 1024-1031 (1998). 4 M. Husseman, E. E. Malmström, M. McNamara, M. Mate, D. Mecerreyes, D. G.

Benoit, J. L. Hedrick, P. Mansky, E. Huang, T. P. Russell, and C. J. Hawker,

Macromolecules 32, 1424-1431 (1999). 5 E. E. Malmström and C. J. Hawker, Macromol. Chem. Phys. 199, 923-935

(1998). 6 K. Matyjaszewski and J. Xia, Chem. Rev. 101, 2921-2990 (2001). 7 K. Matyjaszewski, B. Gobelt, H. Paik, and C. P. Horwitz, Macromolecules 34,

430-440 (2001). 8 S. Perrier, D. Berthier, I. Willoughby, D. Batt-Coutrott, and D. M. Haddleton,

Macromolecules 35 (2002). 9 X. Zhang, J. Xia, and K. Matyjaszewski, Macromolecules 31, 5167 (1998). 10 X. Zhang and K. Matyjaszewski, Macromolecules 32, 1763 (1999). 11 V. L. Covolan, G. Ruggeri, and E. Chiellini, J. Poly. Sci., Part A: Poly. Chem.

38, 2918 (2000). 12 R. P. Kambour and J. T. Bendler, Macromolecules 19, 2679-2682 (1986). 13 R. Oslanec, Thesis, UPENN, 1997. 14 J. Genzer, Thesis, UPENN, 1996. 15 J. Bartulin, G. Cardenas, and H. Maturana, Hydrometallurgy 8, 137 (1982). 16 S. Lee, M. Jorgensen, and J. F. Hartwig, Organic Letters 3, 2729-2732 (2001). 17 R. F. Heck and J. P. Nolley, J. Org. Chem. 37, 2320 (1972). 18 V. C. Gibson and W. Reed, in European Patent Application (UK, 1998), p. 1-9. 19 K. Suzuki, K. Yamaguchi, A. Hirao, and S. Nakahama, Macromolecules 22,

2607-2611 (1989). 20 J. L. Viovy and L. Lesec, Advances in Polymer Science, Vol. 114 (Springer-

Verlag Berlin Heidelberg, Berlin, 1994). 21 E. Schroder, G. Muller, and K.-F. Arndt, Polymer characterization (Hanser

Publishers, Munich, 1989).

36

22 R. M. Silverstein, G. C. Bassler, and T. C. Morril, in Spectrometric

Identification of Organic Compounds (John Wiley & Sons, Inc., New York,

1991). 23 J. E. Mark, Physical Properties of Polymers Handbook (AIP Press, New York,

1996). 24 J. M. Farral and J. M. J. Frechet, Macromolecules 12, 426-428 (1979). 25 S. Pyne.

CHAPTER 3

Polymer Chains

as Coupling Agents

39

3.1 Coupling Agents

Coupling agents are the adhesion promoters used to improve the adhesion between

organic polymers and inorganic surfaces, and to maintain this improved adhesion for a

period of time in various environmental conditions, including moisture or water and heat.

Glass fibre–resin composites are the classical examples (Fig. 3-1) of the use of such

materials. Coupling agents are often used as primers, applied as 0.1 – 10 µm thick layers

during different stages of the pretreatments to improve the performance of the bonded

components. The main roles of the primer are to establish a strong initial bond, and to

increase the service life of the adhesive joint by protecting the substrate from hydration and

corrosion. The result is a strong and moisture resistant interfacial bond. Coupling agents

can be also used as a surface finish (coating) or a surface modifier. The thickness of these

layers varies from one to several monolayers. They are also used as an additive.

Organosilanes, organozirconates, organozirco-aluminates and chrome complexes are some

examples of coupling agents used in adhesives technology1. Zirconate-based primers are

used for pretreatment of polyolefin films prior to printing. Zircoaluminates are applied as

primers for coating fillers in resins and polymers. Chrome complexes are used for example

for increasing the strength and durability of aluminium/polyethylene interfaces.

3.1.1 Commercially-available Coupling Agents

The most commonly used coupling agents are organosilanes. They are primarily

used to promote adhesion between metals, oxides, glass and polymers (resins, coatings,

adhesives etc.). Organosilanes are hybrids of silica and organic polymers, making them

ideal candidates for the bonding together of organic resins and mineral surfaces. Their

history goes back to the 1940’s when they were developed to pretreat glass fibres for fibre

composites, to increase the stability of the glass/polymer interface and to improve the water

resistance of these composite materials1. The general structure can be written as RnSiX(4-n)2.

R is a nonhydrolysable organofunctional group that reacts with organic resins and

40

polymers. X is a hydrolysable group which reacts with an inorganic substrate. Some of the

widely used silanes are amino-functional, epoxy-functional, mercapto-functional, carboxyl-

functional, hydroxyl-functional and vinyl-functional silanes. Epoxy-functional and

carboxyl-functional silanes are excellent coupling agents for thermosets like epoxies,

phenolics, melamines and urethanes3.

Figure 3-1: Schematics of coupling mechanism using coupling agents.

Commercially available silanes can be classified according to the type (-chloro, -

methoxy, -ethoxy) and functionality (mono-, di-, tri-) of the hydrolysable group X (Fig. 3-

2), non-hydrolysable group R (amino-, epoxy-, vinyl-, chloro) (Table III-1) and the chain

length (Fig. 3-3). The choice of silane depends on the particular application. The reactivity

of the hydrolysable group decreases in the order: chlorosilane > methoxysilane >

ethoxysilane. Monofuctional silanes yield a single monolayer, but the layer tends to lack

long term hydrolytic stability. In contrast, trifunctional silanes have high hydrolytic

stability but have a tendency to form multilayers. Silanes with two hydrolysable groups

form more flexible interfaces than trifunctional silanes. The chain length has a significant

effect on the way the silane molecules self-organise on the substrate. The longer the chain

the greater the degree of self-assembly or orientation due to the van der Waals forces acting

between the neighbouring chains. Park and Kim4 investigated the surface energies of four

Silicon oxide

EPOXY RESIN

41

different n-alkyl monochlorosilanes: trimethyl chlorosilane (C1), n-butyldimethyl

chlorosilane (C4), n-octyldimethyl chlorosilane (C8) and n-octadecyldimethyl chlorosilane

(C18). By using the dynamic contact angle measurement technique, they evaluated the

surface energies of these silanes deposited on glass surfaces. They also found that C1 and

C18 formed well organised layers whereas C4 and C8 formed more disordered layers. As

the alkyl chain length decreased, the surface molecules became more liquid-like with lower

surface density and coverage.

Figure 3-2: Functionalities of silane hydrolysable groups.

ClSi

Cl

Cl

H3COSi

OCH3

OCH3

EtOSi

OEt

OEt

ClSi

CH3

Cl

H3CSi

Cl

CH3

-trichlorosilane -trimethoxysilane -triethoxysilane

-dichlorosilane -monochlorosilane

42

Table III-1: Examples of silanes with different non-hydrolysable groups.

Cl(H3CO)3SiH2CH2C

H2C CHSi(OC2H5)3

H2NCH2CH2CH2Si(OC2H5)3

CH2O(CH2)3Si(OCH3)3

O

CH3CH2CH2CH2SiCl3

3-Aminopropyltriethoxysilane

3-Glycidoxypropyltrimethoxysilane

Vinyltriethoxysilane

4-Chlorophenylethyltrimethoxysilane

Butyltrichlorosilane

Figure 3-3: Examples of short and long chain silane molecules.

Organosilanes can be applied from solution (aqueous or non-aqueous), from vapour

phase or by plasma polymerisation. More details about these methods, including their

advantages and disadvantages can be found elsewhere5-10.

ClSi

Cl

CH2

Cl

CH3

ClSi

Cl

Cl

CH2

CH2

CH3

(OTS)

3

16

Butyltrichlorosilane Octadecyltrichlorosilane(n-BTMS)

43

3.1.2 Mechanism of coupling through silane coupling agents

Several theories have been advanced to explain the mechanism of adhesion between

organic polymers and hydrophilic inorganic surfaces3. They include Chemical Bonding,

Surface Wetting, Deformable Layers, Restrained Layer, and Reversible Hydrolysis, each of

which is briefly described below:

The Chemical Bonding Theory

As noted above, coupling agents contain hydrolysable end-functional groups (X)

that can react with a substrate. These are typically alkoxy (MeOSi≡), acyloxy (AcOSi≡),

amino (Me2NSi≡) or chloro (Cl3Si≡) groups. The surfaces of glass or silicon wafers with a

native silicon oxide layer have silanol groups (Si-OH) which are available for covalent

bonding with the end-functional groups of coupling agents.

According to Arkles2, this reaction involves 4 steps. Firstly, hydrolysis of X groups

attached to Si (Fig. 3-4). Water needed for hydrolysis is either added or is present on the

substrate surface. The reactive silanol groups are formed which leads to condensation of

the hydrolysed molecules into oligomers (Fig. 3-5). These oligomers bond to surface

silanol groups via hydrogen bonding (Fig. 3-6a). In the last step, which occurs during the

drying or curing process, covalent bonding with the substrate is achieved (Fig. 3-6b). If

the substrate is silica or glass, siloxane interfacial bonds (Si-O-Si) are formed. However, if

the substrate is metallic, the interfacial Si-O-metal bonds are susceptible to hydrolysis11.

The other end of the coupling agent has at least one different functional group that

reacts with a laminating resin during the curing process. This way, the coupling agent forms

a chemical bridge between two materials creating a strong interfacial bond (Fig. 3-1).

44

Figure 3-4: Hydrolysis of a silane coupling agent with three methoxy end-functional

groups.

Figure 3-5: Formation of oligomers and their condensation to the substrate.

Figure 3-6: a) H-bonding to the substrate, b) consequent covalent bond formation upon

curing.

MeOSi

OMe

R

OMeH2O

HOSi

OH

R

OH

HOSi

OH

R

OH

HOSi

OH

R

OH

HOSi

OH

R

OH

- 2H2OHO Si O

R

OH

Si O

R

OH

Si

R

OH

OH

OH OH OH+

HO Si O

R

O

Si O

R

O

Si

R

OH

OH H H H H H

O O O

HO Si O

R

O

Si O

R

O

Si

R

OH

O

- 2H2O

H

H O

b

a

45

Surface Wetting Theory

According to this theory, adhesion between the lower-energy organic resins and the

high-energy mineral surfaces is strong as the system has a tendency to achieve as low a

surface energy as possible. Complete wetting is necessary to achieve this interfacial

strength. Low viscosity silane coupling agents can be applied to assist the wetting of the

substrate. It was found however, that once the silanes bond to the substrate it is their

reactivity rather than their surface energy that has the greatest effect on the final interfacial

strength between the resin and the surface12,13. In other words, the reactivity of the silane

coupling agent with the polymer resin has greater significance than the polarity or

wettability of the treated mineral surface3.

Deformable Layer Theory

Due to a difference in thermal coefficients between the resin and the mineral

substrate, the resulting differential thermal shrinkage causes stress build-up along the

interface upon cooling. According to the deformable layer theory, when a silane coupling

agents is applied, a flexible interfacial layer is formed which may relieve some of this stress

without bond rupture3. The maximum toughness of such composites is achieved by this

deformable layer at the interface.

Restrained Layer Theory

This theory supposes that an interfacial region is formed with a modulus value

between that of the resin and the substrate14. It was suggested that when the silane coupling

agent is applied it tightens up the polymeric resin at the interface, which helps to achieve

maximum bonding strength and resistance to hydrolytic debonding. The maximum

chemical resistance of a composite is obtained with a restrained layer at the interface.

Reversible Hydrolysis Mechanism

This theory suggests that the bonds of the coupling agents to the substrate break and

reform reversibly in order to release the interfacial stresses15.

46

3.1.3 Nature of Bonding with Silane Coupling Agents

The silane coupling agent forms oxane bonds M-O-Si with various mineral surfaces

(M) such as oxides of silicon, aluminium, zirconium, tin, titanium, nickel, iron, boron and

carbon2. There is no obvious reason to expect a great improvement in water resistance

because some of these oxane bonds, as in the case of Si, Fe and Al, are hydrolysable3.

However, it was experimentally shown that stable bonds are formed with the oxides of Si,

Al, Zr, Sn, Ti, and Ni. Less stable bonds were formed with the oxides of B, Fe and C.

Alkali metal oxides and carbonates do not form stable oxane bonds2. More information on

bonding of the silane coupling agents to different surfaces can be found elsewhere16-20.

From a chemical point of view, the commercially available γ-

Aminopropyltriethoxysilane (γ-APS) would be a logical candidate for the bonding of a

silicon substrate and an epoxy resin because of its amino functionality. However, the nature

of aminosilanes is rather complicated. It has been shown that there is a strong hydrogen

bonding between silanol groups Si-OH of the substrate and the amine groups21 (Fig. 3-7).

Kanan et. al.22 showed, using infrared spectroscopy, that each molecule of (3-

aminopropyl)dimethylethoxysilane deposited on the silica powders from the vapour phase

formed surface bonds with two silanol sites. One Si-O-Si bond was formed through the

dimethylethoxysilane end-group, the second was a H-bond with the amine. As a result,

there were no free amino groups dangling from the silica surface. In this study, the fact that

not all amino groups might be available for chemical reaction with the epoxy resin would

result in incomplete information about the reactivity of the modified silicon surfaces.

Figure 3-7: Bonding of γ-APS to silicon substrate.

SiO

H2C

SiOO O O

H2C CH2

CH2

NH2

CH2

H2C

H2N

Si Si Si SiOH

OHSi

Sisilica

47

3.1.4 Self-Assembled Monolayers

There are two mainstream methods for the deposition of the organic molecules onto

a solid substrate in order to form organised monolayers: the Langmuir-Blodgett technique,

and the technique of self-assembled monolayers (SAM). The Langmuir-Blodgett technique

is a well established method for the preparation of physisorbed monolayers on solid

substrates. These monolayers are linked to the surface via physical bonds, and therefore are

not very stable when exposed to certain chemical environments. Recently this technique

was used by Devaux et al23,24 to form an initiator monolayer chemisorbed on the surface as

a precursor for TEMPO-mediated polymerisation of polystyrene brushes on an AFM tip24.

Sagiv25 developed a technique for preparation of SAM on planar surfaces by

adsorption of amphipathic molecules26. These monolayers are stronger and more

chemically resistant than Langmuir-Blodgett monolayers because they are attached to the

substrate via a chemical bond. As an example, the SAM were constructed on gold surfaces

via a thiol group27 and on silica surfaces via a trichlorosilane group28. The most typical

substrate for studying SAM is a silicon wafer with a native silicon oxide layer. The typical

thickness of this layer is around 18Å, and the surface has smoothness on an atomic level.

This enables surface property modification using SAM without any roughness effects.

Trichlorosilane coupling agents can be used on any silica-type surfaces such as glass and

metallic oxides. The widely assumed mechanism of the silanation reactions is shown in

Figure 3-8.

The silica surfaces are hydrated, and the silanol groups are available for chemical

reaction with the trichlorosilane functional groups. The suggested steps of the silanation

reaction are 7:

1. Physisorption - the polar trichlorosilane molecules are strongly attracted to the

hydrolysed silica surface.

2. Hydrolysis – at the surface, the silane groups are hydrolysed and they form

hydrogen bonds to the silanol groups on the silica substrate.

48

3. Condensation – water is eliminated and each chain is anchored to the substrate and

to the adjacent chains, creating a network.

Figure 3-8: Schematics of the silanation reaction7.

49

Such layers can be very stable and well-oriented. There are several important

factors influencing the final quality of the SAM such as temperature that affects the rate of

reaction as well as the solubility parameter of the trichloromolecules in a particular solvent.

Chain length, and a presence of an end group, have a profound effect on the surface quality

and the final composition of the resulting monolayer.

Sagiv suggests, in his paper on organised monolayers prepared by adsorption of

octadecyltrichlorosilane (OTS)25, that it is very unlikely that every OTS molecule would

form covalent bonds both with the substrate and the adjacent OTS molecules within one

monolayer. The more probable scenario is that some OTS molecules are covalently bonded

to the substrate and form H-bonds to the neighbouring molecules. The resulting monolayer

would be a combination of both chemically and physically adsorbed molecules, creating a

stable 2-D network of closely packed molecules. Therefore in the case of OTS, such

monolayers can be prepared on many types of polar surfaces even if density and

distribution of the surface OH groups vary. To prepare well-defined planar monolayers, it is

important to use anhydrous solutions. If there is more water present than just a surface

monolayer, the trifunctional silanes tend to condense in the bulk and form thick polymeric

films on the surface.

3.1.5 Mixed Monolayers

The studies reported in this thesis involved the adsorption of mixed monolayers

onto silicon wafers. In his work on adsorption of mixed monolayers, Sagiv25 describes the

properties of SAM prepared by deposition of long-chain saturated fatty acids, long-chain

substituted cyanine dyes and OTS molecules mixed at different combinations and ratios.

Cleaned glass slides were dipped in the solutions of these molecules in a mixture of

chloroform, tetrachloromethane and n-hexane. Deposition time ranged from a few seconds

to 10 minutes, at which point the layers were saturated. The saturation was confirmed by

absorption spectroscopy. The monolayer composition vs. solution composition was also

studied in the same work. By plotting the molar ratio of OTS/dye in a monolayer (RM) vs.

50

the molar ratio of OTS/dye in a solution (RS), Sagiv showed that RM values are

substantially smaller than the corresponding RS. For example, for RS = 50 the

corresponding value of RM = 16. However, when the same solution of OTS-dye was

applied to an aluminium substrate, RM was found to be larger than the corresponding RS 29.

Regarding the reproducibility of the resulting layers, the author devoted great care to the

reproducible cleaning of the glass surfaces. However, there was still a large variation in the

RM values reported even with glass slides from the same cleaning batch. This is a relevant

issue because the same was found in studies described in this thesis of mixed monolayers

deposited on silicon wafers. Sagiv concluded that the deposition of mixed monolayers is a

very complex and dynamic process. The final composition and properties of the final

monolayer depend on the interactions of the molecules in solution, and on the interactions

of the adsorbed molecules with a particular substrate. Although it is possible to prepare

mixed monolayers in a one step procedure, not any arbitrary monolayer composition might

be achievable. As the monolayers consist of both physisorbed and chemisorbed molecules,

their final composition is usually different from the solution concentration, and the

composition is a specific function of the particular solute-substrate system and therefore

difficult to predict.

51

3.2 Adsorption of Polymer Chains

In a dilute solution, the polymer chains are swollen and have a conformation of

random coils. As the chains adsorb at a surface, their conformation changes in order to

minimise the free energy of the surface. The excluded volume interactions cause the

random coils to expand. Jenkel and Rumbach30 first proposed three types of chain

conformations: trains (all segments are in contact with the surface), loops (no contact with

the substrate) and tails (free chain ends) (Fig. 3-9). Neutron scattering and neutron

reflectivity techniques have been used to determine the profiles of physisorbed

homopolymers, block copolymers and end-functionalised polymer chains. If there is a low

concentration of adsorbed polymers, the chains flatten in order to decrease the surface

energy. As the coverage and molar mass of chains increases, the length of loops and tails

increases. A structure with a large number of tails with approximately equal length is called

a polymer brush. The brush can be both physisorbed or chemisorbed, depending on the

nature of the functional groups.

Figure 3-9: Conformations of adsorbed polymer chains.

When polymer chains are deposited onto a solid impenetrable substrate, the polymer

chains undergo three different regimes in a solution. First, when the polymer chains are a

large distance from the surface in a very dilute solution, the chains are swollen, it is

assumed they are not interacting, and the distance between the molecules is greater than the

Tail

Loop

Train

Loop

Tail

52

Radius of gyration Rg of a single molecule. This concept is described in Section 3.3. Jones

and Richards found the concentration gradient to be non-uniform in this regime31. As the

chains approach the surface, they become less swollen and begin to overlap. This is the

regime known as a semi-dilute solution, where the chains interact strongly with each other

and the concentration fluctuations are much smaller. When the molecules adsorb onto the

surface, they are in the regime of a high concentration solution where there are no

fluctuations in the spatial concentration31.

A quantity characterising the adsorption process is called the surface excess z*.

(1)

where φ(z) is the volume fraction of polymers as a function of distance from the substrate,

and φ(∞) is the bulk volume fraction of polymer. In dilute solution, z* corresponds to the

adsorbed amount of polymer. A plot of the equilibrium amount of adsorbed polymer (z*)

vs. bulk solution concentration (φ(∞)) is called an adsorption isotherm. A typical example

of a high affinity adsorption isotherm is shown in Figure 3-10.

Figure 3-10: General adsorption isotherm of a polymer with a high affinity for a substrate.

Polymer segment/solvent molecule, polymer segment/surface and solvent/surface

interactions determine the amount of polymer adsorbed at the surface. In a good solvent the

total amount of polymer adsorbed increases with molar mass (M) for polymers of low M.

z*

φ(∞)

∫∞

∞−=0

* )()(( dzzz φφ

53

For polymers with high M the adsorbed amount reaches a plateau. There is no plateau in

adsorption from theta solvents and the amount of polymers increases continuously with M.

As a result, the adsorbed amounts in theta solvents are usually higher than in good

solvents32.

As noted above, the adsorption of polymer chains on a solid impenetrable surface involves

the loss in conformational and translational entropies due to the localisation of the chains

on the surface. The process might also involve a displacement of small molecules (solvent

molecules for example) from the surface. According to polymer adsorption theory32, a

polymer segment will adsorb if its adsorption (interaction) energy ( au2 ) is more negative

than that of the solvent molecule ( au1 ). A dimensionless interaction energy parameter χs for

the adsorption of polymer from solvent is defined as:

(2)

where au1 is the adsorption energy of solvent and au2 is the adsorption energy of a polymer

segment. The polymer adsorbs if χs is positive. There is a critical value of χs that a long

polymer chain has to exceed in order to adsorb. Below this critical value the entropy pulls

polymer chains away from the surface. This is due to the conformational and translational

entropy loss caused by a localisation of the polymer segment at the solid surface.

3.2.1 Competitive Adsorption of Polymer Chains

When more than one polymeric species competes for surface sites, this is referred to

as competitive adsorption. There are two main reasons for different polymers to compete:

1. Entropy effect: simultaneous adsorption of chains of broad range of molecular mass

2. Enthalpy effect: polymer chains contain segments of different adsorption energies

kTuu aa /)( 21s −=χ

54

In the first case, it has been found33 that longer polymer chains adsorb

preferentially, displacing shorter ones from the surface. This displacement is due to the

increased entropy by exchanging one long chain for a number of short chains. Hence free

energy is reduced, which thus favours adsorption of longer chains. This entropy tendency to

adsorb the larger chains is valid regardless of the polymer shape (linear, branched, flexible

etc). As a result, when polymer samples of high polydispersity are used the high molar

mass species adsorb selectively. In the second case, preferential adsorption occurs when

one of the polymers has a higher surface affinity, and therefore a more negative segmental

adsorption energy parameter. The au2 of a particular polymer with respect to the solvent can

be calculated using the self-consistent-field method (SCF). Enthalpy also acts when the

large amount of small molecules (monomers, solvent molecules etc.) are attracted to the

surface more strongly than the polymer segments which causes polymer desorption30. In

both cases, there might be a dynamic process of adsorption and desorption of some

molecules until the equilibrium is reached. The rate of exchange depends on the

temperature, molar mass and structure of adsorbing and desorbing chains 33. Adsorption-

desorption kinetics can be thought of as occurring in four steps: 1. diffusion to a surface 2.

adsorption onto a surface 3. surface rearrangements (conformational reorganisation) and 4.

displacement of polymer chains by other species. The first two steps are rapid whereas the

third and fourth steps are slow and history dependent33.

The field of competitive adsorption of various polymer species was extensively

researched by Frantz and Granick33-35. In studies of the enthalpic effects in competitive

adsorption of hydrogenated polystyrene (hPS), deuterated polystyrene (dPS) and carboxylic

acid terminated polystyrene (hPS-COOH) onto silicon oxide34, the authors used infrared

spectroscopy to determine the surface excess of these polymers. It was found that the

chemical structure of the polymer chains can outweigh the entropic tendency towards a

preferential adsorption of longer chains. In a non-competitive adsorption study of PS of

various M it was found that for fractions with M smaller then 115 000 g/mol the adsorption

was rapid, reaching a plateau within 10 minutes. Larger molecules reached equilibrium

within 2 hours. In a competitive adsorption experiment of a 1 : 1 mixture (by weight) of

hPS (12K) and dPS (12.5K), dPS displaced hPS almost completely within 15 minutes.

55

When higher molecular masses (> 500K) were used, hPS was not detectable on the surface.

However, when the molecular mass of hPS (12K) was larger than that of dPS (6K), the

protonated hPS adsorbed selectively. For polymers with higher M the equilibration time

was considerably longer than for samples of low M, usually around 2 hours. FTIR-ATR

measurements suggest that initially the species of low M occupy the surface. Then, as the

adsorption proceeds, the higher M chains start to displace the lower M chains until all or

almost all of the surface area is occupied by high M species. The authors also explored the

effect of polar chain end-functionality on kinetics of adsorption. In a competitive adsorption

using a 1:1 mixture of hPS-COOH (7K) and dPS (6K), carboxyl terminated chains

segregated to the silicon oxide surface almost completely, with almost no trace of dPS.

Even when the M of hPS-COOH was 2x higher than M of dPS, hPS-COOH still segregated

to the surface almost completely. Only when the mismatch between M of dPS (550K) and

M of hPS-COOH (7K) was very large, dPS adsorbed preferentially but the surface excess

of hPS-COOH was still significant, at 25% of the surface excess measured during a non-

competitive adsorption. The magnitude of the differential sticking enthalpy was calculated

to be 0.03 kT for dPS vs. hPS and 6.4 kT for hPS-COOH vs. hPS. These results showed

that even a small difference in chemical structure can have a major effect on surface

segregation (adsorption). In other words, even a small preferential adsorption per polymer

segment causes large preferential adsorption of the whole chain due to enthalpy.

In addition to the adsorption – desorption kinetic studies of various polymer species,

Schneider and Granick35 showed the effect of solvent quality on the interaction energy

parameter χs, defined as the difference in adsorption energy between a polymer segment

and a solvent molecule. In cyclohexane, which is a theta solvent for PS, the difference in χs

for hPS and dPS is equal to 0.03 kT. In CCl4, a good solvent, the difference in χs for hPS

and dPS is much smaller, around 2 x 10-3 kT.

56

3.2.2 Physisorption vs. Chemisorption

When the nature of the interactions between the polymer molecules and the

substrate is physical (H-bonds, Van der Waals, dispersion and dipolar forces), the

adsorption process is called physisorption. Examples of this process are the adsorption of

PS, PS-OH and PS-COOH onto a silicon substrate36. Bond strength of the adsorbed chains

can be expressed in terms of a surface sticking energy per monomer ε with units of [kBT].

In the case of hydrogen bonding, ε ~ 4 kBT which demonstrates that strong physical bonds

are quite common37. However, these bonds are reversible under certain conditions such as

exposure to a good solvent, displacement by other polymers or molecules in a process of

competitive adsorption. They can also be thermally unstable and prone to dewetting which

makes the surface layer non-homogeneous. When a chemical bond is formed with the

substrate, the process is called chemisorption. As shown before, silanes form chemical

bonds with solid surfaces and therefore silane-terminated polymer chains chemisorb38. The

adsorption of silane end-functionalised polymers with other functionality introduced along

the chain, for example bromine, therefore involves both physisorption and chemisorption

processes.

3.3 Tethered Polymer Chains

When free polymer chains are dissolved in a theta solvent, they assume the

Gaussian coil configuration, and their dimensions can be expressed in terms of the radius of

gyration Rg31:

Rg ~ aN 1/2 (3)

where N is a number of the repeat units in the polymer chain and a is a statistical segment

length. When the solvent is good and excluded volume interactions are taken into account,

where the individual polymer segments exclude other segments from the space pervaded by

the polymer segment, then the above expression becomes:

57

Rg ~ aN ν (4)

where ν is an excluded volume exponent. For dilute solution ν = 0.588.

When the polymer chains are tethered by one end to a solid surface, their

dimensions change. If the grafting density σ of polymer chains adsorbed to a solid surface

is low (average distance between grafting points > Rg), the polymers chains are isolated

from their neighbors and they adopt either ‘pancake’ or ‘mushroom’ conformation. The

pancake conformation is expected when the polymer chains have a high affinity for the

surface39. In that case the enthalpic polymer/chain interactions are stronger than the entropy

and the chains lie flat on the surface (Fig. 3-11a). When the polymer’s affinity to the

surface is very low, the polymer chains adopt the mushroom conformation (Fig. 3-11b).

Dimensions of these ‘mushrooms’ are similar to the radius of gyration of the free chains (Rg

~ N 3/5). As the grafting density σ increases, the neighboring chains start to overlap. The

excluded volume interactions cause chains to stretch away from the surface until the

entropy starts to dominate due to the changes in the chain configuration. The resulting

arrangement of polymer chains is called a polymer ‘brush’ (Fig. 3-11c). The equilibrium

thickness h of a polymer brush varies linearly with the degree of polymerisation40:

h ~ Naσ 1/3 (5)

where a is the statistical segment length and σ is the grafting density.

Figure 3-11: Pancake, mushroom and brush conformations.

b) Mushroom c) Brush a) Pancake

58

3.3.1 Grafting to vs. Grafting from Techniques

As mentioned previously, the polymer chains can be attached to a solid surface by

physisorption or chemisorption. Physisorption is a reversible process that involves physical

interactions, Van der Waals and hydrogen bonding, whereas chemisorption involves a

formation of the covalent bonds. There are two approaches to formation of the covalent

bonds: grafting to and grafting from. The grafting to technique involves a reaction of the

end-functionalised polymer chains with a suitable substrate. This method can be relatively

simple, but does not yield high density brush structures because the chains must diffuse

through the already-adsorbed polymer layer to reach the reactive sites on the surface. This

adsorption barrier increases as the thickness of the adsorbed layer increases. As a result, the

grafing to procedure produces a low density brush of low thickness41 as only small amounts

of the polymer can be deposited. For example, Zhao et al42 investigated ultra thin layers of

Si(OH)3-terminated polystyrene prepared by spin coating onto silicon wafers using AFM.

Geoghegan et al43 prepared triethoxysilane-PS brushes to study the kinetics of brush

penetration into a rubber network. Clarke et al38 studied the structure of grafted PS layers in

various polymers matrices. Various aspects of the grafting to mechanism have been

reviewed extensively elsewhere44-48.

Many researchers focused on developing various techniques for forming a brush-

substrate system using the grafting from approach. The main advantage of this technique is

that it is possible to prepare thick brushes with a grafting density much higher than that

possible using the grafting to technique. The basic principle of all these methods is the

deposition of surface initiators onto the surface, and then growing the polymer chains from

these sites using cationic, anionic or “living” free radical polymerisation techniques.

59

3.4 Techniques for Synthesis of Grafted Polymer Chains

Research focusing on tailoring the surface properties of inorganic materials using

chemically attached polymer brushes boomed in the 1990’s and it is still a growing field.

The methods of preparing the polymer brushes of various physical and chemical properties

in situ evolved dramatically. Initial research involved a multiple step process for generating

surface-imobilised initiator monolayers using azo compounds for the free radical

polymerisation of vinyl polymers41. This multiple step process, however, led to low

grafting densities of the surface attached initiator, incomplete conversions and side

reactions. As quantitative analysis of organic monolayers is very challenging41, therefore it

became impossible, without knowing the exact composition of the initiator layers, to

understand the mechanism of the subsequent polymerisation. Also, there was no simple

way to detach the grafted polymer chains from the surface for further analysis. Therefore

another approach was taken.

Prucker and Rühe41 developed a procedure for the attachment of a complete initiator

in a single reaction step, using azo initiators containing a chlorosilane group. These

initiators consist of an anchoring group (mono-, di- or trifunctional chlorosilane) which

links the initiator to the surface (silica gel), the initiator itself, and a cleavable group that

allows detachment of the polymer chains after polymerisation for further analysis. A use of

di- and trichlorosilane anchoring groups, however, resulted in crosslinked and disordered

silane multilayers, and the graft densities varied from experiment to experiment.

Structurally well organised monolayers were achieved by using the monochlorosilane

anchoring group. The grafting densities of this monofunctional initiator were varied by

varying the silane/silica ratio during the deposition procedure. Polystyrene chains were

grown from the surface-imobilised initiator by free radical polymerisation of styrene. The

polystyrene chains were then degrafted by a transesterification reaction carried out with

methanol/p-toluenesulfonic acid in toluene, and analysed by GPC and light scattering

techniques. Silica gel was used in this study because it has a higher surface area than planar

silicon wafer, and therefore a larger amount of material was available for analysis. The

60

major weakness of this method is the thermal instability of azo-based initiators. Another

drawback is that use of a traditional free radical polymerisation technique is suitable for

preparing homopolymer brushes only. It doesn’t allow the formation of polymer brushes

with accurately controlled structure, or copolymer brushes. These disadvantages can be

overcome by using “living” free radical polymerisation techniques as shown below.

3.4.1 Grafting of Polymer Brushes using NMP

The “living” free radical polymerisation techniques (NMP and ATRP), described in

Section 2.1, can be used for growing polymer brushes from solid substrates. First, the

surface active initiators are applied to the surface (Fig. 3-12) and then polymer brushes are

grown from these active sites (Fig. 3-13). Choice of monomers and their ratio to a bulk

initiator determines the final structure and properties of the grafted brush. This method has

been explored by Hawker et. al.49 in their work on the controlled synthesis of polymer

brushes.

Figure 3-12: Chemisorption of surface active initiators to Si substrate.

61

The aim of the present study was to use these brushes as coupling agents between a

cross-linked polymer (epoxy) and silicon substrates. Therefore it was necessary to vary the

concentration of the active sites on the substrate and the area density of polymer chains

grown from the substrate. An in-active initiator was synthesised by adapting the procedure

for the synthesis of an active initiator published elsewhere50. The schematics for both the

active and the in-active initiator synthesis can be found in Appendix E. These active and

inactive initiators were mixed at different ratios and chemisorbed to the silicon wafers from

the anhydrous toluene solutions. The polystyrene based chains were grown under NMP

conditions from the active sites (Fig.3-13).

Figure 3-13: a) deposition of mixed initiators; b) polymer brush.

Si Si Si Si Si

X X X

Active Initiator

Si Si Si Si Si

X X X

Polymer brush

a)

b)

62

3.4.2 Grafting of Polymer Brushes using ATRP

In this work, the approach was to attach to a silicon wafer a molecule with an active

functionality suitable for reaction with an ATRP agent. In this case γ-aminopropyl

trimethoxysilane (γ-APS) was used. The main advantage of this approach is that γ-APS is a

commercially available reagent, widely used in the industry, for example for bonding steel

to epoxy. In order to change the concentration of this active initiator, it was mixed with n-

butyltrimethoxysilane (n-BTMS) at different ratios (Fig.3-14a). This silane has a terminal

methyl group, and therefore is not able to attach itself to the ATRP agent. 2-

Bromoisobutyryl bromide was reacted with the surface bound γ-APS, forming a surface

attached ATRP initiators (Fig. 3-14b) from which the polymer brushes were grown. The

preparative conditions for polymerisation can be found in Table II-1, Section 2.1.1.2.

Figure 3-14: a) Deposition of γ-APS and n-BTMS on Si wafer

SiH3CO

OOCH3

H3C

SiH3CO

OOCH3

H2N

SiH3CO

OOCH3

H3Ca)

SiH3CO

OOCH3

H2N

Br

O

Br

+

SiH3CO

OOCH3

HN O

Brb)

63

b) ATRP agent attachment.

3.4.3 Grafting of Polystyrene Chains and their Functionalisation

Some of these procedures were reported in Section 2.1 where the process of the

functionalisation of free polystyrene was described. In this section, the process of “grafting

to” of the brominated silane, end-functionalised, polystyrene and the consequent conversion

to amino-functional polystyrene is described.

Brominated, silane-terminated, polystyrene [PSBr0.1-SiClMe2 (MW = 8000, x =

0.1)], mixed with PS-SiClMe2 in various ratios, was deposited onto silicon wafers from

toluene solution, forming the mixed monolayer (Fig. 3-15).

Figure 3-15: Chemically adsorbed PSBr and PS chains on Si wafer. The conversion of the bromine to the amino functional group is performed by

adapting a palladium-catalysed method of converting aryl halides to anilines using lithium

bis(trimethylsilyl)amide (LiN(SiMe3)2) (Fig. 3-16). The reaction is catalysed by

bis(dibenzylideneacetone)palladium Pd(dba)2 and tri-t-butylphosphine P(t-Bu)3 (Fig. 3-17)

and is known as the Heck reaction. Details of the preparation can be found in section 3.6.3.

Br

Br

Br

Br

BrBr

X NH2

R NMe3Si SiMe3

LiR+

1. Pd(dba)2/P(t-Bu)3 (1:1)

2. HCl, neutralisation

X = Br or Cl

64

Figure 3-16: Heck reaction using LiN(SiMe3)2.

a) b)

Figure 3-17: Catalysts for Heck reaction: a) Pd(dba)2 , b) P(t-Bu)3.

The conversion procedure was carried out in two steps. In the first step, the bromine

group was converted to the bis(trimethylsilyl)amide protective group (Fig. 3-18). The

second step consists of de-protection of -N(SiMe3)2, and conversion to –NH2 by

acidification and neutralisation (Fig. 3-19).

Figure 3-18: Conversion of the PSBr to the silane protected PS.

Figure 3-19: Conversion of the silane terminated PS to PS-NH2.

O

OPd P

C

C C

CH3 CH3H3C

CH3

CH3CH3

H3CH3C

H3C

NSiMe3Me3Si

Br

Li

NMe3Si SiMe3

nn

+Pd(dba)2/ P(t-Bu)3

NSiMe3Me3Si

NH3+Cl- NH2

NaOH

n

HCl

n n

65

3.5 Analysis of Grafted Surfaces

3.5.1 Ellipsometry

Ellipsometry is an optical technique used to analyse transparent thin films. It has

been used mainly in semiconductor research and fabrication to determine properties of

layer stacks of thin films and the interfaces between the layers. Ellipsometry is also utilised

by researchers from other fields such as biology, medicine and polymer coatings. The main

advantage of ellipsometry is that it is a non-destructive, highly sensitive technique with a

measurement range from monolayers to micrometers. The upper limit of film thickness

amenable to ellipsometric measurement is determined by the film uniformity and

homogeneity. The precision of the measurement is also given by the optical properties of

both substrate and the transparent film. To achieve the best possible resolution (units of

angstroms) the sample should be optically flat, non-scattering, highly reflecting, have

uniform film refractive index (homogeneous and isotropic), large differences between film

and substrate refractive indices, and the film should be transparent.

Figure 3-20: Single-film model. ϕ is the angle of the incident and reflected beams, n and k

are the real and imaginary parts of the refractive index of the substrate and the film, t is the

film thickness.

φ φ

λ

φ2φ2

Substrate: n3, k3

Film: n2, k2, t

Ambient

66

A highly collimated beam of linearly polarised light generated by a laser travels

through the transparent film/films to a substrate where it is reflected and analysed by a

photodetector. The case of a single-film model is shown in Figure 3-20. From the measured

difference between the states of polarisation of the incident and reflected beam, various

properties of the reflecting surface can be computed. Single wavelength ellipsometry can

measure only a film thickness and refractive index whereas Spectroscopic ellipsometers

(multiple wavelength beam) can analyse complex multilayered structures, interface

roughness, inhomogeneous and anisotropic layers.

Figure 3-21: Schematic of the analysing section of the AutoEl-II Ellipsometer.

A Rudolph Research AutoEl-II Ellipsometer was used to measure the average

thickness of the functionalised polystyrene chain deposited on Si wafers. The schematic

setup is shown in Figure 3-21. The AutoEl-II is a nulling type Ellipsometer. It is equipped

with a helium-neon laser generating collimated, monochromatic light at a wavelength 632.8

nm. The incident beam angle was set to 70o. A double-film model was used for the analysis

(Fig. 3-22).

Autocollimator

Eyepiece

Laser head

Polariser

Sample

Analyser

Adjustable stageCompensator slide

67

Figure 3-22: Double-film model. ϕ is the angle of the incident and reflected beams, n and

k are the real and imaginary parts of the refractive index of the substrate and the film, t is

the film thickness.

For a single film model and when the refractive index n of an analysed layer is

known, the equations below can be used to determine the film thickness t:

where: pR and sR = total reflection coefficients

pr12 , sr12 = the Fresnel reflection coefficients for the ambient medium-film

interface parallel and perpendicular to plane of incidence

pr23 , sr23 = the Fresnel reflection coefficients for the film-substrate

interface parallel and perpendicular to plane of incidence

φ φ

λ

φ1φ1

φ2

Substrate: n3, k3

Film: n2, k2, t

Film: n1, k1, t

Ambient

)cos(2

)exp)(exp1()exp1)(exp(exptan

22

22312

22312

22312

22312

φλπβ

ββ

ββ

nt

rrrrrrrr

RR

jssjpp

jssjpp

s

pi

=

+++++==Ψ −−

−−∆

68

In this study, the single film model was used for the thickness calculations. The

silicon oxide layer thickness on the silicon wafers was determined prior the surface

chemical modification. These thickness values were included in a software for calculating

the deposited polymer layer thickness. The refractive index of polystyrene n = 1.592 was

used for both pure and mixed polymer monolayers.

3.5.2 Surface Energy by Contact Angle Measurement

In the field of adhesion and adhesives technology, knowledge of the surface energy

of a given material can be used to predict its surface properties and its interactions with

other materials. Contact angle measurements provide a very sensitive tool for examination

of surface chemistry and surface homogeneity, which are used to monitor surface

treatments and surface cleanliness as well as surface wettability. The main advantages of

this technique are its simplicity, rapidity of analysis and inexpensive equipment.

A surface free energy of solids γ is a measure of the attractive intermolecular forces

between a surface layer and a liquid phase. As a drop of a testing liquid is placed on top of

a solid surface, it either spreads or retracts, depending on the interfacial energy between the

solid and the liquid. The initial spreading coefficient So can be defined in terms of the

difference between the solid surface energy (γsv) and the liquid surface tension (γlv) together

with the interfacial tension (γsl) as51:

So = γsv - γlv - γsl (6)

The liquid spreads spontaneously on the solid when So > 0 (Fig. 3-23a). When So < 0 the

liquid wets the surface only partially (Fig. 3-23b).

69

Figure 3-23: a) So > 0, b) So < 0. Θ is an equilibrium contact angle.

The Young equation relates the equilibrium contact angle Θ to all three components, γsv, γsl

and γlv:

γsv = γsl + γlv cos Θ (7)

When the solid surface energy is higher than the liquid surface tension, then the surface free

energy γsv of the substrate resulting from adsorption of vapour from the liquid may be

considerably lower than the surface free energy in vacuum (γs). In order to balance this

discrepancy, an equilibrium spreading pressure πe is introduced52:

γsv = γs - πe (8)

πe can be neglected in the case where a high surface energy liquid wets a low surface

energy substrate. In the present study, water and diiodomethane were used as the testing

liquids for polymer coated surfaces. These surfaces generally have lower surface energies

and therefore πe was neglected in further calculations.

The intermolecular forces present when two dissimilar materials are brought into contact

are expressed in the Dupré equation as the work of adhesion WA 52:

γsl = γsv + γlv - WA (9)

When the dispersive and polar forces are considered, then according to the theory of

fractional polarity53:

Θγsl γsv

γlv

Θ

γsl γsv

γlv

a) b)

70

γ = γd + γp (10)

where γd and γp are the dispersive and polar components of the surface energy. The polar

component of surface energy includes dipole forces, induction forces and hydrogen

bonding54.

Using a geometric mean assumption, for a liquid drop forming a contact angle (θ) on a

solid surface, the combined Owens-Wendt/Young equations 55,56 state that:

(11)

At least 2 testing liquids of different polarities are used for the contact angle measurement

when this method is used. Dispersive (γlvd) and polar (γlv

p) surface tensions for many testing

liquids are well known. The values for water and diiodomethane (the test liquids used in

this work) are shown in Table III-2.

A Ramé-Hart Model 100 Goniometer system was used in this study to measure static

contact angles using the testing liquids mentioned above.

Table III-2: Surface Tension of Test Liquids

Liquid γld (mJ/m2) γl

p (mJ/m2) γl (mJ/m2)

Water 21.8 51.0 72.8 Diiodomethane 48.3 2.5 50.8

Another method used by researchers to determine the composition of thin polymeric

films is X-ray Photoelectron Spectroscopy (XPS)57. In this study, however, XPS studies did

not provide satisfactory results due to lack of sensitivity.

+=Θ+

dlv

plvP

svdsvd

lv

lv

γγγγ

γγ

2)cos1(

71

3.6 Results and Discussion

3.6.1 NMP Results

The procedures for the synthesis of polystyrene brushes and their derivatives can be

found in work of Hawker et. al.49. Identical conditions were employed in the present work.

The molecular mass of the non-grafted (free) polymer generated during the synthesis due to

the addition of “free” alkoxyamine initiator was analysed by GPC. The results for PS and

PS/PHEMA brushes are shown in Table III-3.

Table III-3: PS and PS/PHEMA brushes

Monomers M1/ M2 M/Init MWcalc MW tbrush [Å] σ [# chains/m2]

Styrene 1/0 500/1 47K 26K 83 2.0x1017

Styrene/HEMA 9/1 200/1 20K 3K 51 8.5x1017

M1/ M2 ………molar ratio of monomers in a reaction solution

M/Init………..molar ratio of all monomers to initiator molecules

MWcalc ………theoretical molecular mass (MW = M x MWaverage + MTEMPO) x conversion

MW………….molar mass determined by GPC

tbrush………….thickness of polymer brush measured by Ellipsometry

σ……………..areal chain density

The areal chain density σ has been calculated from the following equation:

MWNt

AA××== ρσ # , (12)

where A is the sample area (typically 1 cm2)

72

ρ is polymer density (ρPS = 1.05 g/cm3)

t is the measured thickness of the polymer layer

NA is Avogadro’s number (6.023x1023)

MW is the molecular mass of the polymer

# is number of polymer chains

Hawker49 reported that the relationship between brush thickness and molecular mass

of free polymer is almost linear. Therefore it can be assumed that the molecular mass of the

covalently attached polymer chains is consistent with the molecular mass of the free

polymer. In comparison, the chain density σ = 5x1017 chains/m2 of the polystyrene brush

grown from silicon wafer reported by Hawker et. al49 is somewhat higher than the value

found in the present study. The chain density of PS/PHEMA brush was found to be slightly

higher, implying a more dense brush.

Table III-3 shows poor agreement between the calculated MW and the measured

MW for Styrene/HEMA copolymers. The refractive index detector used in this study is

very useful for determining molar mass of homopolymers. Introducing copolymer might

affect an elution volume Ve and therefore resulting MW would deviate more from the

calculated value. In this study, however, there was only 10 mole% of HEMA incorporated

into the copolymer and therefore it is more likely that the large difference between the

calculated and measured MW was caused by incomplete conversion during the

copolymerisation process resulting in copolymers with a short chain length.

The aim of the next experiment was to obtain PS brushes of various densities. The

mixtures with different ratios between the active and in-active initiators were prepared. The

initiators were deposited onto Si wafers from anhydrous toluene solutions and the surfaces

were analysed by ellipsometry and contact angle goniometry. The results are summarised in

Table III-4. The initiator synthesised at the University of Wollongong (UOW) produced

layers of the same thickness as the initiator supplied by IBM Almaden Research Centre but

the surface energy was lower, especially the polar component γinitiator (polar). The polystyrene

brushes were grown by NMP from these surface bound initiators. The chains grown from

73

the UOW prepared initiators produced the thickness and areal density of almost one

magnitude higher than the chains grown from the IBM prepared initiators. The reason for

the differences is not clear.

Table III-4: PS brush on Si wafer prepared by NMP

Active Init

[%]

tinitiator

[Å]

γinitiator(total)

[mJ/m2]

γinitiator (polar)

[mJ/m2]

tPS

[Å]

σ

[# chains/m2]

0 70.6 46.4 5.3 15.4 3.7x1016

10 a 14.2 66.6 25.2 23.8 5.7x1016

20 a 11.0 63.4 26.4 9.3 2.2x1016

100 a 16.6 56.2 22.2 10.8 2.6x1016

100 b 16.2 40.2 10.9 94.8 2.3x1017

a active initiator prepared at IBM Almaden Research Centre b active initiator synthesised at UOW

Even though the brush synthesis was successful, the chain densities were lower than

in the previous experiment and the results show that there is a poor control over the

polymer chain density.

3.6.2 ATRP Results

For grafting of PS, several mixtures of various ratios of γ-APS/n-BTMS in

anhydrous toluene were prepared and deposited onto the silicon wafers at room temperature

overnight. The reaction conditions are summarised in Table III-5. After Soxhlet extraction

to remove any non-reacted polymer, the final polymer layer was analysed by Ellipsometry.

The “free” polymer was analysed by GPC. As shown in Table III-6, the thickness and the

chain density increases with increasing γ-APS concentration. Samples with 100% γ-APS

concentration are an exception.

74

Table III-5: Grafting of PS chains

Sample % APS M:I:C:L Ligand Temp [oC] Solvent

I 10 200:1:1:2.5 Bipy 110 bulk

II 25 200:1:1:2.5 Bipy 110 bulk

III 50 200:1:1:2.5 Bipy 110 bulk

IV 100 200:1:1:2.5 Bipy 110 bulk

IM 10 200:1:1:1 Me6-TREN 110 bulk

IIM 25 200:1:1:1 Me6-TREN 110 bulk

IIIM 50 200:1:1:1 Me6-TREN 110 bulk

IVM 100 200:1:1:1 Me6-TREN 110 bulk

Table III-6: PS brush on Si wafer

Sample Thickness [Å] MW PD σ [# chains/m2]

I 77 8609 1.18 5.7x1016

II 261 9850 1.24 1.7x1017

III 701 9948 1.18 4.5x1017

IV 289 9753 1.22 1.8x1017

IM 101 21306 1.14 3.0x1016

IIM 251 26260 1.16 6.0x1016

IIIM 561 27759 1.13 1.3x1017

IVM 450 28458 1.16 1.0x1017

75

Because of inconsistency in PS growth, the coupling agent itself was investigated.

From the thickness measurements of the deposited γ-APS/n-BTMS layers (Fig. 3-24), it is

clear that for the lower concentrations of γ-APS the increase in the layer thickness is very

low. For the higher concentrations, however, the thickness increases dramatically due to a

thick multilayer formation. This corresponds to the findings of Prucker and Rühe41. The

authors report that the tri-functional silanes exhibit a tendency towards the formation of

surface attached networks. In addition, the amino groups can have strong interactions with

the appropriate sites at the surface. If such layers are used for additional surface reactions,

the structure of the resulting monolayer is difficult to analyse and reproduce41.

Figure 3-24: γ-APS thickness vs. γ-APS concentration deposited on Si wafer.

In the following experiment, poly(2-(trimethyl(silyloxy)ethyl methacrylate)

(PHEMA-TMS) and PS copolymer brushes were prepared in the same way as described

above for pure PS brushes (HEMA-TMS/St ratio was 1/9, Me6-TREN was used as ligand,

reaction was carried out in a bulk, at 110oC for 4 hours). The concentration of γ-APS and

the ATRP reaction time were varied. The final layers were analysed by Ellipsometry. The

results are summarised in Table III-7.

APS/BTMS

0

100

200

300

400

0 20 40 60 80 100

[%] APS

Thi

ckne

ss [Å

]

76

Table III-7: Thickness of PHEMA-TMS/PS brush

γ-APS [%] Time [h] Thickness [Å]

5 4 149

10 4 126

25 4 123

50 4 112

50 2 95

50 4 93

50 6 88

50 8 95

Neither γ-APS concentration or reaction time produced the expected increase in

thickness (and therefore increased chain density). As mentioned previously, it is very

challenging to measure quantitatively the composition of the silane layer. Therefore it is

uncertain if these results are due to the lack of control over the silane deposition or the

ATRP process itself.

77

3.6.3 Bromination of PS-SiClMe2 and Deposition of Mixed PSBr0.1-SiClMe2/PS-

SiClMe2 monolayers

Silane terminated polystyrene (MW=8000) was dissolved in nitrobenzene (7 wt%)

and put into a test tube wrapped in aluminium foil. The calculated amount of bromine was

added and the solution was stirred in dark for 24 hour. The brominated polymer was

precipitated into MeOH. Due to the low molecular mass of the initial polystyrene, it was

not possible to filter the precipitated polymer in the traditional way using filter paper,

therefore the polymer was left to settle at the bottom of a beaker, the solvent was decanted

and the remainder was left to evaporate. The polymer was dried at 75oC under a vacuum

overnight. This step was repeated at least twice.

Silicon wafers were cut into 1.5 x 7.5 cm pieces using a diamond scriber and a

scalpel. Each wafer was cleaned using chloroform followed by piranha (H2SO4/H2O2)

solution in an ultrasonic bath. The wafers were washed with tap water, de-ionised water,

acetone and toluene and then dried with a nitrogen flow. They were then exposed to

ultraviolet ozone cleaning system (UVO) treatment for 1 hour, to remove the last traces of

organic contaminants, and used for the solution depositions immediately.

Various ratios of PSBr(0.1)SiClMe2/PS-SiClMe2 were dissolved in dry toluene to

form 1% solution. Cleaned Si wafers were dipped in the above polymer solutions and

stirred over night. The next day the wafers were baked in the oven at 120oC for several

hours and then cleaned by Soxhlet extraction for at least 12 hours to remove any physically

adsorbed polymer chains from the surfaces. Surface energy (γ) and thickness (t) of

chemically adsorbed mixed polymer monolayers were determined using contact angle and

ellipsometry measurements. The contact angles were measured using water and

diiodomethane.

The brominated polystyrene chains were deposited onto the silicon wafers before

any further chemical modification, in order to “preserve” the reactivity of the polymer’s

78

silane end-functional group and also to prevent the reaction of the amino groups with the

silicon substrate.

Experiment 1

The 1% solutions of PSBr(0.1)SiClMe2/PS-SiClMe2 (1, 2.5, 5 and 10 wt% of

PSBr(0.1)SiClMe2) in toluene were deposited on Si wafers. The ellipsometry and contact

angle measurements (CAM) of the resulting layers are summarised in Table III-8.

Table III-8: Mixed PSBr(0.1)SiClMe2 and PS-SiClMe2 layers (Exp. 1)

PSBr [%] t [Å] σ [# chains/m2] γtot [mJ/m2] γdisp [mJ/m2] γpolar [mJ/m2] 1 9.1 7.2x1016 41.98 35.24 6.73 1 8.6 6.1x1016 41.88 35.05 6.83

2.5 9.6 7.6x1016 41.89 35.54 6.34 2.5 9.0 7.1x1016 41.72 33.81 7.9 5 7.3 5.8x1016 44.52 31.22 13.29 5 7.4 5.9x1016 43.41 30.81 12.6

10 6.4 5.1x1016 42.72 29.92 12.8 10 6.4 5.1x1016 43.38 29.08 14.3

The thickness of the PSBr(0.1)SiClMe2/PS-SiClMe2 layers was in the range 6Å – 10Å, which

corresponds to a chain density ranging from 5.1x1016/m2 to 7.6x1016/m2. These values are

in the same order of magnitude as the polystyrene brushes of similar molecular mass

prepared by ATRP (Section 3.6.2). The total surface energy of these layers was relatively

uniform, but the polar component of the surface energy increased with an increasing

amount of brominated polystyrene in the layer. The value of 14.3 mJ/m2 for 10%

PSBr(0.1)SiClMe2 represents a 100% increase compared to 6.7 mJ/m2 for 1%

PSBr(0.1)SiClMe2.

79

Experiment 2

The previous experiment was repeated for different ratios of PSBr(0.1)SiClMe2/PS-SiClMe2

2, 4, 6, 8, 10, 13, 16, 20 wt% of PSBr(0.1)SiClMe2). Results of the Ellipsometry

measurements and CAM are summarised in Table III-9.


PSBr [%] t [Å] σ [# chains/m2] γtot [mJ/m2] γdisp [mJ/m2] γpolar [mJ/m2]2 21.9 1.7x1017 43.98 42.86 1.12 4 13.3 1.1x1017 43.38 42.3 1.08 6 21.9 1.7x1017 43.69 42.78 0.91 8 23.3 1.8x1017 43.01 42.18 0.83

10 19.7 1.6x1017 42.58 40.71 1.88 13 15.8 1.3x1017 43.02 40.9 2.12 16 21.8 1.7x1017 42.91 41.14 1.77 20 24.4 1.9x1017 42.57 41.24 1.34

In this experiment, the values of the polar surface energy are scattered and lower than in the

previous experiment. The thickness of the PSBr(0.1)SiClMe2/PS-SiClMe2 layers was in the

range 13Å – 25Å, which corresponds to a chain density ranging from 1.1x1017/m2 to

1.9x1017/m2.

80

Experiment 3

1% toluene solutions of 2, 4, 6, 8, 10, 12, 16, 20 wt% of PSBr(0.1)SiClMe2/ PS-SiClMe2

were prepared and deposited on Si wafers. Table III-10 shows the thickness and the areal

chain densities of the resulting layers. The surface energies were not analysed during this

experiment.


PSBr [%] t [Å] σ [# chains/m2] 2 21.0 1.7x1017 4 15.7 1.3x1017 6 15.3 1.2x1017 8 9.6 7.6x1016

10 16.0 1.3x1017 13 18.6 1.5x1017 16 13.3 1.1x1017 20 14.7 1.2x1017

Experiment 4

This experiment was run under the same conditions as Experiment 3. Table III-11 shows

the thickness and surface energies of the resulting layers.


PSBr [%] t [Å] σ [# chains/m2] γtot [mJ/m2] γdisp [mJ/m2] γpolar [mJ/m2] 2 14.7 1.2x1017 44.83 31.56 13.27 4 16.9 1.3x1017 41.8 36.76 5.04 6 13.0 1.0x1017 44.93 32.43 12.5 8 11.3 8.9x1016 43.87 33.47 10.37

10 13.1 1.0x1017 41.58 36.04 5.54 13 11.3 8.9x1016 45.04 32.51 12.53 16 11.2 8.9x1016 45.89 31.77 14.12 20 15.4 1.2x1017 41.37 37.27 4.1

81

The PSBr and PS chains used in this study had the same length, therefore it was not

clear if the thickness of the grafted layers would change with an increasing concentration of

PSBr chains in the grafted layer. The trend of a decreasing thickness with an increasing

PSBr concentration could be related to the attractive forces between the silica surface and

bromine atoms on the PSBr chains. As mentioned earlier, due to these attractions, the

polymer chains have a greater tendency to lie flat on the surface. This flat conformation

prevents other chains from coming into a contact with the surface. As a result of this

behaviour, it would be expected that the thickness and the chain density would decrease

with increasing PSBr concentration. The results, however, do not support this expectation,

and the thickness varied slightly from experiment to experiment showing no trend.

The next step was the conversion of the bromine to the amino- functional group

using Heck’s procedure. The modified Si wafers were put into containers with sealable lids

and placed in a glove box filled with argon. About 30 mL of dry toluene was added,

followed by LiN(SiMe3)2, P(t-Bu)3 and Pd(dba)2. For 1.0 mmol of brominated units,

1.1mmol of LiN(SiMe3)2 and 0.05 - 0.002mmol of P(t-Bu)3 and Pd(dba)2 were used. The

samples were stirred for 24 hours under an inert atmosphere, and then they were washed

with fresh toluene and dried under a nitrogen flow.

To de-protect the silane-protected amino group, the samples were placed into

containers with THF/MeOH (9/1) solution. A few drops of 1M methanolic HCl were added

to the solutions to convert the silane group to the -NH3+Cl- salt. Then the samples were

washed with diethyl ether, dried with N2 flow and placed into fresh THF/MeOH (9/1)

solution. 1M methanolic NaOH was added to convert -NH3+Cl- groups to -NH2 groups. The

samples were washed with water containing triethylamine, dried in the oven at 110oC and

stored in a desiccator under an inert atmosphere. The thickness of the final polymer layers

was measured by ellipsometry.

The thicknesses measured by ellipsometry before and after Heck reaction are

summarised in Tables III-(12-15). With an exception of Experiment 1 (Table III-12), the

82

monolayer thickness was dramatically reduced after the final conversion step, when the

monolayers were exposed to acidic and then basic environments.

Table III-12: Change in thickness of the polymer layer after Heck reaction (Exp 1)

PSBr [%]

Thickness after Soxhlet extract

[Å]

Thickness after HCl and NaOH

[Å]

[%] retained

1 8.9 9.3 98 2.5 9.3 9.3 98 5 7.4 12.5 144

10 6.4 9.3 125

Note: Thicknesses shown in this table are the averages values.


PSBr [%]


[Å]


[Å]

[%] retained

2 21.9 9.4 43 4 13.3 7.6 57 6 21.9 6.8 31 8 23.3 5.8 25

10 19.7 8.8 45 13 15.8 8.6 54 16 21.8 8.4 39 20 24.4 7.2 30


PSBr [%]


[Å]


[Å]

[%] retained

2 21 5.4 26 4 15.7 6.4 41 6 15.3 4.2 28 8 9.6 4.8 50

10 16 3.8 24 12 18.6 3.4 18 16 13.3 3.8 29 20 14.7 3.2 22

83


PSBr [%]


[Å]


[Å]

[%] retained

2 14.7 5.8 40 4 16.9 5.6 33 6 13 7 54 8 11.3 6.4 57

10 13.1 5.8 44 12 11.3 5.8 51 16 11.2 6.4 57 20 15.4 6.2 40

A control test was performed in order to investigate the cause of the polymer

monolayer reduction. Pure PS-SiClMe2 was deposited on a Si wafer from 1% toluene

solution under the same conditions as for the mixed monolayers, as described in the

previous experiments. The sample was subjected to Soxhlet extraction for 3 days to remove

any free polymer chains. The average monolayer thickness, as measured by ellipsometry,

was 40±1Å. Then, the sample was placed in pure toluene and 3 drops of 1M HCl in MeOH

were added while stirring. After the exposure to this acidic environment, the average

thickness was 39±1Å. The sample was put back into the toluene solution and an additional

3 drops of 1M HCl in MeOH were added. After removing the sample, the thickness was

measured again and it was found to be 38±1Å. From these observations, it is clear that the

sample containing only PS-SiClMe2 retained its original monolayer thickness which means

that no chains were detached from the Si surface.

In contrast to the single monolayer, the mixed monolayer thickness was reduced

after the samples were exposed to the acidic conditions. This reduction appeared to be a

completely random process varying from experiment to experiment even though the

experimental conditions were kept the same. The randomness of the washing off of the

polymer chains was apparent within each experiment as well, showing no trend with an

increasing concentration of the amino groups. Possible explanations for the washing off of

84

the polymer chains during this last step of conversion to amino- groups, along with its

implications for the adhesion studies, are discussed in Chapter 5.

These grafted chains were further used for coupling with diglycidylether of

bisphenol A (DGEBA). It is assumed that the amino groups are randomly distributed along

the polystyrene chain. These amino groups are available for the reaction with two

component epoxy resins where one of the components is the amine hardener.

In this study, DGEBA was used as the pre-polymer and either 1,5-diamino-1-

methylpentane (DMP) or Jeffamine D230 were used as curing agents (Chapter 4; Fig.4.4).

Figure 3-25: Bonding of Si to epoxy resin via PS-NH2 chains.

Upon curing, which is a process involving heating, epoxy groups react with amino

groups forming a highly cross-linked network. The unreacted DGEBA/DMP mixture was

used as a glue to bond cured DGEBA/DMP blocks to the polymer modified silicon wafers.

Upon curing, the amino groups of the functionalised polystyrene react with the epoxy

groups in the glue and a permanent bond between the polymer and the epoxy resin is

formed (Fig. 3-25).

O

Si

O

Si

O

Si

NH2

NH2H2N

H2NNH2

H2N HN

H2N NH2

NH2

NH2

H2N

O OHC OH

CH2

DGEBA/DMP

85

3.7 References

1 A. J. Kinloch, Adhesion and Adhesives-Science and Technology (Springer, 1987). 2 B. Arkles; Vol. 7 (Petrarch Systems, Inc., 1977). 3 E. P. Plueddemann, Silane Coupling Agents (Plenum Press, 1982). 4 J. M. Park and J. H. Kim, J. of Colloid & Int. Sci. 168, 103-110 (1994). 5 M. E. McGovern, K. M. R. Kallury, and M. Thompson, Langmuir 10, 3607-3614

(1994). 6 J. B. Brzoska, J. B. Azouz, and F. Rondelez, Langmuir 10, 4367-4373 (1994). 7 P. Silberzan, L. Leger, D. Ausserre, and J. J. Benattar, Langmuir 7, 1647-1651

(1991). 8 C. R. Kessel and S. Granick, Langmuir 7, 532-538 (1991). 9 X. Xiao, J. Hu, D. H. Charych, and M. Salmeron, Langmuir 12, 235-237 (1996). 10 P. K. C., R. W. Johnson, and T. A. Desai, Surf. and Coatings Tech. 154, 253-261

(2002). 11 E. P. Plueddemann, Int. J. Adhes. 1, 305 (1981). 12 W. D. Bascom, J. Colloid Interface Sci. 27, 789 (1968). 13 L. H. Lee, J. Colloid Interface Sci. 27, 751 (1968). 14 C. A. Kumins and J. Roteman, J. Polymer Sci. 1, 527 (1963). 15 A. J. Kinloch, Structural Adhesives: Developments in Resins and Primers (Elsevier

Applied Science Publishers, London). 16 J. S. Quinton and P. C. Dastoor, Appl. Surf. Sci. 152, 131 (1999). 17 M. R. Turner, E. Duguet, and C. Labruere, Surf. and Int. Analysis 25, 917 (1997). 18 R. Hild, C. David, H. U. Muller, B. Vokel, D. R. Kayser, and M. Grunze, Langmuir

14, 342 (1998). 19 A. Kuznetsova, E. A. Wovcho, and J. T. Yates, Langmuir 13, 5322 (1997). 20 B. L. Kropman, D. H. Blank, and H. Rogalla, Langmuir 16, 1469 (2000). 21 H. Ishida, C. Chiang, and J. L. Koenig, Polymer 23, 251-257 (1982). 22 S. M. Kanan, W. T. Y. Tze, and C. P. Tripp, Langmuir 18, 6623-6627 (2002). 23 C. Devaux, J. P. Chapel, E. Beyou, and P. Chaumont, Eur. Phys. J. E 7, 345 (2002). 24 C. Devaux and J. P. Chapel, The Eur. Phys. J. E 10, 77-81 (2003).

86

25 J. Sagiv, J. Am. Chem. Soc. 102, 92-98 (1980). 26 in An amphipathic (or amphiphilic) molecule is a compound having both a

hydrophilic and a hydrophobic end. Some examples of these molecules are soaps,

detergents and some coupling agents. Due to their unique structure, they can

organised themselves either into the structures called micelles or to form SAM as

mentioned above. 27 C. D. Bain and G. M. Whitesides, J. Am. Chem. Soc. 111, 7164 (1989). 28 R. Maoz and J. Sagiv, J. Colloid Interface Sci. 100, 465 (1984). 29 E. E. Polymeropoulos and J. Sagiv, J. Chem. Phys. 69, 1836 (1978). 30 E. Jenkel and B. Rumbach, Electrochem. 55, 612 (1951). 31 R. A. L. Jones and R. W. Richards, Polymers at Surfaces and Interfaces (1999). 32 G. J. Fleer and M. A. Cohen Stuart, Polymers at Interfaces (Chapman & Hall,

London, 1993). 33 P. Frantz and S. Granick, Physical Review Letters 66, 899-902 (1991). 34 P. Frantz, D. C. Leonhardt, and S. Granick, Macromolecules 24, 1868-1875 (1991). 35 H. M. Schneider and S. Granick, Macomolecules 25, 5054-5059 (1992). 36 D. R. Iyengar and T. J. McCarthy, Macromolecules 23, 4344-4346 (1990). 37 B. O'Shaughnessy and D. Vavylonis, Eur. Phys. J. E 11, 213-230 (2003). 38 C. J. Clarke, R. A. L. Jones, J. L. Edwards, K. R. Shull, and J. Penfold,

Macromolecules 28, 2024-2049 (1995). 39 M. A. Plunkett, Thesis, Royal Institute of Technology, 2002. 40 B. Zhao and W. J. Brittain, Prog. Polym. Sci. 25, 667-710 (2000). 41 O. Prucker and J. Rühe, Macromolecules 31, 592-601 (1998). 42 W. Zhao, G. Kraush, M. H. Rafailovich, and J. Sokolov, Macromolecules 27, 2933-

2935 (1994). 43 M. Geoghegan, C. J. Clarke, F. Boue, A. Menellw, T. Russ, and D. G. Bucknall,

Macomolecules 32, 5106-5114 (1999). 44 K. Ebata, K. Furukawa, and N. J. Matsumoto, J. Am. Chem. Soc. 120, 7367 (1998). 45 X. Yang, J. Shi, S. Johnson, and B. Swanson, Langmuir 14, 1505 (1998). 46 D. E. Bergbreiter, J. G. Franchina, and K. Kabza, Macomolecules 32, 4993 (1999).

87

47 V. Koutos, E. M. Van der Vegte, E. Pelletier, A. Stamouli, and G. Hadziioannou,

Macromolecules 30, 4719 (1997). 48 V. Koutos and E. M. Van der Vegte, Macromolecules 32, 1233 (1999). 49 M. Husseman, E. E. Malmström, M. McNamara, M. Mate, D. Mecerreyes, D. G.

Benoit, J. L. Hedrick, P. Mansky, E. Huang, T. P. Russell, and C. J. Hawker,

Macromolecules 32, 1424-1431 (1999). 50 J. Dao, D. G. Benoit, and C. J. Hawker, J. of Polym. Sci.: Part A: Polymer

Chemistry 36, 2161-2167 (1998). 51 M. E. Schrader, in Contact angle, Wettability and Adhesion, edited by K. L. Mittal

(Koninklijke Wohrmann BV, Zeist, 1993). 52 A. V. Pocius, Adhesion and Adhesives Technology (Hanser, 1997). 53 F. M. Fowkes, J. Phys. Chem. 67, 2538 (1963). 54 R. Menescal, R. West, and C. Murray, Macromolecules 24, 329 (1991). 55 H. Kobayashi and M. J. Owen, Macromolecules 23, 4929-4933 (1990). 56 J. H. Clint and A. C. Wicks, J. Adhes. Adhesives 21, 267 (2001). 57 K. M. R. Kallury, M. Thompson, C. P. Tripp, and M. L. Hair, Langmuir 8, 947-954

(1992).

CHAPTER 4

Interfacial Toughness

Measurements

91

4.1 Mechanism of Adhesion

Four main mechanisms of polymer adhesion have been suggested: mechanical

interlocking, diffusion, electrostatic forces and adsorption1.

Mechanical Interlocking

In this mechanism, also referred to as mechanical keying, the adhesive fills the

irregularities in the surface, and upon solidification the adhesive and the surface interlock.

Mercury amalgam for filling tooth cavities is a typical example of this adhesion

mechanism.

Diffusion Mechanism

This mechanism is based on interdiffusion of polymer chains across the interface.

Interdiffusion is possible when polymers are mutually soluble and the polymer segments

have sufficient mobility, which occurs at temperatures above the glass transition (Tg).

When the temperature of a polymer is above Tg, the free volume that exists between the

segments is large enough for other polymer chains to diffuse through the material. For

simple organic materials, solubility can be expressed in terms of solubility parameters (δ),

which are tabulated2. Two materials are soluble in each other when the values of δ are

similar. This is valid for any polymer/polymer or polymer/solvent system. Interdiffusion

usually does not occur when one of the polymers is highly cross-linked, crystalline, below

Tg or if the δ values of the two components are too dissimilar.

Electrostatic Forces

Particles with the same polarity repel each other whereas particles with opposite

polarities attract each other. Electrostatic forces are the second strongest molecular

interactions after covalent bonding3. They are responsible for the formation of ionic bonds

between atoms and molecules which bear a charge, as in a case of ionic crystals.

92

Adsorption Mechanism

When a good molecular contact is established between two materials, bonds may

form due to interatomic and intermolecular forces acting along the interface. Depending on

the nature of the interactions, the following interfacial bonds are formed: a) strong primary

bonds, due to chemisorption, including ionic, covalent and metallic bonds; b) weak

secondary bonds, due to physisorption, including van der Waals (dipole-dipole interaction

and dispersion forces) and hydrogen bonds; c) donor-acceptor interactions (acid-base

interactions) with an intermediate strength.

4.1.1 Donor-Acceptor Interactions

According to the Lewis nomenclature, a Lewis acid is an electron acceptor and a

Lewis base is an electron donor. Partially halogenated ethylene based polymers (PVC),

acrylic acid copolymers, solvents such as dichloromethane (CH2Cl2) and chloroform

(CHCl3), and surfaces of silica and Fe2O3 type are Lewis acids. PMMA, PS,

polycarbonates, polyimides, aromatic solvents (benzene and tetrahydofuran), and surfaces

of Al2O3 type are Lewis bases. Amines, alcohols and polyamides are the examples of both

electron acceptors and donors.

Acceptor-donor interactions can have a profound effect on adhesion through

changing adsorption characteristics, as demonstrated in Figure 4-1. This figure, showing

PMMA deposited on a silica surface from various solvents1, demonstrates that with an

increase of both basicity and acidity of solvent the amount of PMMA adsorbed on the silica

surface decreases. This is because basic solvents compete with basic polymer for the acidic

SiOH groups and, on the other hand, acidic solvents compete with acidic silica substrate for

the basic polymer1.

The system used in the present study consists of basic PS, PS(Br) which is expected

to be slightly more basic than PS (in terms of Lewis classification) due to the bromine atom

93

in the para position, a basic solvent (toluene) and acidic silica. Both PS and PS(Br) compete

with toluene for SiOH sites on silica. It can be assumed that more basic PS(Br) adsorbs

preferentially. This is also supported by the fact that toluene is a good solvent for PS

whereas the solubility of PS(Br) in toluene decreases with an increasing degree of

bromination4.

Figure 4-1: Adsorption of basic PMMA onto acidic silica from basic, neutral

and acidic solvents1.

The interaction parameters χ for PS/toluene and PS(Br)/toluene systems were calculated

according to following equation2:

2)(3.0 jii

RTV δδχ −+= (1)

where Vi is the molar volume of polymer, δi and δj are the tabulated solubility parameters of

polymer and solvent. Calculated interaction parameters for PS and PS(Br) are χPS = 0.35

and χPS(Br) = 0.36, showing not a big difference. It has to be noted that the solubility

parameters are semi-empirical values and therefore the χ’s are only estimates.

94

4.2 Interfacial Fracture Toughness

Fracture occurs when a sufficient amount of energy is released by the growth of a

crack to decrease the overall energy of the system. The energy released comes from the

stored elastic (potential) energy and depends on the loading system. The interfacial fracture

energy Gc, also called the critical strain-energy release rate, can be defined as the energy

required per unit area to propagate a crack. It depends on the rate of crack propagation and

the temperature. Fracture energy can be divided into two components1: intrinsic adhesive

fracture energy Go and energy lossΨ.

Gc = Go + Ψ (1)

Go is the energy required to fracture the bonds in order to propagate a crack through a unit

area. Ψ is the energy dissipated in viscoelastic and plastic deformations during the crack

propagation. When viscoelastic and plastic energy losses (Ψ) are negligible, then Go is a

direct measure of interatomic and intermolecular bonding forces along the interface. When

only secondary bonds are present and the crack propagates through a homogeneous

material, then:

Go = 2γ (2)

where γ is the surface free energy of the material.

In the case of fracture of an adhesive joint, the crack may propagate either in the

materials forming the joint or along the interface. Taking this behaviour into consideration,

Go can be expressed as:

Go = iGo(interfacial) + aGo(adhesive) + sGo(substrate) (3)

where Go(interfacial) is the intrinsic fracture energy of interfacial failure, Go(adhesive) is

the intrinsic fracture energy of cohesive failure in the adhesive, and Go(substrate) is the

intrinsic fracture energy of cohesive failure in the substrate. The pronumerals i, a and s are

95

the area fractions (i + a + s = 1). When fracture occurs only in the interface then Go =

Go(interfacial).

Stress that acts on the material around the crack tip whilst the chemical bonds and physical

bonds are still intact can be expressed as the stress intensity factor (K). Fracture occurs

when K exceeds a critical value Kc, called critical stress intensity factor or fracture

toughness. Gc and Kc are related material properties.

4.2.1 Modes of Fracture

During fracture, new surfaces are created within the fractured material. Fracture can

occur under various conditions defined as modes of fracture5:

Tensile fracture

A direct mechanical load is applied to a material, and the stress is increased continuously

until the material breaks. This type of fracture is uncontrolled and catastrophic.

Fatigue

An alternating stress is applied to a material, simulating loading cycles. After a certain

number of these loading cycles the material fractures. The stress applied is usually

considerably smaller than that required for a tensile fracture.

Creep fracture

Also called static fatigue, this type of fracture occurs when a constant load is applied for a

certain amount of time. The longer the time allowed, the smaller the stress required to

fracture the material. Creep fracture is also a function of temperature and environmental

conditions.

96

Wear/Abrasion

This is a special case of fracture, where small particles of the material are broken off the

surface. An example of this fracture would be the wear of rubber.

Environmental Stress Cracking

Material fractures under small stresses when exposed to a chemically active environment as

in the case of corrosion in metals, glass and polymers. In polymers, the cracking can also be

caused by swelling in water. In this case, water does not react chemically with the polymer

but causes a stress build up within the material.

Crack propagation

Crack propagation occurs in all of the above fracture modes. Depending on how the

stress is applied to a crack, there are three different modes of loading – mode I, II and III6.

Mode I is a cleavage or tensile opening mode, mode II is an inplane-shear mode and mode

III is an antiplane-shear mode (Fig. 4-2).

Figure 4-2: Modes of loading6: a) mode I, b) mode II, c) mode III.

97

4.3 Thermosets

Thermosets are cross-linked polymers forming a random three-dimensional network

during the curing process. Epoxy resins, unsaturated polyesters, phenol-formaldehyde and

amino resins are typical examples of thermosets. These materials have to be molded into a

desired shape during polymerisation and cannot be re-molded after the solidification

because they do not melt or flow at high temperatures. High modulus and resistance to

solvents makes them ideal for applications such as high strength adhesives and matrix

materials for fibre-reinforced composites.

Epoxy resins are the toughest thermosets commercially available. Fracture

toughness Gc is usually in a range 100-300 J/m2. Critical stress intensity factor KIc,

commonly increases with an increasing degree of cure, which is achieved either by higher

curing temperature or longer cure times6[and ref. there in].

There are three modes of crack propagation in epoxy resins6:

Stable brittle propagation

This mechanism is typical for fully cured polymers tested at low temperatures. The

fractured surfaces exhibit no features and Gc is low.

Stable ductile propagation

This crack propagation mechanism occurs in under-cured resins at high temperatures.

Fracture toughness is generally high. Fractured surfaces have a ridged structure.

Unstable brittle propagation

The crack propagates by a stick/slip mode. In this mode the crack jumps upon initiation and

then stops. The fracture surface shows crack arrest lines or broad bands. At the arrest zone,

98

the crack deviates from its original trajectory and the fractured surfaces tend to interlock.

This interlocking along with the plastic deformation around the crack tip results in a

blunting of the crack tip. Initially, the crack grows very slowly then it slips rapidly and

stops when the applied KI has decreased. KI is defined in section 4.6.

It was shown that resins with a high yield stress exhibit continuous crack propagation

whereas resins with low yield stress undergo stick/slip crack propagation (Fig. 4-3) 6.

Figure 4-3: Schematics of a) stable and b) stick/slip mode of crack propagation6.

99

4.4 System Studied

The kinetics and overall reaction of diglycidylether of bisphenol A (DGEBA) epoxy

resin cured with amino hardeners has been well characterised by several researchers using

differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA), near infra-

red spectroscopy and solid-state nuclear magnetic resonance7-11. DGEBA fully cross-linked

with aliphatic amines has a relatively low Tg but it is well above room temperature (see

results section).

In the present study, monomeric DGEBA [DER resin 332] from Fluka (MWav = 340

g/mol) was cross-linked with two different curing agents: 1,5-diamino-3-methylpentane

(DMP) [DYTEK A amine] from Sigma-Aldrich and Jeffamine D230 from Aldrich. DMP

was chosen as an analogy to hexamethylene diamine (HMDA) which is known to form

highly cross-linked networks12. DMP is a liquid at room temperature and therefore easier to

mix with DGEBA than HMDA which is crystalline at room temperature. In addition,

HMDA had a tendency to exude from DGEBA during the curing process and to form

“blooms” on the surface. DMP showed no sign of such segregation. Jeffamines are curing

agents used to form systems with a very low Tg ranging from around -55oC to 85oC.

Chemical structures of DGEBA, DMP and D230 are shown in Figure 4-4.

4.4.1 Mechanism of Cross-linking

Epoxy resins contain reactive oxirane groups capable of reacting with various

chemical groups (hydroxyl, carboxyl, amino and cyano) via nucleophilic substitution. In

this study, aliphatic primary diamines (DMP and D230) were used as curing agents (Fig. 4-

4). In the initial step, the primary amine adds its active hydrogen to the epoxy group. The

resulting secondary amine then reacts with another epoxy group. Because there are two

active epoxy groups on each DGEBA molecule and two amino groups on both DMP and

D230 molecules, a three-dimensional network is formed upon curing (Fig. 4-5).

100

Figure 4-4: Structures of epoxy and diamino curing agents.

C

CH3

CH3

O

H2C

CHH2C

O OCH2

HCCH2

O

Monomeric diglycidylether of bisphenol A (DGEBA)

H2N NH2

CH3

1,5-diamino-methylpentane (DMP)

H2NHC C

H2O

CH3

CH2

HC NH2

CH3

x

General formula of Jeffamine

101

Figure 4-5: Mechanism of Cross-linking between di-epoxy resin and diamino hardener

system.

H2CHC

O

HC CH2

O

NH2H2N+

HC

HC

OHOH

CH2

CH2

NH

NH

NH2H2N

H2CHC

O

HC CH2

O

HC

HC

OHOH

CH2H2C

N

CH2

N

H2C

N

CH2

CH2

N

H2C

H2C

***

*

*

*

2

102

4.4.2 Characterisation of Epoxy System

The DGEBA/DMP system was analysed using a Perkin-Elmer DMA 7e in three

point bend mode under the following conditions: temperature scan rate 10oC/min, load 100

mN, dynamic load 80 mN, frequency 1Hz and nitrogen purge 20cc/min. The ratio of loss

modulus to storage modulus (E’’/E’ known as tan δ) was measured. The temperature when

tan δ starts to increase rapidly was taken as an analogue to the Tg measured by DSC. This

work was done in collaboration with Nathan Jones (UOW). The samples of various cure

regimes were tested by DMA (Table IV-1). Figure 4-6 shows that the values of Tg do not

change substantially when the time of the post-cure regime is varied. The average Tg value

is about 118oC. The Tg of Sample A (cured at 40oC for 24 hour) was near 105oC which is

much higher than expected. Typically, Tg of cured epoxies tends to be from 10oC to 15oC

above the cure temperature. In this case, Tg was more than 60oC above the cure

temperature. This could be caused by an exothermic reaction during the cure. Such

behaviour was observed when HMDA was used as a crosslinker. The system had to be

heated during mixing to keep HMDA in a liquid form. HMDA and DMP are structurally

very similar and therefore it is possible that there was an exothermic reaction during cure

that caused such a large increase in Tg.

DGEBA/D230 was characterised by Gaillard13. Reported Tg measured by Differential

Temperature Analysis (DTA) at heating rate 10oC/min was 82 ± 3oC.

Table IV-1: Cure regimes of DGEBA/DMP system

Sample ID Pre-cure regime Post-cure regime

A 24h @ 40oC

B 24h @ 40oC 24h @130oC

C 24h @ 40oC 36h @130oC

D 24h @ 40oC 48h @130oC

E 24h @ 40oC 60h @130oC

F 24h @ 40oC 72h @130oC

103

Figure 4-6: Tg of various cure regimes of DGEBA/DMP measured by DMA.

4.5 The Fracture Mechanism of Thermosets

It has been suggested that thermosets (cross-linked polymers) undergo brittle

fracture via a localised shear yielding mechanism6. When a crack is initiated, the energy

stored in the material starts to dissipate around the crack tip. As this amount of energy

increases, due to an increasing applied force, the growing crack becomes more and more

blunt and the toughness of the material in a vicinity of the crack tip increases. As the force

increases, the material around the crack tip starts to yield until fracture occurs. This fracture

can be either brittle or ductile. This mechanism is different from crazing which occurs

predominantly in thermoplastic (non-crosslinked) polymers, where the volume of the

crazed material increases. During the process of crack propagation, features called crazes

are formed. They are planar crack-like defects (microcracks) bridged by nanometer size

fibrils. These fibrils are load-bearing. They break down by either chain scission or dis-

entanglement14-16. The crazing mechanism has been well described elsewhere14-17.

A B C D E F100

105

110

115

120

125 Tg by Tan delta Onset [oC]

Tem

pera

ture

[o C]

Sample ID

104

Both mechanisms involve plastic deformation around the crack tip. During crazing

the volume of a material increases whereas during shear yielding the volume is constant.

Both crazing and yielding contribute to brittle fracture. It has been shown6 [and ref. there in] that

cross-linking of polymer chains inhibits crazing and leads to yielding. Thermosets are a

typical example of materials exhibiting such behaviour. When the polymer is tough and

capable of absorbing a large amount of energy via plastic deformation before it fails, the

homogeneous shear yielding results in ductile fracture. When shear yielding is localised

around the crack tip, the crack grows by chain pull-out or scission resulting in brittle

fracture.

4.6 Interfacial Toughness Measurements

Design for the interfacial toughness measurements used in this study was based on

work by Smith et al18,19. The authors studied adhesion between polymer/non-polymer

interfaces using an asymmetric double cantilever test (ADCT), where a specimen is

“sandwiched” between two glass beams. The aim of their work was to promote adhesion

between PS/glass and PS/Si, by modifying the glass and silicon surfaces with PS-PMMA

and PS-PVP block copolymers. The asymmetric double cantilever beam specimen is shown

in Figure 4-7. When a razor blade is inserted into the sample between the polystyrene and

the glass, as indicated in Figure 4-7, both mode I and II (tensile and shear) stress patterns

are generated which affect the propagating crack. This is due to asymmetry of the beams

caused by different beam stiffnesses.

A property combining both tensile and shear stress components is known as mode

mixity, Ψ. It can be expressed in terms of KII and KI (mode II and I) stress intensity factors

(Eq. 4). When Ψ is very small, the crack propagates along the weakest interface. By

adjusting thicknesses of the beams, the crack can be driven into a desired interface. The

modes of crack propagation I and II are described in section 4.2.

I

II

KK1tan−=Ψ (4)

105

Figure 4-7: Schematics of the “sandwich” structure for testing the PS/glass interface19.

Interfacial fracture toughness Gc depends strongly on the ratio of the beam thicknesses

h1/h2, where h1 is the thickness of the top beam and h2 is the thickness of the bottom beam.

Smith et al18 found that when the aspect ratio (h1/h2) > 0.86, the sample fails at high Gc and

there is high scatter in the data. The crack tends to oscillate between the PS layer and the

PS/epoxy interface. When (h1/h2) = 0.71, the crack is driven along the PS/glass interface

and Gc reaches a plateau. For (h1/h2) < 0.71 the thinner upper glass beam breaks. Smith et

all 18 reported that adding copolymers increased Gc from 1 J/m2 (unmodified PS/glass) up

to 25 J/m2.

Smith18 applied the same sample geometry to test the PS/silicon interface. The only

difference was that the silicon had to be affixed to a rigid inflexible foundation because

silicon is very brittle and fractures upon small displacement in bending. The sample

geometry is shown in Figure 4-8. In some cases, the authors were able to measure the

interfacial toughness even though it would be expected that the shear component of the

stress would drive the crack towards the more compliant polymer side19. A razor blade was

inserted into the sample and left in place for 24 hours. Initially, the crack propagated into

the polymer but the energy stored in the bent beam turned the crack back to the interface.

106

Figure 4-8: Schematics of the “sandwich” structure for testing PS/silicon interface18.

Gc of the PS/silicon interface was determined to be around 1 J/m2. Addition of PS-PMMA

block copolymer increased Gc up to 44 J/m2. The interfacial toughness was found to be

independent of the epoxy thickness but increased with increasing thickness of the block

copolymer.

Creton et al16 used the ADCT to study polymer/polymer interfaces reinforced with

block copolymers. The authors observed four different mechanisms of interfacial failure: a)

chain pullout when chains were shorter than their entanglement length; b) small plastic

deformation around the crack tip when the chain density (σ) (equation 12 in Section 3.6)

was low (low Gc); c) craze breakdown of the interface with higher σ; and d) chain scission

when the interface was very strong due to high σ. The authors concluded that Gc increased

with increasing areal density of the copolymer chains.

The mechanisms of the chain pull-out and scission were also reported by Norton et

al20 in a study of adhesion between epoxy/PS interfaces modified with PS-COOH chains

end-anchored to the epoxy. The authors found that Gc increased with increasing PS-COOH

107

length up to N = 838. Beyond this length, Gc started to decline. Other applications of

ADCT can be found elsewhere17,21-25.

4.6.1 Asymmetric Double Cantilever Test

The Asymmetric Double Cantilever Test (ADCT) is a special form of the double

cantilever beam specimen used for testing adhesion strength of adhesive joints6. Due to the

existence of a bi-material interface, deformation induced by both tensile and shear stress

components contribute to the interfacial toughness.

The applied interfacial fracture toughness is a function of the Young’s modulus of

both the epoxy E1 and the silicon wafer E2, the thicknesses of the beams h1 and h2, the crack

opening displacement ∆ (which equals the thickness of the razorblade), and the measured

crack length a.

As reported in the work by Smith18, the asymmetric double cantilever beam

geometry makes it possible to direct the crack along the preferred interface. Gc can be

calculated from an expression based on a model of a cantilever beam on an elastic

foundation derived by Kanninen26:

(5)

ahC 1

1 64.01+= ahC 2

2 64.01+=

Fracture energy is determined from the crack length a, which is measured during the test.

++∆= 3

1322

32

311

31

322

22

311

4

322

311

2

83

ChEChEChEChE

ahEhEGc

108

In this study, the razor blade was inserted between the epoxy and the silicon wafer

modified with a polymer coupling agent. A schematics of the system is shown in Figure 4-

9. The crack propagates by pushing the razor blade further along the interface.

Figure 4-9: Asymmetric double cantilever beam specimen; Gc = f (E, h, ∆2, a4).

As shown in Figure 4-9, the sample consists of two dissimilar beams. The

asymmetry arises not only from the beams having different thicknesses but also from

differences in the material properties of these beams. As a result, there are both tensile and

shear components of the stress acting ahead of the crack tip (mode I and II). The

propagating crack tends to deviate towards the more compliant material.

4.7 Tailoring the Interfacial Toughness using Polymeric Coupling Agents

In a similar manner to other studies where the block copolymers were used to

improve adhesion between two dissimilar materials 16-24, in this investigation the mixtures

of functionalised (active) PS and plain (inactive) PS chains were deposited onto silicon

substrates via the silane end-functional groups. Functionalisation of PS was achieved by

bromination, as described in section 2.1.2, and further conversion to amino groups as

described in section 3.6.3. The mechanism of depositing the mixed monolayers of PSBr0.1-

SiClMe2 (active polymer) and PS-SiClMe2 (inactive polymer) was explained in section

Razor blade Epoxy

Silicon wafer

Coupling agent

h1

h2

∆

Crack length a

109

3.4.2. By increasing the ratio of active/inactive polymer chains chemically anchored to the

silicon surface, the areal density σ of the active chains and therefore the number of

functional groups available for reaction was increased. A random distribution of functional

groups along each functionalised polymer chain was assumed.

When a large amount of epoxy is cast directly onto a glass or a silicon surface and

cured, the stresses created by the epoxy shrinking and by the difference in the thermal

expansion coefficients between the epoxy and the substrate are large enough to “strip” the

top of the glass surface or break the silicon wafer. When the block epoxy is pre-cured and

only a thin layer of the same or similar epoxy is used to glue the epoxy to the substrate,

these stresses are minimised. The pre-cured DGEBA/DMP blocks were cut into 7mm x

5mm x 75mm (H x W x L) beams27 and polished. The modified silicon substrates 15mm x

75mm (W x L) were glued to pre-cured epoxy beams using DGEBA/DMP or

DGEBA/D230 as an adhesive. One end of silicon wafer was covered by PTFE foil for an

easier crack initiation. The samples were pre-cured at 40oC, removed from the oven, and

the Si/epoxy interface was pre-cracked by inserting a razor blade. The razor blade was

removed and the samples were put back into the oven, cured for 3h at 80oC and post-cured

for 3h at 150oC 13. DGEBA and DMP were mixed at 2.1 : 1 molar ratio, DGEBA and D230

were mixed at 2 : 1 molar ratio. After curing, Gc was measured by ADCT. The photographs

of prepared samples are shown (Fig. 4-10).

Figure 4-10: Photograph of an ADCT specimen.

Propagated crack Si wafer

Epoxy beam

Razor blade

110

A razor blade was inserted into the pre-cracked interface and pushed forward until

the crack started to propagate. The razor blade was left inserted in a sample overnight to

allow enough time for the crack to equilibrate. The next day the crack length a was

measured, looking through the top of the epoxy beam, using an optical microscope. The

razor blade was then pushed further until the crack started to propagate again, and the

process was repeated as many times as possible. The limitations were the sample length

and/or premature fracture of the silicon beam.

111

4.8 ADCT Experiments

Experiment 1: ADCT of the epoxy/Si beam samples with mixed (PSBr/PS)-SiCl(Me)2

monolayers deposited on Si wafers

Mixtures of 1, 2.5, 5 and 10 wt % PSBr0.1-SiClMe2 in PS-SiClMe2 were dissolved in

toluene at a total concentration of 1 wt % polymer. Cleaned silicon wafers were dipped into

these solutions and stirred overnight using an orbital stirrer. Then the samples were

removed and cleaned using toluene by Soxhlet extraction overnight to remove any

physisorbed chains. After extraction, the samples were dried under a nitrogen flow and

stored in a desiccator. Thickness and surface energy of the modified silicon surfaces were

measured by ellipsometry and contact angle measurement. The results are shown in Figures

4-11 and 4-12. After conversion the bromine groups to the amino groups by the Heck

reaction, the layer thickness was found to be either identical or slightly higher. The small

increase in the thickness is probably caused by some impurities remaining on the surfaces

after the Heck reaction.

The DGEBA/DMP blocks were prepared and glued to the modified silicon wafers by

methods mentioned above. Both DMP and D230 were used as curing agents for DGEBA to

glue a DGEBA/DMP block to the modified silicon wafers. Gc values, calculated using

equation 5, are shown in Figure 4-13. The sample parameters are shown in Table IV-2.

Note that the BS(Br) [%] in all the plots below relates to the initial concentrations of PSBr

chains in the toluene deposition solutions. Note also that 2 different cross-linkers were

used.

Table IV-2: Gc and sample parameters

Sample [%]

∆ [m]

a [m]

h1(top) [m]

h2(bottom) [m]

E1(top) [N/m2]

E2(bottom) [N/m2]

Gc(average) [J/m2]

5 Jeff 1.00E-04 3.80E-03 7.00E-03 3.70E-04 3.00E+09 1.30E+11 73 10 Jeff 1.00E-04 3.40E-03 7.00E-03 3.70E-04 3.00E+09 1.30E+11 101 2.5 DMP 1.00E-04 3.90E-03 7.00E-03 3.70E-04 3.00E+09 1.30E+11 103 5 DMP 1.00E-04 3.40E-03 7.00E-03 3.70E-04 3.00E+09 1.30E+11 143 10 DMP 1.00E-04 3.20E-03 7.00E-03 3.70E-04 3.00E+09 1.30E+11 230

112

Figure 4-11: Monolayer thickness of the modified Si surfaces (Exp 1).

Figure 4-12: Surface energy of the modified Si surfaces (Exp 1).

0 2 4 6 8 100

2

4

6

8

10

12

14

16

18

20

Original thickness Thickness after Heck

Mixed PS(Br) and PS monolayersTh

ickn

ess

[Å]

PS(Br) [%]

0 2 4 6 8 100

5

10

15

20

25

30

35

40

45

50

Surface energies of PS(Br)/PS monolayers on Si wafers

γ [m

J/m

2 ]

PS(Br) [%]

γ Total γ Disp γ Polar

113

Figure 4-13: Interfacial fracture toughness measurements using ADCT (Exp 1).


monolayers deposited on Si wafers using DGEBA/DMP as a glue layer

Toluene solutions (1 wt %) of PSBr0.1-SiClMe2 (2, 4, 6, 8, 10, 13, 16 and 20%) in

PS-SiClMe2 were prepared. The same procedures as in the experiment 1 were used to

prepare the polymer modified Si surfaces for ADCT. DGEBA/DMP was used as the glue

layer. Thickness and surface energy of the mixed polymer monolayers are shown in

Figures 4-14 and 4-15. Figure 4-16 shows the variation of interfacial toughness with

increasing amount of the active polymer on the silicon surface. Note that the red line is only

a guide for the eye.

0 1 2 3 4 5 6 7 8 9 10 110

50

100

150

200

250

300

350

400

DGEBA/DMP and DGEBA/D230 glue

Gc [

J/m

2 ]

PS(Br) [%]

DGEBA/DMP DGEBA/Jeff

114



0 2 4 6 8 10 12 14 16 18 20 2202468

1012141618202224262830

Original thickness Thickness after Heck


ickn

ess

[Å]

PS(Br) [%]

0 2 4 6 8 10 12 14 16 18 20 220

5

10

15

20

25

30

35

40

45

50

Surface energies of PS(Br)/PS monolayers on Si wafers

γ [m

J/m

2 ]

PS(Br) [%]

polar disp total

115



monolayers deposited on Si wafers using DGEBA/D230 as a glue layer

Experiment 2 was repeated using fresh solutions of PSBr0.1-SiClMe2 (2, 4, 6, 8, 10,

12, 16 and 20%) in PS-SiClMe2. Thickness of the mixed polymer monolayers after the

Soxhlet extraction and after the Heck reaction is shown in Figure 4-17. DGEBA/D230 was

used as the glue layer. The resulting plot of Gc vs. solution concentration of PSBr0.1-

SiClMe2 is shown in Figure 4-18. The points that are not showing error bars were obtained

from only one measurement before the silicon substrate fractured. The surface energies of

these samples were not measured.

0 2 4 6 8 10 12 14 16 18 20 22

-10123456789

10111213

DGEBA/DMP layer

Gc [

J/m

2 ]

PS(Br) [%]

116



0 2 4 6 8 10 12 14 16 18 20 22

0

10

20

30

40

50DGEBA/D230 layer

Gc [

J/m

2 ]

PSBr [%]

0 2 4 6 8 10 12 14 16 18 20 222

4

6

8

10

12

14

16

18

20

22


ickn

ess

[Å]

PS(Br) [%]

Original thickness THickness after Heck

117


monolayers deposited on Si wafers using DGEBA/D230 as a glue layer

The same conditions and procedures were used as in experiment 3, using fresh

solutions. Figures 4-19 and 4-20 show surface analysis of the mixed polymer layer. Figure

4-20 shows Gc measurements.

The control sample CS1 was prepared by depositing pure PS-SiClMe2 (inactive PS)

on the Si substrate. The sample was then glued to the DGEBA/DMP beam with

DGEBA/D230 glue and tested by ADCT as described before. Gc results were plotted along

with the results for the mixed polymer layers. The red line in Figure 4-20 serves only as a

guide.


0 2 4 6 8 10 12 14 16 18 20 220

2

4

6

8

10

12

14

16

18

20 Original thickness Thickness after Heck

Mixed PS(Br) and PS monolayers

Thic

knes

s [Å

]

PS(Br) [%]

118



0 2 4 6 8 10 12 14 16 18 20 22

5

10

15

20

25

30

35

40

45

50

Surface energies of PS(Br)/PS monolayers on Si wafersγ

[mJ/

m2 ]

PS(Br) [%]

Disp Polar Total

0 2 4 6 8 10 12 14 16 18 20 220

20

40

60

80

100

120

Mixed polymer layer Control sample

DGEBA/D230 Layer

Gc [

J/m

2 ]

PS(Br) [%]

119

4.8.1 Discussion of Gc

In Experiment 1, the thickness of the mixed polymer layers was found to be quite

uniform, averaging around 8Å. The polar component of surface energy shows a slightly

increasing trend with an increasing concentration of brominated PS. Several samples broke

before any reasonable crack lengths could be measured. The remaining samples show

increasing Gc with increasing concentration of PS(Brx). The slope is steeper, and the values

generally higher, for the samples where DGEBA/DMP glue was used. The average

interfacial toughness values of these samples were calculated to be between 100 and 230

J/m2. In the case of DGEBA/D230, the values were between 75 and 100 J/m2.

In Experiment 2, the thickness data are more scattered, oscillating around 20Å. The surface

energies were uniform across the range of concentrations, and the polar surface energy was

very low. Interfacial fracture toughness of these epoxy/Si interfaces was very low

averaging between 2 and 8 J/m2. The general trend is a slight increase in Gc with increasing

concentration of PS(Brx).

In experiment 3 (D230), The Gc values are very low as well, averaging between 5

and 20 J/m2. After an initial scatter, the Gc values level off and oscillate around 15 J/m2.

In Experiment 4, the average thickness was around 14Å and γpolar are scattered

between 2 and 15 mJ/m2. Gc values were between 70 to 95 J/m2 and remained steady across

the range of concentrations. The interfacial toughness of the control sample CS1

(epoxy/PS) was calculated to be 10 J/m2. These results show that adding the PS(Brx) and

converting it to PS(NH2) increased the interfacial toughness by factor of about four.

Some correlation between γpolar and Gc was observed. The possible meaning of this

observation will be further discussed in Chapter 5.

Two different cross-linkers, DMP and D230, were used to glue DGEBA/DMP

blocks to the polymer modified silicon surfaces, in order to investigate any effect of

120

different degree of cross-linking and different Tg on the characteristics of the crack

propagation. Due to high scatter between the experiments, no conclusions could be made.

The only observation made was that the samples prepared with DGEBA/D230 had a higher

success rate, meaning that there was less breaking of the silicon wafers. This behaviour

would suggest that the more flexible DGEBA/D230 glue relieves more stress that is

generated during the crack initiation process and therefore prevents breaking of the silicon.

The stiffer DGEBA/DMP glue, however, does not have as a high capacity to relieve the

stress, and therefore breaking of the silicon is more frequent.

The expression for Gc calculation, derived by Kanninen (equation 5), assumes both

equal width and length of the beams. In this study, the length of the beams was identical but

the width was different. The silicon beam was wider than the epoxy beam to prevent

breaking of the silicon. This change in the sample geometry would introduce an error in the

final Gc values. It was shown28 that this error is only small and therefore it was neglected in

this study. Another factor that affects the final fracture toughness is the mismatch between

the thermal coefficient of the epoxy and the silicon materials. The detailed description of

this subject can be found elsewhere28. For all experiments conducted in this study, the

sample geometry was kept constant. Therefore it was assumed that the errors caused by the

different sample dimensions and the thermal mismatch would contribute equally to Gc of

each sample and therefore were not taken into the account.

The surface energy measurements were not conducted after the Heck reaction

because of the risks of contamination and deactivation of amino functional groups.

121

4.9 References

1 A. J. Kinloch, Adhesion and Adhesives-Science and Techology (Springer, 1987). 2 J. E. Mark, Physical Properties of Polymers Handbook (AIP Press, New York,

1996). 3 A. V. Pocius, Adhesion and Adhesives Technology (Hanser, 1997). 4 R. Oslanec, (University of Wollongong, 2005). 5 E. H. Andrews, Fracture in Polymers (Oliver & Boyd, Edinburgh and London,

1968). 6 A. J. Kinloch and R. J. Young, Fracture Behaviour of Polymers (Applied Science

Publishers, London and New York, 1983). 7 B. G. Min and Z. H. Stachurski, Polymer 34, 4488 (1993). 8 B. G. Min, Z. H. Stachurski, and J. H. Hodgkin, Polymer 34, 4908 (1993). 9 B. G. Min, Z. H. Stachurski, J. H. Hodgkin, and G. R. Heath, Polymer 34, 3620

(1993). 10 M. E. Merritt, L. Heux, J. L. Halary, and J. Shaefer, Macromolecules 30, 6760-6763

(1997). 11 L. Heux, J. L. Halary, F. Laupretre, and L. Monnerie, Polymer 38, 1767 (1997). 12 L. Heux, F. Laupretre, J. L. Halary, and L. Monnerie, Polymer 39, 1269 (1998). 13 S. Gaillard, “Internship report,” (2001). 14 C.-Y. Hui, A. Ruina, C. Creton, and E. J. Kramer, Macromolecules 25, 3948-3955

(1992). 15 F. Xiao and W. A. Curtin, Macromolecules 28, 1654-1660 (1995). 16 C. Creton, E. J. Kramer, C.-Y. Hui, and H. R. Brown, Macromolecules 25, 3075-

3088 (1992). 17 H. R. Brown, Macromolecules 24, 2752-2756 (1991). 18 J. W. Smith, E. J. Kramer, F. Xiao, C.-Y. Hui, W. Reichert, and H. R. Brown, J. of

Mat. Sci. 28, 4234-4244 (1993). 19 J. W. Smith, E. J. Kramer, and P. J. Mills, J. Polymer Sci.:Part B: Polymer Physics

32, 1731-1744 (1994).

122

20 L. J. Norton, V. Smigolova, M. U. Pralle, A. Hubenko, K. H. Dai, E. J. Kramer, S.

Hahn, C. Berglund, and B. DeKoven, Macromolecules 28, 1999-2008 (1995). 21 J. Washiyama, C. Creton, and E. J. Kramer, Macromolecules 26, 6011-6020 (1993). 22 F. Xiao, C.-Y. Hui, and E. J. Kramer, J. of Mat. Sci. 28, 5620-5629 (1993). 23 K. Char, H. R. Brown, and V. R. Deline, Macromolecules 26, 4164-4171 (1993). 24 H. R. Brown, K. Char, V. R. Deline, and P. F. Green, Macromolecules 26, 4155-

4163 (1993). 25 J. Duchet, J. P. Chapel, and B. Chabert, Macromolecules 31, 8264-8272 (1998). 26 M. F. Kanninen, International Journal of Fracture 9, 83-92 (1973). 27 J. Benkoski, edited by P. communicacion (2003). 28 J. J. Bekonski, E. J. Kramer, H. Yim, M. S. Kent, and J. Hall, Langmuir 20, 3246-

3258 (2004).

CHAPTER 5

Discussion

125

5.1 Discussion

The results for interfacial fracture toughness (Gc) can be looked at as indirect

measurements of the interfacial properties. The initial analysis of Gc vs. initial

concentration of PSBr (Chapter 4) would suggest that within the range of concentrations of

the active PS(NH2) (2 – 20wt%) grafted to Si wafers as mixed monolayers of PS/PS(NH2),

there is a saturation in Gc values. The reproducibility between the experiments is poor,

however, and therefore it is not possible to draw any specific conclusions about the value of

the interfacial fracture toughness of these Si/epoxy interfaces. Possible explanations for this

behaviour are discussed below.

Based on the research of interfacial toughness between two immiscible polymers

and polymer/solid substrate, described in the introduction1-9, it would be expected that Gc

values would increase with increasing concentration of the active chains. In this study, the

interface containing only PS (no PS(NH2)) chains was significantly weaker than the

interfaces with PS(NH2) chains. Within each experiment, all non-zero PS(NH2)

concentrations gave essentially the same Gc. Between the experiments, however, Gc. varied

significantly. Possible reasons for this behaviour are discussed below.

Even though the silane coupling agents are widely used to improve water resistance

of various surfaces and interfaces, Plueddemann10 suggests that the siloxane bonds between

Si and silane coupling agents are not resistant to hydrolysis. This means that the siloxane

covalent bonds can be disrupted by water. The activation energy for this reaction is 98.6

kJ/mol10. The addition of benzoic acid during the hydrolysis process decreases this

activation energy to 25 kJ/mol, which is energy comparable to the strength of a hydrogen

bond. During the process of bonding of the silane coupling agents to the mineral surfaces

(M), there is a continuous process of bonding and de-bonding due to a reversible reaction

with water10:

M−O−Si− + H2O ⇔ M−O−H + H−O−Si−

126

Adding water to the system would drive the reaction to the right. Efficient adhesion can be

achieved when the equilibrium conditions are shifted to the left, towards the formation of

oxane bonds. This can be done by a) maximising the concentration of initial siloxane

(M−O−Si−) bonds, b) minimising the penetration of water to the interface and c) applying

polymeric material that “holds” silane molecules at the surface. The last case was

demonstrated in a study of water-assisted crack propagation between (3-

glycidoxypropyl)trimethoxysilane coated silicon wafers and DGEBA11. The authors

observed almost no effect of pH (varied from 7 to 11) on the interfacial fracture toughness.

In this study, the number of siloxane bonds at the interface was maximised by

treating the silicon wafers with hot piranha solution during the cleaning procedure. This

process is known to create a high concentration of Si-OH sites that are available for

reaction with silanes12. As mentioned in Section 3.6.3, when PS-SiClMe2 was deposited on

the silicon wafers, the polymer layers were stable even when exposed to acidic and basic

conditions. It was confirmed by ellipsometry that after adding 3 drops of 1M HCl in MeOH

solution to a PS-SiClMe2/silicon sample immersed in 25mL of THF/MeOH (9/1) and

stirred for 25 minutes, the PS layer retained its original thickness (~ 40Å). However, when

the mixed monolayers were deposited on silicon wafers, the polymer chains were washed

off during the last step of conversion when exposed to acidic and basic environments

(except the initial experiment 1). It is known that both acids and bases are catalysts for

hydrolysis and re-formation of the oxane bonds10. The ellipsometry measurements showed

that the thickness of the mixed PS(Br)/PS layers before the conversion was lower than that

of the plain PS-SiClMe2. Therefore the areal density of the mixed polymer chains (σ) was

lower than that of the plain PS chains. This might be due to the attraction of the polar

bromine groups of PS(Br) chains to the silicon oxide surface, which causes the chains to lay

flat on the surface and therefore prevent other chains from accessing the surface. The lower

σ might allow for the acid and the base to reach the surface and hydrolyse the siloxane

bonds. In this way, the NH2 groups of the converted polymers may act as conductors for the

acid and the base.

127

This washing off of the grafted chains could perhaps be prevented by using polymer

chains with tri-functional silane end-group instead of mono-functional silane end-group.

The mono-functional silane terminated polymer was ideal for producing a monolayer.

However, as discussed in Chapter 3, mono-functional silane does not produce stable layers.

The more stable layers are formed in case of tri-functional silanes through cross-linking

reactions of adjacent silanol groups. This cross-linking produces a layer that is more

resistant to hydrolysis13.

The disruption of the siloxane bonds found in this study is in agreement with the

results of Duchet et al14, who studied the stability of monofunctional alkylsilane layers with

Cl or N(CH3)2 hydrolysable groups, deposited on quartz slides and on nanometer size silica

particles, either from toluene solution or by vapour method. The samples were exposed to

0.1M KOH solution for various periods of time. The silane layer degradation was observed

by contact angle measurements.

In research done on polyester laminates10, single silanes were compared with mixed

silanes as treatment on glass microbeads in polyester casting. Methacrylate silane and

diamine-functional silane were mixed in a ratio 9/1. The authors found that the initial

adhesion improved by adding the diamine silane, however, after 24h of boiling in water, the

flexural strength decreased compared to the pure methacrylate silane. The explanation was

that the unreacted amine groups were hydrophilic enough to allow water to the interface

region.

Examples of hydrolysis of silane coupling agents can also be found in some

practical applications, such as glass-reinforced composites. Plueddemann10 studied the

flexural strength of glass-cloth-epoxy and glass-cloth-polyester composites after exposure

to hot aqueous solution of pH ranging from 2 to 10. In this case, silanes were used as

adhesion promoters. He found that the best retention of mechanical properties was in acid

or alkaline water. This situation is different to that described in the present study, however,

because the silane coupling agents are fixed between the composite layers so even if the

siloxane bonds are dis-lodged due to hydrolysis they reform upon heating. In the present

128

study, the samples are immersed in a THF/MeOH (9/1) solution during hydrolysis. There is

nothing to “hold” the polymer layers down to the silicon wafer, and hence the chains are

washed off into solution. This process of “washing off” of the chains is completely random

and varies from experiment to experiment.

Several analytical techniques were used by other researchers to examine the grafted

densities of silane layers. Duchet et al15 used solid state 29Si NMR spectroscopy, gas

chromatography and elemental atomic analysis to determine the grafting density of the

silane chains. When solid state 29Si NMR spectroscopy was used, the authors reported only

semi-quantitative results with accuracy about 4-5%. For long polymer chains, the accuracy

of the measurement would be expected to be even lower. In order to carry out gas

chromatographic analysis, the silane molecules were cleaved from the surface by

hydrofluoric acid etches. In the last method, an elemental atomic analysis, the grafting

density was determined from the atomic percentage of carbon. The disadvantage of this

method is that the carbons from the solvent molecules adsorbed at the surface are included

in the percentage and as a result, σ was overestimated. The authors reported that σ

decreased with an increasing length of silane chains because the bonded molecules

sterically hinder the reaction of additional molecules. Another technique used to analyse the

silane monolayers was transmission infrared spectroscopy, where the silane chains were

deposited on Aerosil silica16. The main reason for using silica particles was that the grafting

density is higher on highly curved surface of silica than on flat surfaces. This is due to

already mentioned steric effects of the chains which are more pronounced in case of the flat

surfaces. Also, the amount of grafted chains to analyse is much greater in cases of silica

particles where the surface area is significantly larger than in planar samples.

As shown above, the study of the grafting densities of the uniform silane

monolayers produces rather semi-quantitative results. To perform a quantitative analysis of

the mixed polymer monolayers chemically attached to a solid substrate is even more

challenging because the concentrations of the chains with particular atoms attached is very

low. Such analysis would be extremely demanding on the resolution of the available

surface analysis techniques. In this study, several samples with mixed polymer monolayers

129

were examined by grazing angle IR and XPS techniques, but no satisfactory results were

obtained.

For simplicity, it was assumed in this work, that the deposition of polymer chains

from solutions of different ratios between PS and PS(Br) chains would yield monolayers of

similar surface composition. Unfortunately, due to the washing off of the chains, this initial

information about the monolayer composition was lost. To perform a quantitative analysis

on the remaining partial monolayer on silicon wafer would be extremely challenging for the

reasons described above. The samples were therefore used for the asymmetric double

cantilever test (ADCT) assuming that both types of polymer chains were still present on the

surface.

In the study of adhesion between glass and polyethylene (PE), Duchet et al14

explored different regimes of the interfacial failure. Chlorosilane-terminated PE with

different chain lengths was grafted to the glass substrates and joined to a thick PE film. The

second glass slide was used to reinforce the PE film. The mechanism of failure depends on

the chain length and σ. For short chains with molecular mass below the entanglement

molecular mass, and low σ, the interface fails by chain pull-out. The Gc of such interfaces

was reported to be between 1 and 6 J/m2 14. For long chains and low σ, the interface fails

by chain scission. In both regimes, Gc varies linearly with σ (Gc ~ σ N2)17. In case of higher

σ, crazing involving chain disentanglement occurs, and Gc varies with σ2 (Gc ~ σ2)18.

In this study, the regime of long chains and low grafting density is applicable.

Based on the adhesion studies between the thermoplastic polymers and a solid substrate1-9,

it would be expected that the interfacial fracture toughness would increase with increasing

areal chain density. However, such behaviour was not observed. There appeared to be an

increasing trend in the Experiment 1 (Figure 4-13) when DMP was used as a cross-linker.

However, the plot contains only three measurement data and the error bars are large. For

this reason, the possible increase in Gc was not taken into further consideration. From the

dependence of Gc on initial solution concentration of the grafted mixed polymers layers

(PS/PSBr), it can be concluded that within the studied concentration range (0 – 20%) of the

130

active PS(NH2) chains the interfacial fracture toughness reached a plateau. Fracture

toughness of epoxies is usually in the range 100-300 J/m2. For the specific system of

DGEBA/HMDA (hexamethylene diamine), the reported Gc initiation value is around 570

J/m2 19. HMDA has a very similar structure and glass transition temperature (Tg) to the

DMP cross-linker used in this study (Chapter 4). Except for Experiment 1 (Section 4.7), the

highest Gc value measured was 95 J/m2 which is well below the fracture toughness of

epoxy. Therefore this plateau would be caused by the crack propagating along the interface

rather than within the bulk epoxy.

When the ADCT was performed on the sample with plain PS chains (i.e. no amine

functionality), the razor blade was easily inserted and the crack propagated in a continuous

manner without stopping. This behaviour is different from the samples with mixed

monolayers, where after a difficult crack initiation the crack would jump and then stop. In

this case, an increased force had to be applied in order to reinitiate the crack propagation.

This corresponded to the stick/slip propagation mechanism described in Chapter 4. Since

both plain PS and mixed PS/PSBr monolayer samples were prepared under the same

conditions, including the epoxy cure regime, this difference in crack propagation behaviour

can be related to the differences in the interfacial properties. This is clear evidence of a

higher Gc when PS(NH2) chains are present at the surface. Whether this is due to reaction

of the remaining chains with the epoxy, or if it is due to an increase in the surface

roughness, is not clear. The reason may be a combination of both.

It is not obvious why the samples from the initial experiment (Exp. 1) exhibited

different behaviour compared to the later experiments (Chapter 4). The samples were

prepared in the same way and exposed to same reaction conditions. The only difference

was that the silicon wafers used were sourced from two different suppliers. Silicon wafers

used in the initial experiment were supplied by Silicon Inc. (20Å native oxide layer),

whereas for the subsequent experiments, silicon wafers from Compart Technologies (17Å

native oxide layer) were used. The wafers were cleaned using the same procedure prior to

use. Wafers from both suppliers had the same dopant type and concentration, and

orientation 100. When control experiments were performed with mixed monolayers

131

deposited on both wafer types, the polymer chains were washed off under acidic conditions

from both wafers (Tables III-(13-15)). Considering these results, the difference in the

behaviour between the first initial experiment and the rest of the experiments is very

unlikely to be due to using different Si wafers.

A possible explanation for this behaviour might be in the testing technique itself. It

was shown that when bulk epoxy was tested in tensile mode (I), the stress intensity factor

(KIci) for crack initiation (i) varies with the degree of cure19. The same was observed for

GIci20. Clearly, there was a difference between Kci and Gci values for the crack initiation,

and those for the crack arrest (a), Kca and Gca 19. The evidence for different behaviour of Gci

and Gca was also given in a study of Mijovic et al20. The authors observed that for the

DGEBA/DETA (diethylene triamine) system, the GIci (initiation) increased with increasing

content of curing agent up to a certain concentration and then dropped slightly, whereas

GIca (arrest) changed only slightly. In another study19, the effect of the testing rate on the

fracture toughness was measured. The authors found that as the testing speed increased, KIci

decreased, but KIca remained constant. A variation of KIci with testing temperature was also

reported19. It was found that at low temperatures (below 0oC) the crack propagated in a

continuous manner. As the temperature increased, the crack propagation became unstable,

and KIci increased whereas KIca remained approximately constant.

This behaviour would suggest that GIca is less sensitive to the polymer chemistry

and testing conditions than GIci. On the other hand, in a study of the effect of various post-

cure times on the fracture energy of bulk epoxy20, the values of fracture energies of

initiation and arrest followed the same pattern – higher values of GIci and GIca were

observed within the initial stages of post cure. Later, GIci and GIca dropped slightly and

became nearly constant. The only difference between the initiation and arrest values was

that GIca was lower than GIci.

In this study, the crack in the epoxy/modified SiOx interface propagated in an

unstable stick/slip manner. The ADCT test was carried out manually and the crack length

was measured at the point where the crack arrested. This indicates that the Gc values

132

obtained from these measurements are actually Gca (arrest) values rather than Gci

(initiation). Gca measurements, as described above, are not sensitive to many factors such as

changing testing properties and the properties of the material. It is possible to speculate that

in this study, the GIca values, obtained from the ADCT, are not sensitive to the variation in

concentration of PS(NH2) chains at the interface, and therefore no trend in Gc was

observed. However, considering the fact that the measured GIc varied from batch to batch,

even though there is no trend observed within each batch, the test gives some indication

about the interfacial fracture energy. The large scatter in Gc values is typical for Gca

measurements. The possibility that the scatter in the data is due to the surface

contaminations from de-ionised water during the Si surface preparation (prior to the

monolayer deposition) is highly unlikely, as the UVO removes any residual traces of

organics from the surface. The samples were immersed into the polymeric solutions

immediately after the UVO cleanse.

The difference in the crack propagation behaviour between the PS and PS(NH2)

samples was obvious even though the ADCT was done by hand. As mentioned before, for

PS modified Si wafers, the crack propagated between the polymer layer and the epoxy very

easily in a stable continuous manner. The re-initiation of the crack was very easy. In

comparison, to re-initiate the crack in the PS(NH2) samples was difficult and often led to

breaking of the silicon.

The above observations led to the conclusion that the lack of variation in the

interfacial fracture toughness is more likely caused by the behaviour of the polymer chains

on the silicon surface. It was shown that the PS chains remained on the silicon surface even

after exposure to acidic and basic environments (Section 3.6.3), whereas PS(NH2) chains

were washed off during the step of converting –NH3+Cl- functional groups to -NH2 (Tables

III-13 to III-15). The amount of chains washed off varied randomly from batch to batch,

and therefore Gc results can not be related to the surface concentration of PS(NH2) chains.

The only conclusion that can be drawn at this stage is that the ADCT test measured Gca

values and that these values represent the fracture energy of the interfaces.

133

The main disadvantage of ADCT performed by hand, as used in this study, is that

the tensile/shear loading ratio (mixity) varies easily as the razor blade is pressed in. This

could be prevented by building a motor driven tool for performing the razor blade test,

similar to the one used for example by Smith4,7,21. This would allow for better control of the

speed and force applied, which is not possible when the test is done by hand. As a result the

mixity would be constant and the crack propagation cold be studied more objectively.

Comparing the results from the Experiments 1, 2 and 4 (see Section 4.8), the

correlation between the polar component of surface energy (γpolar) of the chemisorbed

PS/PS(Br) chains (before the conversion to PS/PS(NH2)) and the interfacial fracture energy

(Gc) between the epoxy and PS/PS(NH2) was found. Even though the variation of γpolar did

not reflect various concentrations of PS/PS(Br) chains in the original solutions, it varied

from experiment to experiment, suggesting different surface properties. The surfaces with

lower γpolar, around 3 mJ/m2, produced Gc ranging from 2 to 7 J/m2 (Experiment 2). The

surfaces with higher γpolar, around 10 mJ/m2, gave Gc ranging from 50 to 225 J/m2

(Experiment 1) and 70 to 95 J/m2 (Experiment 4). There might not be any direct connection

between γpolar and Gc but two possible hypotheses outlined below suggest some correlation:

1) After some polymer chains were washed off in the final conversion step,

there was still a sufficient amount of chains remaining on the surface with

-NH2 groups that reacted with the epoxy during cure. In such case, Gc would

reflect the variations in the interfacial chemistry. If γpolar corresponds to the

amount of the PS(Br) chains, it could be assumed that surfaces with a higher

γpolar had higher concentrations of PS(NH2) chains and therefore more links

across the Si/epoxy interface. Hence, Gc would increase with an increasing

γpolar due to the increasing concentration of the amino-functional polystyrene

chains.

2) As in scenario 1, the higher γpolar would suggest higher concentration of

PS(Br) and therefore PS(NH2) chains after the conversion. The hydrophilic

134

NH2 groups attracted water and therefore due to the hydrolysis of the

siloxane bonds, a larger amount of chains were washed off of the silicon

surface. This process could leave silicon wafers with only patches of

polymer. The areas of bare silicon would contribute to the strong interfacial

adhesion due to a good adhesion between epoxy and silicon oxide. In

addition, the increased surface roughness also contributes to stronger

adhesion.

The ellipsometry measurements showed that the final thickness of the polymer

chains was very similar for all three experiments, ranging from 5Å to 8Å. Such a small

variation in polymer thickness is within the instrumental error so it can be neglected.

Therefore the thickness measurements do not suggest differences in the structure of the

layers. Comparing the Experiments 1 and 2, the average thickness of the final polymer

layers was about 8Å. The γpolar and Gc values varied substantially however, suggesting that

the interfacial toughness is a measure of the amount of PS(NH2) between the silicon/epoxy

interface.

Further systematic studies including the characterisation of the deposited polymer

chains are necessary to confirm these observations. These will be outlined in the next

chapter.

135

5.2 References

1 K. Char, H. R. Brown, and V. R. Deline, Macromolecules 26, 4164-4171 (1993). 2 C. Creton, E. J. Kramer, C.-Y. Hui, and H. R. Brown, Macromolecules 25, 3075-

3088 (1992). 3 H. R. Brown, Macromolecules 24, 2752-2756 (1991). 4 J. W. Smith, E. J. Kramer, and P. J. Mills, J. Polymer Sci.:Part B: Polymer Physics

32, 1731-1744 (1994). 5 K. K. Weon and e. al., J. of Modern Physics B: Condensed Matter Physics,

Statistical Physics, Applied Physics 17, 1814-1820 (2003). 6 J. Washiyama, C. Creton, and E. J. Kramer, Macromolecules 26, 6011-6020 (1993). 7 J. W. Smith, E. J. Kramer, F. Xiao, C.-Y. Hui, W. Reichert, and H. R. Brown, J. of

Mat. Sci. 28, 4234-4244 (1993). 8 L. J. Norton, V. Smigolova, M. U. Pralle, A. Hubenko, K. H. Dai, E. J. Kramer, S.

Hahn, C. Berglund, and B. DeKoven, Macromolecules 28, 1999-2008 (1995). 9 Y. Sha, C.-Y. Hui, E. J. Kramer, S. F. Hahn, and C. A. Berglund, Macromolecules

29, 4728 (1996). 10 E. P. Plueddemann, Silane Coupling Agents (Plenum Press, 1982). 11 H. Yim, M. S. Kent, D. R. Tallant, M. J. Garcia, and J. Majewski, Langmuir 21,

4382-4392 (2005). 12 J. Duchet, J.-F. Gerard, J. P. Chapel, and B. Chabert, Composite Interfaces 8, 177-

187 (2001). 13 B. Arkles; Vol. 7 (Petrarch Systems, Inc., 1977). 14 J. Duchet, J. P. Chapel, B. Chabert, and J. F. Gerard, Macromolecules 31, 8264-

8272 (1998). 15 J. Duchet, B. Chabert, J. P. Chapel, J. F. Gerard, J. M. Chovelon, and N. Jaffrezic

Renault, Langmuir 13, 2271-2278 (1997). 16 C. P. Tripp and M. L. Hair, Langmuir 7, 923-927 (1991). 17 P. G. De Gennes, J. Phys. (Paris) 68, 1049 (1990). 18 H. R. Brown, Annu. Rev. Mater. Sci. 21, 463 (1991).

136

19 A. J. Kinloch and R. J. Young, Fracture Behaviour of Polymers (Applied Science

Publishers, Londong and New York, 1983). 20 J. Mijovic and J. A. Koutsky, Polymer 20, 1095-1107 (1979). 21 J. J. Benkoski, E. J. Kramer, H. Yim, M. S. Kent, and J. Hall, Langmuir 20, 3246-

3258 (2004).

CHAPTER 6

Conclusion

139

6.1 Conclusion

Various methods for the chemical modification of silicon surfaces by polymeric

coupling agents were investigated. These included grafting from (NMP and ATRP) and

grafting to techniques. Using the NMP technique, polystyrene (PS) chains were

successfully grown from silicon wafers. However, this method for growing copolymers

with various functionalities was found to be unsatisfactory. The copolymers obtained had

very short chain lengths, and the molar mass could not be satisfactorily controlled. ATRP

syntheses were successful for the preparation of free copolymers, however attempting to

grow polymers from the silicon surfaces did not produce molecules of sufficient length.

To overcome these difficulties, novel polymeric coupling agents were synthesised

by a chemical modification of the existing silane end-functional polystyrene chains with

low polydispersity. The PS-SiClMe2 chains were brominated (molar fraction of brominated

units, x = 0.1) and chemically grafted to the cleaned Si wafers as mixtures of PSBr-

SiClMe2/PS-SiClMe2 at various ratios ranging from 2 to 20% of brominated polystyrene.

The Br groups along the PS chains were then converted to NH2 groups by a chemical

procedure based on the Heck reaction. The conversion from PS to PS(Br) and subsequently

to PS(NH2) was carefully monitored by 13C-NMR and 1H-NMR techniques using free

(unattached) PS. The presence of the NH2 group was qualitatively confirmed by 1H-NMR,

which showed a peak around 3.4 ppm, corresponding to the amino group. In the final

conversion step of NH4+Cl- groups to NH2 groups, some polymer chains were washed off

the surface. It was concluded that because of the presence of hydrophilic NH2 groups, the

siloxane bond was hydrolysed. Because the polymer chains were swollen in a good solvent,

some of them were washed away.

The relationship between the interfacial fracture toughness and molecular coupling

was studied using the asymmetric double cantilever test (ADCT). The crack between the

epoxy and PS(NH2)/PS modified Si wafers was found to propagate is a stick/slip manner.

No significant variation of interfacial fracture toughness (Gc) with an increasing

140

concentration of the reactive PS(NH2) chains was observed. Some explanations for this

unexpected behaviour were suggested.

The ADCT was used previously by other researchers to study systems where

continuous crack propagation was more typical. Therefore, it was speculated that due to the

stick/slip behaviour observed, the Gc values measured were Gca (arrest) rather than Gci

(initiation). From the literature, it is known that GIca is less sensitive to the variations in

testing conditions (temperature, speed) than GIci therefore it was speculated that in this

study, the GIca values do not reflect the actual interfacial properties.

Another reason for not seeing any trend in Gc with an increasing concentration of

PS(NH2) chains within each experiment was concluded to be due to the washing off of the

chains caused by the hydrolysis of the siloxane bonds.

Even though there was no trend in the interfacial fracture toughness with increasing

concentration of amino-functional polystyrene chain concentration, Gc obtained by the

ADCT gives a clear indication of the presence of PS(NH2) chain links between the epoxy

and the silicon substrates because Gc of the samples with the plain PS chains was lower

than Gc of the samples with the mixed PS/PS(NH2) chains. In addition, a correlation

between the polar component of the surface energy (γpolar) of the initial PS/PS(Br) layers

and Gc was observed. It was hypothesised that even though these properties might not be

connected, the results demonstrate the increase of Gc with increasing γpolar.

141

6.2 Suggestions for Future Work

The relationship between the interfacial toughness of cross-linked polymers and

molecular coupling to a solid substrate is a field of great significance. In this work, more

light was shed on the topic, but additional research is needed in order to gain more

understanding of what is happening in the mixed polymer monolayers during the deposition

and the chemical modification procedures.

As described in Chapter 5, the characterisation of mixed polymer layers grafted to

Si surfaces is very challenging because of the low concentration of the amino groups in the

polymer monolayer. One possible solution could be to deposit a mixture of deuterated

polystyrene (dPS) and PS(Br), and analyse the resulting monolayers of dPS/PS(NH2) by

neutron reflectivity (NR), a method sensitive to the difference between deuterium and

hydrogen atoms. Also, the initial concentration of PS(Br) in a monolayer could be detected

by x-ray reflectivity (XR). This analytical method is sensitive to the difference in the

number of electrons, and therefore it would be sensitive to the presence of electron rich

bromine molecules. Both NR and XR along with attenuated total reflection infrared

spectroscopy, were used by other researchers in a study of hydrothermal degradation of

trimethoxysilane films deposited on Si wafers1. Nevertheless, because of the very low

concentrations of the functionalised polymer chains in the present study, the analysis would

be very challenging.

If the lower grafting density is the main reason for washing off of the chains, it

would be possible to achieve the higher σ by depositing the plain PS-SiCl2 on the silicon

surface first and than carrying out the bromination procedure. In order to carry out the

reaction quantitatively, free polymer would also have to be added to the bromination

solution in order to overcome the problem of a very low concentration of chemisorbed PS

chains. Because bromine is a relatively large molecule, the steric hindrance might cause a

difference between the degree of bromination of the free polymer and that of the

chemisorbed chains. The content of bromine in the chemisorbed polymer layer would

142

therefore have to be analysed. As mentioned above, X-ray reflectivity might be a suitable

technique for such analysis because of its sensitivity to electron rich bromine molecules. By

using this method of depositing PS chains first and then carrying out the bromination

procedure, all chains would be brominated. The only factor that could be varied in such

experiments is the degree of bromination of each polymer chain. In this study, however, the

concentration of brominated chains needed to be varied. Therefore, this approach would not

produce the surfaces required for the adhesion studies.

Also, growing functionalised chains using the “living” free polymerisation

techniques, including ATRP and reversible addition-fragmentation chain transfer (RAFT)

could be reassessed. The main advantage of these methods is that the functional groups are

incorporated into the growing polymer chains during the synthesis. This eliminates any

need for an additional chemical modification that might disrupt the siloxane bonds. The

Grafting from approach used in these methods usually yields polymer brushes of a higher

areal density and therefore the resulting monolayers are more resistant towards various

chemical environments. Some very recent progress was made in the field of molecular

brushes grafted from solid surfaces by ATRP2. The authors investigated the initiation

efficiency of n-butylacrylate brushes grown from poly(2-(2-bromopropionyl)oxyethyl

methacrylate) macroinitiator. They found that in comparison to the linear polymerisation of

the free polymer chains, the initiation in the grafting from polymerisation was not

quantitative at low conversion. The authors attributed this behaviour to the congested

environment. This could explain the lack of control reported in the present study, where the

molar mass of the polymer brushes grown by ATRP from the silicon surfaces was much

lower than the molar mass of the free polymer chains of the same composition. The authors

were able to enhance the initiation efficiency by increasing the rate of deactivation of the

growing species or decreasing the rate of propagation by increasing the concentration of the

catalyst (CuBr2) or by reducing the concentration of the monomer. It should be noted that

the polymer chains were grown from polymeric substrates and not from inorganic

substrates, as used in the present study.

143

Utilising the procedures from these recent studies1,2 would allow for growing the

mixed polymer layers from the silicon substrate in a controlled manner which would then

open the door for more systematic studies and understanding in the field of the interfacial

toughness between the cross-linked polymers and solid inorganic substrates.

Very recent references indicate that conditions under which surface initiated ATRP

are curried out are different to conditions for bulk ATRP3,4. At the time of writing these

thesis this information was not known.

6.3 References

1 H. Yim, M. S. Kent, D. R. Tallant, M. J. Garcia, and J. Majewski, Langmuir 21,

4382-4392 (2005). 2 J. J. Bekonski, E. J. Kramer, H. Yim, M. S. Kent, and J. Hall, Langmuir 20, 3246-

3258 (2004). 3 C. Xu, T. Wu, Y. Mei, C. M. Drain, J. D. Batteas, and K. L. Beers, Langmuir 21,

11136-11140 (2005). 4 J. Pietrasik, B. Cusick, T. Kowalewski, and K. Matyjaszewski, Polymer Preprints

46, 335-336 (2005).

Appendices

147

APPENDIX A: Universal calibration curve for a GPC column 1

1 M. P. Stevens, Polymer Chemistry (Oxford University Press, New York,

1990).

148

APPENDIX B: GPC and 1H-NMR results of PS and its copolymers

prepared by NMP

1H-NMR of PS synthesised by NMP using TEMPO

GPC of PS synthesised by NMP using TEMPO

149

1H-NMR of PHEMA/PS synthesised by NMP using TEMPO

GPC of PHEMA/PS synthesised by NMP using TEMPO

150

1H-NMR of PS/PS(NH2) synthesised by NMP using TEMPO

1H-NMR of PS/PS(NH2) synthesised by NMP using TEMPO D2O was added to an NMR solution to observe a disappearance of NH2 peak around 3.5 ppm.

1H-NMR of monomeric aminostyrene

151

GPC of PS/PS(NH2) synthesised by NMP using TEMPO

152

APPENDIX C: 1H-NMR spectra of PS(NH2)(Boc)/PS copolymer

1H-NMR of PS(NH2)(Boc)/PS synthesised by NMP using TEMPO

1H-NMR of synthesised Boc-Aminostyrene monomer

1H-NMR of commercial Aminostyrene monomer

153

APPENDIX D: Results for the copolymers prepared by ATRP

154

APPENDIX E: Schematics for synthesis of an active and an in-active

NMP initiators

Active Initiator

In-Active Initiator

ClCl

O

N

O

ON

HR2SiCl

H2PtCl6

O ON

SiR

ClR

NaH

OH

Alkoxyamine precursor

Surface-active alkoxyamine

NO+

TEMPO

Chloromethyl- styrene

Br

NaH

HO

O

HR2SiCl

H2PtCl6

O

Si R

Cl

R

relation between toughness and molecular coupling at cross

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