CHARACTERIZATION OF POLY(METHYL METHACRYLATE) AND

THERMOPLASTIC POLYURETHANE-CARBON NANOFIBER COMPOSITES

PRODUCED BY CHAOTIC MIXING

A Dissertation

Presented to the

The Graduate Faculty of the University

In Partial Fulfillment

of the Requirements for the Degree

Doctor of Philosophy

Guillermo A. Jimenez

May, 2007

CHARACTERIZATION OF POLY(METHYL METHACRYLATE) AND

THERMOPLASTIC POLYURETHANE-CARBON NANOFIBER COMPOSITES

PRODUCED BY CHAOTIC MIXING

Guillermo A. Jimenez

Dissertation

Approved: Accepted:

Rlawhdeorlawhdeorlawhdorlawheo rlawhdeorlawhdeorlawhdorlawheo Advisor Department Chair Dr. Sadhan C. Jana Dr. Sadhan C. Jana

Rlawhdeorlawhdeorlawhdorlawheo rlawhdeorlawhdeorlawhdorlawheo Committee Member Dean of the College Dr. Avraam I. Isayev Dr. Frank N. Kelley

Rlawhdeorlawhdeorlawhdorlawheo rlawhdeorlawhdeorlawhdorlawheo Committee Member Dean of the Graduate School Dr. Rex D. Ramsier Dr. George R. Newkome

Rlawhdeorlawhdeorlawhdorlawheo rlawhdeorlawhdeorlawhdorlawheo Committee Member Date Dr. Kevin A. Cavicchi

Rlawhdeorlawhdeorlawhdorlawheo Committee Member Dr. Shing-Chung Wong

ii ABSTRACT

Chaotic mixing is a novel mixing technique offering high mixing efficiency even under mild shearing conditions. In this work, chaotic mixing was used to prepare composites of carbon nanofibers and two thermoplastic polymers – poly (methyl methacrylate) (PMMA) and thermoplastic polyurethanes (TPU) – and their electrical, mechanical, and thermal properties were evaluated. The TPU systems were based on the reaction products of 4,4’-diphenylmethane diisocyanate, (MDI), soft segment polyol, and

1,4-butanediol as chain extender. Soft segment polyols in the form of poly(propylene glycol) (PPG), and poly(ε-caprolactone)diol (PCL) were used to obtain respectively amorphous and crystalline soft segments. Of these, the TPU system based on crystalline soft segment exhibited shape memory effects. Both, as-received untreated carbon nanofibers (CNF) with a very low amount of atomic oxygen on the surface, and oxidized carbon nanofibers (CNFOX) were used. CNFOX was also modified by esterifying with

PPG to produce a third type of carbon nanofiber named CNFOL. These carbon nanofibers were examined by X-ray photoelectron spectroscopy to determine the elemental composition of the surface, and by scanning electron microscopy and transmission electron microscopy to determine the surface morphology.

iii The as-received CNF and CNFOX fibers were highly entangled and showed

regions of amorphous carbon which acted as a deterrent to dispersion. The presence of

higher content of oxygen-containing functional groups in CNFOX improved their dispersion in PMMA and TPU, but produced electrical percolation at much higher concentrations except in the case of TPU based on crystalline soft segment. Thermal stability and thermo-oxidative properties of the polymers were improved in the presence of carbon nanofibers. It was found that the aspect ratio of carbon nanofibers was preserved and that carbon nanofibers were aligned along the flow direction during chaotic mixing. This led to electrical percolation at much lower nanofiber content, compared to materials produced in a commercial internal mixer. It was also found that crystallinity and phase separation in TPU materials in a chaotic mixer was affected much less by the presence of carbon nanofibers, while TPU materials produced in the commercial mixer showed more phase mixing. Carbon nanofibers interacted with both soft and hard phases in TPU. Composites with CNFOX showed higher crystallinity which resulted in high stress recovery, high fixity, and recovery ratio. Both CNF and CNFOX improved the tensile properties of TPU when measured above the melting temperature of PCL.

iv ACKNOWLEDGEMENTS

I would like to express my gratitude to my advisor, Dr. Sadhan C. Jana for his

guidance, and his support to complete this work. I would like to extend my thanks to my

committee, Dr. Avraam Isayev, Dr. Erol Sancaktar, Dr. Rex Ramsier, Dr. Shing-Chung

Wong, and Dr. Kevin Cavicchi for their advice. I am deeply thankful to all my group

members, past and present, for their suggestions and help throughout all my studies.

Financial assistantship from National Science Foundation in the form of

CAREER Award to Dr. Jana is gratefully acknowledged. I am deeply thankful with the

following sponsors: Fulbright Program, and LASPAU from the United States, and

Universidad Nacional, CONICIT, and MICIT from Costa Rica, for their financial support.

I would also like to thank my colleagues and supervisors at POLIUNA in Costa Rica, and the staff and faculty at the Department of Polymer Engineering, and the Office of

International Programs at the University of Akron for their support.

This dissertation is dedicated to those who gave me their ultimate support and patience: my wife, Giovanna, and my kids, Gabriel and Laura.

v TABLE OF CONTENTS

Page

LIST OF TABLES………………………………………………….….…………………x

LIST OF FIGURES…………………………………………………….………………..xii

CHAPTER I. INTRODUCTION ...... 1

II. LITERATURE REVIEW ...... 6

2.1. Chaotic mixing...... 7

2.2. Carbon nanofibers...... 10

2.2.1. Structure...... 11

2.2.2. Synthesis ...... 14

2.2.3. Properties ...... 14

2.2.4. Surface chemistry...... 16

2.2.5. Surface treatment ...... 19

2.2.6. Chemical functionalization ...... 22

2.3. Thermoplastic polyurethanes...... 24

2.3.1. Chemistry...... 24

2.3.2. Morphology...... 26

2.3.3. Shape memory polyurethanes...... 28

2.4. Polymer-carbon nanofiber composites ...... 30

2.4.1. General properties...... 31

vi 2.4.1.1. Tensile properties...... 32

2.4.1.2. Thermal management ...... 33

2.4.1.3. Electrical conductivity ...... 34

2.4.2. Preparation and characterization...... 38

2.4.2.1. General polymeric systems...... 38

2.4.2.2. Poly(methyl methacrylate)-carbon nanofiber composites ...... 43

2.4.2.3. Thermoplastic polyurethane-carbon nanofiber composites...... 45

2.4.2.4. Effect of fiber surface modification...... 47

2.4.2.5. Shape memory TPU nanocomposites...... 50

III. EXPERIMENTAL...... 55

3.1. Materials ...... 55

3.1.1. Poly (methyl methacrylate)...... 55

3.1.2. Thermoplastic polyurethanes...... 57

3.1.3. Carbon nanofibers...... 59

3.2. Composites preparation procedures...... 61

3.2.1. PMMA-carbon nanofiber composites...... 62

3.2.2. TPU-carbon nanofiber composites ...... 66

3.3. Characterization techniques...... 68

3.3.1. X-ray photoelectron spectroscopy ...... 68

3.3.2. Scanning electron microscopy ...... 68

3.3.3. Transmission electron microscopy and ultramicrotoming...... 69

3.3.4. Optical microscopy and image analysis...... 69

3.3.5. Electrical conductivity ...... 70 vii 3.3.6. Differential scanning calorimetry ...... 72

3.3.7. Thermogravimetrical analysis...... 72

3.3.8. Tensile test ...... 73

3.3.9. Dynamic mechanical analysis...... 73

3.3.10. Fourier-transform infrared spectroscopy ...... 73

3.3.11. Gel permeation chromatography...... 75

3.3.12. Shape memory properties ...... 75

IV. PREPARATION AND CHARACTERIZATION OF CARBON NANOFIBERS.. 78

4.1. Morphology of CNF and CNFOX ...... 78

4.2. Surface chemistry of CNF and CNFOX ...... 80

4.3. Preparation and characterization of CNFOL ...... 82

4.4. Summary...... 85

V. PMMA-CARBON NANOFIBER COMPOSITES ...... 86

5.1. CNF composition and processing technique...... 86

5.2. Effect of chaotic mixing time ...... 98

5.3. Effect of surface chemistry in CNF ...... 101

5.4. Summary...... 109

VI. TPU-CARBON NANOFIBER COMPOSITES...... 111

6.1. TPU23-carbon nanofiber composites ...... 111

6.1.1. Morphological analysis...... 112

6.1.2. Thermal degradation ...... 114

6.1.3. Molecular weight and molecular weight distribution ...... 116

6.1.4. Thermo-mechanical behavior ...... 116

viii 6.1.5. Electrical conductivity ...... 122

6.1.6. Tensile properties...... 122

6.1.7. Degree of phase separation ...... 124

6.1.8. Thermal transitions ...... 127

6.2. TPU33-carbon nanofiber composites ...... 130

6.2.1. Morphological analysis...... 130

6.2.2. Thermal degradation ...... 132

6.2.3. Thermo-mechanical behavior ...... 134

6.2.4. Electrical conductivity ...... 139

6.2.5. Tensile properties...... 141

6.2.6. Hydrogen bonding ...... 147

6.2.7. Thermal transitions ...... 152

6.3. Summary...... 160

VII. SHAPE MEMORY PROPERTIES IN TPU COMPOSITES ...... 162

7.1. Thermally-induced shape memory effect ...... 162

7.2. Shape memory effect induced by Joule heating ...... 165

7.3. Summary...... 168

VIII. OVERALL SUMMARY...... 169

REFERENCES ...... 172

ix LIST OF TABLES

Table Page

2-1 Properties of VGCFs (micro-, and nano-size) versus other fibers. (Adapted from Ref. 30)...... 15

3-1 Physical and mechanical properties of PMMA.148 ...... 57

3-2 Chemical composition of TPU systems...... 57

3-3 Physical and mechanical properties of CNF.74 ...... 59

3-4 Nomenclature of the composite systems...... 62

4-1 Oxygen/carbon ratio present on the carbon nanofibers...... 81

5-1 Thermo-oxidative degradation of PMMA-CNF composites prepared in the chaotic mixer...... 93

5-2 Elongation at break (%) of PMMA-CNF composites...... 96

5-3 Thermo-oxidative degradation of PMMA and composites prepared in chaotic mixer with different CNFOX contents...... 104

5-4 Bulk electrical conductivity of CNF and CNFOX...... 108

6-1 Storage, E’, and loss, E”, modulus at room temperature for TPU23 composites prepared in the chaotic mixer...... 120

6-2 Glass transition temperature of soft segments, Tg,SS from E” and tan δ plots of TPU23 composites prepared in the chaotic mixer...... 121

x 6-3 Tg values obtained from E” and tan δ plots for TPU33-CNF composites prepared in the chaotic mixer and Brabender Plasticorder...... 135

6-4 Ratio of the area under the peak of hydrogen-bonded NH to aliphatic CH, ANH/ACH of TPU33-CNF and TPU33-CNFOX composites (±0.05)...... 148

6-5 Effect of the temperature on the ratio of the area under the peak of hydrogen-bonded NH to aliphatic CH, ANH/ACH in neat TPU33 and TPU33-CNF7 (±0.05)...... 150

xi LIST OF FIGURES

Figure Page

2-1 Elliptic and hyperbolic points in blinking vortex flow...... 9

2-2 Dimensionality of carbon materials...... 11

2-3 Sketch of a single-walled carbon nanotube...... 12

2-4 Aspect ratio of several carbon materials. (Adapted from Ref. 31)...... 13

2-5 Diameter of several carbon materials. (Adapted from Ref. 31) ...... 13

2-6 Growth mechanism of a CNF. (Adapted from Ref. 31) ...... 14

2-7 Electrical resistivities of various forms of carbons compared to that of copper (HHT: heat treatment temperature). (Adapted from Ref. 32) ...... 16

2-8 Functional surface groups containing oxygen in carbon nanofibers. (Adapted from Ref. 34)...... 18

2-9 Reaction mechanism of the DCC-aided condensation of a polyol onto the surface of CNF...... 23

2-10 Molecular structures of most common industrial isocyanates...... 25

2-11 Chemical unit of a typical TPU hard segment...... 27

2-12 Shape deformation and recovery of a SMP...... 29

2-13 Dependence of the electrical conductivity on the filler content...... 35

3-1 Chemical structure of PMMA...... 56

xii 3-2 Chemical structure of PPG...... 58

3-3 Chemical structure of PCL diol...... 58

3-4 (a) Set-up of chaotic mixer 1 set-up and (b) chaotic mixing chamber. Scale bar is given to provide an idea about the dimension of the equipment...... 63

3-5 Mixing head of a Brabender Plasticorder. Scale bar is given to provide an idea about the dimension of the equipment...... 63

3-6 Diagrams showing geometry and dimensions of two mixing heads: (a)Chaotic mixer, and (b)Brabender Plasticorder mixer...... 64

3-7 Sketch of the setup to prepare a TPU prepolymer...... 66

3-8 Mixing chamber of chaotic mixer 2...... 67

3-9 Composite specimen and electrode set up for measurement of volume electrical conductivity in PMMA composites (a) 8-shaped specimen, (b) one half of 8-shaped specimen for conductivity measurement along flow direction, and (c) conductivity along the thickness direction...... 71

4-1 SEM images of (a) CNF and (b) CNFOX...... 79

4-2 Dispersion in water of CNF and CNFOX...... 79

4-3 XPS spectra of both types of carbon nanofibers...... 80

4-4 Narrow XPS spectra of the C(1s) region in (a) CNF, and (b) CNFOX...... 81

4-5 Weight loss for CNF, CNFOX and CNFOL...... 82

4-6 Full and narrow (inset) XPS spectra of CNFOL...... 83

4-7 Morphology of CNFOL as seen by (a) SEM, and (b) TEM...... 84

5-1 Volume electrical conductivity of PMMA-CNF composites prepared in chaotic mixer and Brabender Plasticorder...... 87

xiii 5-2 Optical micrographs of PMMA-CNF composites in the lower and upper limit of their electrical percolation threshold according the processing technique. Arrow indicates flow direction...... 89

5-3 Particle size distribution for PMMA-CNF composites prepared by two processing techniques...... 90

5-4 Fiber length distribution of fibers extracted from a PMMA composite with 4 wt. % of CNF prepared in chaotic and Brabender Plasti-corder...... 90

5-5 Optical micrograph showing conductive networks in 4 wt. % CNF composite prepared in the chaotic mixer. Arrow indicates flow direction...... 91

5-6 Thermo-oxidative stability of PMMA-CNF composites prepared by chaotic mixing...... 92

5-7 Storage modulus of PMMA-CNF composites prepared in (a) chaotic mixer, and (b) Brabender Plasticorder...... 94

5-8 Variation of tan δ with temperature for PMMA-CNF composites prepared in (a) chaotic mixer, and (b) Brabender Plasticorder...... 95

5-9 Maximum values of loss tangent (at glass transition temperature) of PMMA-CNF composites prepared in chaotic mixer, and Brabender Plasticorder...... 95

5-10 Tensile properties of PMMA-CNF composites (a) Stress at break, and (b) Young’s modulus...... 96

5-11 Effect of mixing time on the electrical volume conductivity along flow direction for composites prepared in the chaotic mixer...... 99

5-12 Effect of mixing time on morphological characteristics of PMMA-CNF 2 wt. % prepared in the chaotic mixer...... 100

5-13 Effect of mixing time on the storage modulus of PMMA-CNF 2 wt. % prepared in the chaotic mixer...... 100

5-14 Optical micrographs of PMMA 2 wt. % composites with (a) CNF and (b) CNFOX. Arrows indicate flow direction...... 101

xiv 5-15 Particle size distribution histogram for PMMA 2 wt.% composites with CNF and CNFOX...... 102

5-16 SEM images of the fractured surface of PMMA composites having 2 wt. % of (a) CNF and (b) CNFOX. Some nanofibers are circled for comparison...... 102

5-17 Thermo-oxidative stability of composites of PMMA-CNFOX...... 103

5-18 Dynamic mechanical properties of PMMA-CNFOX composites prepared by chaotic mixing. (a) Storage modulus and (b) tan δ...... 105

5-19 Dynamic mechanical properties of PMMA with 4 wt.% of CNF and CNFOX prepared in a chaotic mixer (a) Storage modulus, E’, and (b) tan δ...... 106

5-20 Volume electrical conductivity of PMMA-CNF and PMMA-CNFOX materials prepared by chaotic mixing...... 107

5-21 Sketch of the carbon nanofibers morphology (a) as-received, (b) after chaotic mixing, and (c) after Brabender plasticorder mixing...... 110

6-1 Mixing torque of TPU23 with different types of carbon nanofibers...... 112

6-2 TEM images of TPU23 composites with 3 wt. % of (a) CNF, (b) CNFOX, and (c) CNFOL. Arrows indicate the flow direction. Scale bars are 2 μm in (a) and (b) and 1 μm for (c)...... 113

6-3 TEM image of a cast film of TPU-CNFOX composite with 3 wt.% nanofibers...... 114

6-4 Thermogravimetric behavior in N2 gas of TPU composites with 0.5 wt. % of CNF, CNFOX and CNFOL (a) Mass loss (b) First derivative of mass loss...... 115

6-5 (a) Molecular weight and (b) molecular weight distribution of TPU23 and its corresponding 0.5 wt. % composites...... 116

6-6 Thermo-mechanical properties of TPU23-CNF composites: (a) storage modulus, (b) loss modulus and (c) tan δ...... 117

xv 6-7 Thermo-mechanical properties of TPU23-CNFOX composites: (a) storage modulus, (b) loss modulus and (c) tan δ...... 118

6-8 Thermo-mechanical properties of TPU23-CNFOL composites: (a) storage modulus, (b) loss modulus and (c) tan δ...... 119

6-9 Tensile properties TPU composites of CNF, CNFOX, and CNFOL (a) Stress at break (b) strain at break, and (c) Young’s modulus...... 123

6-10 Typical infrared spectra of a TPU23 carbon nanofiber composite with principal peak assignments (3 wt. % CNFOL)...... 125

6-11 Deconvolution of the carbonyl peak into free and hydrogen-bonded...... 125

6-12 (a) Hydrogen bonding index and (b) degree of phase separation of TPU23 composites...... 126

6-13 Thermal transitions of TPU composites...... 127

6-14 Sketch of the morphology of hard domains in TPU23 composites with surface treated CNF...... 129

6-15 Optical micrographs of TPU33 composites prepared in chaotic mixer with 1 wt. % of (a) CNF, and (b) CNFOX...... 131

6-16 SEM pictures of (a) TPU33-CNF5 prepared in chaotic mixer, (b) TPU33- CNFOX5 prepared in chaotic mixer and (c) TPU33-CNF5 prepared in Brabender Plasticorder...... 132

6-17 Thermal degradation stability of TPU33-CNF and TPU33-CNFOX composites (a) T1, and (b) T2 values...... 133

6-18 Thermo-mechanical properties of TPU33-CNF composites prepared in chaotic mixer: (a) storage modulus, (b) loss modulus and (c) tan δ...... 136

6-19 Thermo-mechanical properties of TPU33-CNF composites prepared in Brabender mixer: (a) storage modulus, (b) loss modulus and (c) tan δ...... 137

xvi 6-20 Thermo-mechanical properties of TPU33-CNFOX composites prepared in chaotic mixer: (a) storage modulus, (b) loss modulus and (c) tan δ...... 138

6-21 Electrical conductivity of TPU33-CNF composites prepared in (a) chaotic mixer and (b) Brabender Plasticorder and (c) TPU-CNFOX composites prepared in chaotic mixer...... 140

6-22 Tensile properties at room temperature of TPU33-CNF and TPU33- CNFOX composites: (a) stress at break, (b) strain at break, and (c) Young’s modulus...... 143

6-23 Typical stress-strain curves at 60 °C for (a) TPU33-CNF and (b) TPU33- CNFOX composites prepared in a chaotic mixer...... 145

6-24 Tensile properties at 60 °C for TPU33-CNF composites prepared in chaotic mixer and Brabender Plasticorder, and TPU33-CNFOX in chaotic mixer. (a) Maximum stress and (b) Young’s modulus...... 146

6-25 Typical FTIR spectra of TPU33-CNF5 prepared in Brabender Plasticorder with principal peak assignments...... 147

6-26 Values of AHCO/ACO in TPU33 composites according to (a) mixing method, and (b) fiber oxidation...... 149

6-27 FT-IR spectra of C=O stretching region at different temperatures for (a) pristine TPU33 and (b) TPU33-CNF 7 wt. % prepared in the chaotic mixer...... 151

6-28 Values of AHCO/ACO of neat TPU33 and its 7 wt. % composites at several temperatures...... 152

6-29 Full DSC heating scan of TPU33-CNF composites prepared ...... 153

6-30 (a) First and (b) second heating scan of specimens of TPU33-CNF composites mixed in a chaotic mixer...... 154

6-31 (a) First and (b) second heating scan of specimens of TPU33-CNF composites mixed in the Brabender mixer...... 155

6-32 Percentage of crystallinity of TPU33-CNF prepared by two processing techniques, (a) first and (b) second DSC scan...... 156 xvii 6-33 Glass transition temperature, Tg, of TPU33 materials synthesized in two different mixers (from first DSC scan)...... 156

6-34 (a) First and (b) second heating scan of specimens of TPU33-CNFOX composites mixed in a chaotic mixer...... 157

6-35 Percentage of crystallinity of TPU33-CNFOX prepared by chaotic mixing, (a) first and (b) second DSC scan...... 158

6-36 Heating and cooling DSC scan of TPU-CNF 5 wt. % prepared in (a) Brabender and (b) chaotic mixer...... 159

6-37 Crystallization temperature of TPU33-CNF composites...... 159

7-1 Percent of shape retention or fixity of TPU33-CNF and TPU33-CNFOX composites after 50 % strain at 60 °C...... 163

7-2 Stress recovery in (a) TPU33-CNF and (b) TPU33-CNFOX composites prepared in a chaotic mixer...... 164

7-3 Recovery ratio of TPU33-CNF and TPU33-CNFOX composites...... 165

7-4 Increase of temperature with applied voltage for (a) TPU33-CNF 5 and (b) TPU33-CNF 7 composites...... 166

7-5 Shape recovery of TPU33-CNF 5 triggered by Joule heating...... 167

7-6 Change with time of the deflection angle, θ, in a specimen bar of TPU- CNF 5. Straight lines are used to guide the eye...... 167

xviii CHAPTER I

INTRODUCTION

A composite can be defined as the combination of two or more different

materials.1 A subclass of these materials consists of a polymer matrix, and inorganic filler

with dimensions either in the micro-, or nano-scale. The embedded inorganic filler may impart properties not attainable with the matrix alone. In essence, a polymer matrix provides ductility, while the inorganic filler enhances the mechanical or thermal properties. One of the main concerns in composites is to limit the amounts of inorganic fillers to maintain a desired processability. The polymer melt viscosity invariably increases with the addition of fillers. In addition, many manufacturing related issues are encountered with the increase of filler content e.g., surface roughness, shear heating, and breakage of fillers. In some instances, a balance between two complementary properties, e.g., electrical vs. mechanical, must be maintained with the addition of single conductive

filler particles. An example might be conductive compounds such as carbon black, which

offer desired electrical conductivity, but most mechanical properties do not improve or

even deteriorate in some cases. This can be circumvented by using fibrillar conductive

particles, such as carbon nanofibers or carbon nanotubes.

1 In many instances, tensile modulus is dramatically increased by the reinforcing

filler, e.g., clay, talc, carbon and glass fibers, or organic ones like Kevlar, while

elongation at break might be sacrificed. Consequently, the composite industry struggles

with various trade-off issues.

Nanoscience and nanotechnology have flourished in the last decade because of

new methods of production of nanomaterials, and novel characterization and

manipulation techniques that are currently available.2 These nanomateriales form the

basis of a new class of composites that have been known as nanocomposites.

Nanocomposites have revolutionized the composite world since it was first observed that

by dispersing very small amounts of nanoclay, mechanical and barrier properties of nylon

6 can be significantly enhanced.3-5 Such improvements were attributed to the increase of

the interfacial area between polymer matrix and the filler when the clay platelets were individually dispersed.

The nanocomposite boom has also positively influenced the field of carbon

materials. Carbon black has been used for decades in the composite area. However, its

research with focus on nanotechnology obtained new impetus with the discovery of a

new allotrope of carbon, fullerenes in 1985.6 That discovery led to find another amazing

carbon structure, carbon nanotube in 1991.7 Carbon nanotubes are regarded as the ultimate carbon fiber. A single wall nanotube consists of a single layer of graphene rolled into a cylinder and capped with fullerenes. These discoveries promoted mass production of chemical vapor grown carbon materials such as carbon nanofibers, and carbon nanotubes.

2 General applications of carbon fibers can be found in fields like sporting goods,

aerospace, telecommunications, electronics, and structural reinforcement.8 Carbon nanotubes and carbon nanofibers have brought to the carbon fiber industry a new impetus.

The nano-scale dispersion of these nano-fibers in a polymer matrix may produce composite materials with outstanding properties never seen before in their micro- counterparts.

Despite such positive outlook and rapid development of understanding of the synthesis processes, dispersion of nanofibers or nanotubes is one of the major issues due to the bundling nature of these materials. There are several reasons why carbon nanofibers are hard to disperse and distribute into a polymer matrix - poor fiber to matrix interaction, presence of amorphous carbon, and polymer melt viscosity producing poor polymer infiltration into the carbon bundles. Conventional compounding methods normally use high shear force to produce conductive compounds. Dispersion quality in these processes may be excellent in some cases. However, fibers such as carbon nanofibers are susceptible to damage leading to reduction of aspect ratio. Shortening of these fibers is detrimental to complete realization of high modulus and high strength capabilities. On the other hand, a good dispersion is not always a necessity according to the final application of a composite. By controlling the level of dispersion quality, it may be possible to enhance either the mechanical or electrical properties. Conventional polymer mixing techniques are basically designed to maximize the level of dispersion of a filler in a polymeric matrix, not allowing structured or other levels of dispersion.

3 In this research project, we utilized chaotic advection as an alternative mixing

scheme to obtain polymer-carbon nanofiber composites. Chaotic stirring or mixing is a

low-shear mixing process, whose fundamentals were first laid down by Aref in 1984.9

Its mixing efficiency is derived from the rapid interface generation capabilities, thereby exposing fibers to the polymer matrix more efficiently. The rapid interface generation feature was capitalized in this study by promoting fiber-matrix interactions, especially, via functional groups chemically placed on the fiber surface. These functional groups produced favorable, energetic interactions with the polymeric matrix via generation of large interfacial areas aided by chaotic mixing.

The relationship between morphology and properties was evaluated in terms of how a structured dispersion is formed in the chaotic mixer. This structured dispersion is not possible in conventional mixing techniques, e.g. in internal mixer, in screw extrusion, or in injection molding.

The best mixing conditions to produce several polymer-carbon nanofiber composites in a chaotic mixer were determined in this study. An appropriate balance between well dispersed fibers and not so well-dispersed fibers was sought in order to produce multifunctional composites showing good electrical and mechanical properties.

Materials prepared in conventional processing equipment were evaluated together with those obtained by the chaotic mixing process. Such evaluations included thermal degradation, mechanical, electrical, and morphological characteristics. In addition, the surface chemistry of the fibers was assessed and its role on the morphological and final properties of the composites was evaluated.

4 This research was divided into two areas. In the first, the morphology developed in the polymer-carbon nanofibers systems produced in the chaotic mixer was investigated. In the second the effects of carbon nanofiber surface treatments on the extent of dispersion, electrical conductivity, shape memory, and mechanical properties of the composite materials were evaluated. In addition, composite materials were produced in a commercial mixing device under similar conditions in order to establish the potential advantages and limitations of the chaotic mixing technique.

5 CHAPTER II

LITERATURE REVIEW

Tuning the dispersion quality allows us to broaden the scope of applications of

polymer-carbon nanofiber composites. Chaotic mixing is a novel technique that can aid

such tuning with additional features such as self-similar morphology, alignment, and low-

shear mixing. In order to have control on the level of dispersion of the carbon nanofibers

in a polymer matrix, we need to know the nature of such nanofibers in terms of structure,

morphology, texture, and chemistry.

In this Chapter, we concentrate on basic aspects of structure, characterization, and surface chemistry of carbon nanofibers. Surface treatments and functionalization of

carbon nanofibers and carbon nanotubes are reviewed. Also some fundamentals on

chemistry, morphological aspects, and properties of thermoplastic polyurethanes are

provided with emphasis on shape-memory behavior. Most important preparation methods

of polymer composites of carbon nanofibers and carbon nanotubes are reviewed. Special

attention is placed on poly(methyl methacrylate) and thermoplastic polyurethane

6 composites. A basic understanding of the chaotic mixing approach is given, as well as its

usage in dealing with polymer composites especially those having carbon fillers is

discussed.

2.1. Chaotic mixing

The efficiency of mixing in conventional mixing devices can be improved using chaotic mixing principles. Aref first reported the fundamental connection between fluid flow and dynamical systems.9 In his theoretical paper, dynamical system was defined by

the advection equations of a passive marker particle in a prescribed flow. It was observed

that for the two-dimensional flow, the system was Hamiltonian, having only one degree

of freedom, while, for unsteady flow, the system was non-autonomous and in general

chaos was expected. He demonstrated that for steady velocity field, i.e. for

streamfunction being independent of time, the dynamical system is integrable and fluid

particle trajectories are non-chaotic, while for time-periodic velocity field, the particle

trajectories become chaotic and good mixing can be expected. Aref considered the action

of two blinking point vortices in his numerical mixing experiments and observed that

dynamical system became non-autonomous and produced chaotic particle trajectories due

to time periodicity. Subsequent to the research by Aref, a series of studies were

completed on model two-dimensional flows on the Lagrangian description of mixing and

transport to establish the viability of chaotic mixing. A historical account on the field of

chaotic mixing is available in a book written by Ottino.10

7 In chaotic advection, adjacent fluid elements separate from each other

exponentially with time. This has a tremendous potential in polymer mixing as the

domain of one polymer will be exponentially dispersed in another polymer to produce a

series of morphological signatures. The operating Reynolds number and shear rates in

low shear processing by chaotic mixing may be small. Nevertheless, the interfacial area

between the fluids can be created at exponential rates due to repeated alignment,

stretching, and folding of the interfaces by the action of some low order hyperbolic

periodic points as reported by a series of research studies by Ottino and co-workers.11-13

Repeated folding of the interfaces creates self-similar local microstructures, which are retained with the progress of mixing as finer and finer scales are added to them. This is contrary to turbulent mixing, where mixing occurs primarily through randomization of local microstructures. Also, turbulent mixing is much less energy efficient than chaotic mixing.14 Chaotic mixing possesses tremendous potential for applications in

polymerization, reactive functionalization, reactive and non-reactive compatibilization,

blending, and mixing of pigments, and fillers with matrix polymers.

Chaotic mixing is usually generated by time-dependent velocity fields where particle trajectories lose regularity that characterizes steady two-dimensional flows.

Therefore, better mixing is obtained under chaotic advection. A system can be classified as chaotic if it satisfies any of the following criteria:15

a. The flow produces either transverse homoclinic or transverse heteroclinic inter-

sections, or

b. the flow has positive Liapunov exponents, or

c. the flow is able to stretch and fold materials, thus producing horseshoe maps.

8 Figure 2-1 presents the typical form of the streamlines for the case of blinking vortex. It is seen that elliptic points in the neighborhood of the fluid just circulates around, while around hyperbolic points the fluid moves towards it in one direction and away from it in the other. A transverse homoclinic point is defined as the point where the inflow and outflow of a single hyperbolic point intersects. On the other hand, when the crossing takes place from flows of two different hyperbolic points, it is a transverse heteroclinic point. A positive Liapunov exponent indicates that adjacent fluid elements separate exponentially from each other, which is one of the most remarkable and attractive features of chaotic mixing. Horseshoe maps are formed when a flow is able to stretch and fold the fluid elements and return to its initial location.

Streamline

Elliptic Hyperbolic Elliptic

Figure 2-1 Elliptic and hyperbolic points in blinking vortex flow.

In chaotic mixing, even with low Reynolds numbers, the repeated alignment, stretching and folding generates interfacial area at exponential rates. This fact has been proved to enhance the rates of transport of mass and energy, as well as chemical reaction rates in non-polymeric systems.16 Additionally, chaotic mixing produces more ordered structures than random mixing, reducing the percolation concentration in conductive

9 blends.17 The principles of chaotic mixing have also been applied to improve the mixing

in single screw extruders.18,19

Chaotic flows are likely to generate self-similar mixing microstructures. This

feature has been exploited by Zumbrunnen and co-workers17,20-22 and Jana and co-

workers 23-27 in the production of conductive polymer composites 17,20,21,27 and lamellar and fibrillar morphological forms in immiscible polymer systems.22-26,28 Danescu and

Zumbrunnen used two-,21 and three-dimensional17 batch chaotic mixers to produce

conductive composites of polystyrene and observed electrical conductivity at much

reduced carbon black loadings. Similar observations were made in the mixing of pre-

concentrate of polyethylene with the host polymer in a continuous chaotic mixer.20 In all cases, conductivity of composites decreased with prolonged mixing, purportedly due to better dispersion of carbon black particles at long times, which in turn adversely affected conductive networks.17,20,21 Dharaiya et al.27 exploited the fibrillar morphology of

polypropylene phase to induce double percolation networks in a composite of 1 wt. % carbon black, and a blend of polypropylene with polyamide-6 produced by chaotic

mixing. Although these studies provided basic understanding of how chaotic mixing can

be used to produce conductive networks of carbon black in polymer blends, no studies

exist on chaotic mixing on high aspect ratio conductive fillers, such as carbon nanofibers.

2.2. Carbon nanofibers

Carbon nanofibers are vapor grown carbon fibers (VGCF) which can be

considered as carbon fibers (CF) with unique physical characteristics, and properties.

10 2.2.1. Structure

In order to have a clear idea about the structure of carbon nanofibers, a review of

the allotropes of carbon is presented below. Before 1985, only two allotropes of carbon

were acknowledged; graphite and diamond. Graphite shows a 2D hexagonal arrangement

of sp2 hybridization, while diamond is a 3D material (sp3 bonding) with isotropic

properties. Recently, two new forms of carbon joined graphite and diamond; fullerenes (a

zero-dimensional, 0D), and carbon nanotubes, and carbon nanofibers (1D forms).29

Figure 2-2 shows the dimensionality of the carbon materials mentioned above.

3D 2D 1D 0D Diamond Graphite Fibers and tubes Fullerene

Figure 2-2 Dimensionality of carbon materials.

A carbon nanotube (CNT) can be seen as a cylinder made up a graphene layer, capped by hemispheres of fullerenes (Figure 2-3). The curvature in the graphene layers increases the energy of the tubes per carbon atom. However, the lack of dangling bonds at the edges of the graphene layers lowers the total energy.

11 Rolling of the graphene sheet

Figure 2-3 Sketch of a single-walled carbon nanotube.

Carbon fibers on the other hand, represent an important class of graphite-related

materials. There are several precursors that can be used to synthesize carbon fibers, each

producing fibers with different morphologies. The mechanical strength of these fibers is

based on the fiber axes lying close to the in-plane direction of a graphene layer.30

Typical diameters for individual commercial fibers are ~10 μm. VGCF can be prepared over a wide range of diameters; from less than 1000 Å (nano fibers) to more than 100 μm (micro fibers). VGCFs have hollow cores similar to carbon nanotubes. A comparison of the aspect ratio and diameter among all the carbon fibers is given in

Figures 2-4 and 2-5 respectively. Carbon nanotubes have aspect ratios comparable to

VGCFs, while fullerenes have aspect ratios close to unity similar to carbon blacks.

Carbon nanotubes possess much smaller diameters than VGCF, and consequently showed a reduced quantity of graphene cylinders. A multi-walled nanotube for example, shows several concentric graphene tubules. 12 10000

8000

6000

4000 Aspect ratio (L/D) Aspect ratio 2000

0 Carbon black VGCF Carbon fibers

Figure 2-4 Aspect ratio of several carbon materials. (Adapted from Ref. 31)

10000

1000

100 Diameter (nm) 10

1 Fullerenes CNT VGCF Carbon fiber

Figure 2-5 Diameter of several carbon materials. (Adapted from Ref. 31)

Compared to conventional carbon microfibers, VGCF shows high purity graphene layers with a lack of cross-linking between them. This fact makes VGCF to have a lack of active sites on the fiber surface, which makes it resistant to oxidation, and less prone to bond to matrix materials.8

13 2.2.2. Synthesis

Vapor-grown carbon fibers are synthesized based on the decomposition of

hydrocarbons at temperatures ranging between 700 to 2500 °C. They may grow on substrates (fixed catalyst method), or without a substrate (floating catalyst method). The growth mechanism of VGCFs occurring via a fixed catalysis dehydrogenation reaction of

a hydrocarbon can be divided into several steps as presented in Figure 2-6. The diameter

of the fibers is related to the diameter of the catalyst particles. There are several mechanisms leading to VGCF growth resulting in different structural morphologies.

Growing fibers with the catalyst fixed normally yields micro and continuous fibers, while by using the floating catalyst method yields nano-size discontinuous filaments.31

Metal Metal Metal

Activation Initial Filament formation growth and thickening

Figure 2-6 Growth mechanism of a CNF. (Adapted from Ref. 31)

2.2.3. Properties

The physical properties of VGCF in some instances can approach those of single-

crystal graphite. Properties of vapor-grown carbon fibers produced by the fixed catalyst

and floating method are summarized in Table 2-1 together with a comparison with some 14 macro-fibers produced from other precursors.30 This table shows how approaching to the single-crystal graphite morphology, VGCFs can have the highest tensile properties and the lowest electrical resistivity among all the carbon fibers.

Table 2-1 Properties of VGCFs (micro-, and nano-size) versus other fibers. (Adapted from Ref. 30)

Isotropic PAN VGCF VGCF Property pitch-based fiber (micro) (nano) fiber

Diameter, μm 14.5 10 5-8 0.05

Density, g/cm3 1.57 2.0 2.0 2.1

Tensile strength, GPa 0.6 2.1 4 12

Young’s modulus, GPa 30 520 300 600

Resistivity, μΩ-cm 5000 500 50 20 (single fiber)

Single-fiber characterization of macro-fibers and micro-VGCF fibers is

commonly performed with direct experimental techniques. However, properties of nano-

VGCF produced by the floating catalyst technique are more difficult to determine due to

their smaller dimensions. In order to obtain an idea of their intrinsic properties such as

mechanical and electrical, the graphitic index of those nanofibers can be compared to that

of the microfibers, and thus their properties might be inferred.

15 The electrical conductivity of carbon materials is very sensitive to lattice

perfection; the higher the crystal perfection, the higher the conductivity. This feature can

be observed in Figure 2-7 which depicts a comparison in terms of electrical resistivity, ρ, of various carbonaceous materials against copper. It can also be observed that thermal treatment can have a significant effect on the electrical resistivity by improving the graphitic character of a carbon fiber.32

5.1E-05

m ) 4.6E-05 Ω 4.1E-05 3.6E-05 3.1E-05 2.6E-05 2.1E-05 1.6E-05 1.1E-05 5.5E-06 Electrical resistivity at 25 °C ( 5.0E-07 Copper Intercalated Graphite PAN VGCF graphite carbon (HTT) fibers (HTT)

Figure 2-7 Electrical resistivities of various forms of carbons compared to that of copper (HHT: heat treatment temperature). (Adapted from Ref. 32)

2.2.4. Surface chemistry

Carbon nanofibers are composed of 90 % by weight of elemental carbon and the rest is made various elements mainly oxygen and hydrogen. Different types of defects are present on the surface such as vacancies, dislocations, edges and steps.33 There are

mainly two approaches to address the surface characteristics of the carbon fibers. One is 16 the ‘solid state chemistry’ approach which considers the defects as ‘active sites’ on the

carbon fiber surface. The other approach is the ‘chemical functionality’ concept which deals with functional groups containing mainly oxygen and other atoms like nitrogen or

sulfur.

There are two categories of functional groups or surface oxides: acid and basic surface groups. Acidic and basic groups are simultaneously present on carbons, however the first ones are the most commonly found on the fiber surface. Among the most

important oxides with acidic character are the carboxyl and phenol groups. From these

two structures, a few other groups may be found at the carbon surface: anhydrides,

lactones, and lactols.34 Titration methods are generally used for the measurement of

acidic groups. Since the acidity constants of the various groups (carboxyls, phenols, and

lactones) differ by several orders of magnitude, an estimate of their relative amounts can

be obtained by titration with bases of different strength.35 Carbonyl groups are also

present on the carbon surface as isolated or conjugated structures like quinones.

Potentiometric techniques and cyclic voltammetry are suitable for the determination of

weak acidic groups and quinones. Figure 2-8 shows a scheme of typical functional groups

having oxygen in their structures that are possible to find on the surface of carbon

nanofibers.

Characterization of surface groups can also be carried out by X-ray photoelectron spectroscopy, XPS. This technique has been used extensively to characterize carbon materials, and in special carbon nanofibers which have been surface treated.34-38 From

XPS data, chemical information up to a 10 nm depth can be obtained when core electrons are excited by X-ray irradiation to leave the atoms.39 Depending on the nucleus attraction,

17 and the electron shielding (due to bonding to electronegative or electropositive atoms),

every element will produce a set of peaks in the photoelectron spectrum at kinetic

energies determined by the photon energy and the respective binding energies. The exact binding energy of an electron therefore depends not only upon the level from which the photoemission is occurring, but also upon the formal oxidation state of the atom and the local chemical and physical environment.

Anhydride O O O Carboxylic O

HO

Ether O O Lactone O O Quinone

O

O Carbonyl OH Hydroxyl

Figure 2-8 Functional surface groups containing oxygen in carbon nanofibers. (Adapted from Ref. 34)

Changes in the last two factors give rise to small shifts in the peak positions in the spectrum called chemical shifts. These chemical shifts energies can be quite small compared to the line width, thus making necessary to perform a deconvolution of the overlapping peaks. Normally carbon materials show two main peaks assigned to carbon and oxygen, C(1s) and O(1s) respectively. Most XPS studies on the surface chemistry of 18 carbon materials involve C(1s) analysis which shows one graphitic peak at 284.6 eV and four oxide peaks: 286.3 eV assigned to -C-OH or -C-O-C-; 287.7 eV assigned to -C=O;

289.4 eV corresponding to -COOH or -COOR, and 290.6 eV corresponding to –COO- and the π-π* shake-up satellite.35,36,38,40,41

The decomposition of the oxygen peak, O1s, at 528-536 eV in individual peaks is generally more difficult than for the C1s peak. XPS analysis of oxidized carbon fibers confirms that by moderate treatment mainly hydroxyl groups are formed and that after extensive oxidation carboxyl groups are also detected. The amount of oxygenated groups present on carbon fibers may be estimated from the ratio of the intensity of the O1s to the

C1s peak.

Other surface characterization techniques that have been applied to carbon fibers are Auger spectroscopy, secondary ion mass spectroscopy (SIMS), infrared spectroscopy

(IR), Raman spectroscopy and surface enhanced Raman spectroscopy (SERS), surface energy through wetting experiments, inverse gas chromatography, surface area and pore structure by gas or liquid adsorption, and scanning tunneling spectroscopy to obtain information on the surface roughness.39

2.2.5. Surface treatment

Unlike traditional CFs such as PAN-, and pitch-based fibers, the growth mechanism of VGCF also introduces lower concentrations of surface functional groups.42

VGCFs are produced by catalytic chemical vapor deposition that leads to two zones in the material with different aspects: an internal catalytic phase with a regular and oriented

19 structure and a pyrolytic phase characterized by irregular graphite layers and amorphous

carbon. In order to improve the surface chemical reactivity of the carbon fibers, surface treatment is essential. Surface treatment of carbon fibers is generally done in order to

increase the polymer melt infiltration into the fiber bundles, thereby increasing the level

of dispersion. This will result in enhancement of most of the mechanical properties. Also,

removal of the pyrolytic carbon is important to improve the electric and thermal conductivity of the composite. The effect of the surface treatments can be formation of functional groups and/or creation of steps, pits or defects in the surface.43

There are three main processes for modifying the oxygen surface groups on

carbon nanofibers: (i) thermal treatment, (ii) coating, and (iii) mild oxidation.

Thermal treatment may remove the surface groups on the fiber, and can improve

the perfection of the graphite lattice. Therefore, the surface fiber will become more

perfect and smooth, but more inert. Fibers after heat treatment normally show ultra high

modulus and strength and high electrical conductivity. Coating by means of chemical or

physical vapor deposition produces an enlargement of the fiber surface. Surface

treatments by mild oxidation in turn fall into the following categories: dry gaseous

oxidation, liquid oxidation, electrolytic oxidation, plasma treatment, ion or cluster

bombardment, and covalent linkage of biomolecules. Gas phase oxidation essentially

involves using air or oxygen diluted in an inert gas. Since a prerequisite of fiber treatment is to maintain its mechanical strength, overoxidation of carbon filament must be carefully avoided and the oxidation time must be strictly controlled. In the liquid phase oxidation, almost all oxidation agents currently used for carbon and graphites have been applied to carbon fibers and carbon nanofibers, especially nitric acid, sodium hypochlorite,

20 potassium permanganate and potassium dichromate.33 Electrolytic (anodic) etching is

well adapted for conductive materials like carbon filaments. It allows a continuous

treatment process with great flexibility in the operating conditions. In fact, anodic

oxidation is largely used for surface treatment of carbon filaments and yarns. Low-

pressure plasma treatments employ gases such as nitrogen, argon, oxygen, and air among

others. Plasmas contain highly energetic ions and radicals with very high average

temperatures. Therefore, chemically inactive basal planes of graphite may be

functionalized.

Among all the modification methods cited above, liquid and plasma oxidation are

the most used to alter the surface of carbon nanofibers. Nitric acid and oxygen plasma

treatments of VGCF can be used to increase the concentration of surface oxygen without

changing the morphology of these fibers. By performing an XPS survey of the O1s

region, Serp et al. 44 concluded that an oxygen plasma treatment might improve the adhesion of this type of fibers to polymeric matrices. Later on, Figueiredo et al. 45 using

Atomic Force Microscopy (AFM) determined that nitric oxidation of the VGCFs did not change the topography of the fiber. However, plasma treatment increased the size of the grains in the granular texture of the carbon nanofibers.

Several surface treatments, however, may introduce significant defects into the fibers, potentially affecting their intrinsic properties. Darmstadt et al.42 found that air-

oxidation of CNF affected not only its texture and chemical surface but also its graphitic

structure. Acid and plasma treatment have been found not to harm severely the graphitic

structure of CNF as reported by Serp et al.44 However, Toebes et al.46 recently concluded

21 that an increment of the specific surface area and pore volume upon acid treatment was

due to opening of the inner tubes in CNF.

2.2.6. Chemical functionalization

Functionalization allows for the segregation of entangled or bundled CNF for

their subsequent alignment. The surface modification of CNF plays an important role in

their use in composites, providing strong fiber-matrix bonding and thus improving the mechanical properties of the material.

Two main paths are usually followed for functionalization of CNF; attachment of

organic moieties to carboxylic groups that were previously formed by oxidation and

direct bonding to the surface double bonds. The covalent bonding can be realized via

chemical or electrochemical reactions. The chemical functionalization involves mainly

oxidation, fluorination, amidation, and esterification.47

Functional groups such as hydroxyl and carboxyl that have been produced after surface oxidation treatment can be used in carbon nanotubes or carbon nanofibers in order to covalently attach small molecules, oligomers or polymers.48 Two basic grafting approaches namely “grafting from” and “grafting to” are followed.49 In the “grafting

from”, a polymer builds up from a monomer that has been chemically attached to the

surface of a CNF or CNT. This method is basically a polymerization reaction occurring

on the fiber surface. In the “grafting to” approach, polymers with reactive end groups

react with the functional groups on the nanofiber or nanotube surface. This method

produces low grafting density due to the polymer chains diffusing to the fiber surface.

22 Therefore, the concentration of carboxyl groups on the surface is very low. This, together

with the fact of the low chemical reactivity of the carboxyl groups, makes it necessary to use chemical aids to speed up the reaction and to obtain maximum grafting density. One of those chemicals is thionyl chloride, SOCl2, which transform a carboxylic acid group

into a much more reactive acyl chloride. Poly(ethylene oxide) 50, poly(propionylethylenimine-co-ethylenimine)51, and polyurea 52 among many others

have been grafted onto multi-walled nanotubes (MWNT) using SOCl2. Few papers have

been published to date dealing with chemical derivatization of CNF. In one of those, Wei

et al. 53 grafted a poly(ethylene glycol) (PEG) onto the surface of VGCF by using N,N’-

dicyclohexylcarbodiimide (DCC) as a condensation agent. Mechanism for this reaction is

sketched in Figure 2-9. By using thermogravimetrical analysis, they determined that 11

mol % of grafting was equivalent to 1 mol % of COOH used in the reaction. However,

they observed that in the absence of DCC the grafting reaction does not occur appreciably.

O-Acylisourea H N O O O N N O H O H C N N N H DCC Dicyclohexylurea Oxidized carbon OH nanofiber having carboxylic acid Polyol groups on its surface O O OH

Figure 2-9 Reaction mechanism of the DCC-aided condensation of a polyol onto the surface of CNF.

23 2.3. Thermoplastic polyurethanes

Thermoplastic polyurethanes are versatile and to some extent complex materials in that their chemistry, structure and morphology can be tweaked to obtain the desired final properties. Therefore, a detailed review of their characteristics is presented below.

2.3.1. Chemistry

Thermoplastic polyurethanes (TPU) are based on the exothermic reaction of polyisocyanates with polyol molecules, containing hydroxyl groups. Relatively few basic isocyanates and a range of polyols of different molecular weights and functionalities are used to produce the whole spectrum of polyurethane materials. The chemically efficient polymer reaction may be catalyzed, allowing extremely fast cycle times and large production rates.54 Two types of diisocyanates are employed in polyurethane preparations, aromatic and aliphatic ones. Figure 2-10 shows chemical structures of the most commonly used aromatic diisocyanates; toluene diisocyanate (TDI), and 4,4’- diphenylmethane diisocyanate (MDI).55 TPU are usually made from pure MDI which is reacted with a substantially linear polyether or polyester diol and with a chain-extending diol of a low molecular weight, such as 1,4-butanediol, (BDO) in either a one-step or a two-step reaction process.56 The polyester diols are usually the condensation products of adipic acid and one or more simple aliphatic diols ranging from ethylene glycol to 1,6- hexanediol.

24 CH3 CH3

NCO OCN NCO

NCO

2,4-TDI isomer 2,6-TDI isomer

OCN NCO

MDI

Figure 2-10 Molecular structures of most common industrial isocyanates.

To obtain improved low-temperature performance with relatively high resistance to hydrolysis, 6-hydroxycaproic acid polyesters, made by the polymerization of ε- caprolactone are also used.

In this specific case, poly(ε-caprolactone) (PCL) – based polyols are produced.

Polyether TPU is usually based upon polyethyleneglycols, polypropyleneglycols, polytetramethyleneglycols, or polytetrahydrofurans.

The reaction between isocyanates and polyols is accelerated by the addition of catalysts such as acids, bases (mostly aliphatic tertiary amines) and metal complexes

(organo tin compounds).56 Two general types of organotins are used as polyurethane catalysts, tin II (stannous) and tin IV (stannic). The major stannous compound used is stannous 2-ethylhexanoate, more commonly referred to as stannous octanoate. The main tin IV compounds used are dialkyltin dicarboxylates or dialkyltin mercaptides. The

25 proposed mechanism for tin IV catalysts, dialkyltin dicarbonates and dialkyltin

dialkylthiolates, is the reaction of the tin with a polyol forming a tin alkoxide, which can

then react with the isocyanate to form a complex. Transfer of the alkoxide anion onto the

co-ordinated isocyanate affords an N-stannylurethane, which then undergoes alcoholysis

to produce the urethane group and the original tin alkoxide

2.3.2. Morphology

Thermoplastic polyurethanes have a combination of high elongation and tensile

strength and Young’s modulus, and so form a bridge between rubbery polymers and

thermoplastics and in addition their toughness provides excellent abrasion and tear

resistance.

The polyurethanes are segmented or block copolymers with an (AB)n structure,

where A represents the hard block with a degree of polymerization, (number of repeat

units, DP) of one to 10, and B represents the soft block which has a degree of

polymerization in the range 15 to 30. They consist of alternating hard segments (which

may be glassy or semicrystalline) and soft (elastomeric) chain segments. The microphase

separation of these two chemically distinct components gives rise to the unusual and

useful physical and mechanical properties of the polyurethanes.57 The resulting hard block domains act as cross-linked, physically-bonded crystalline centers dispersed in the soft segment domain of flexible polyether or polyester chains. The hard block domains, therefore, act as molecular reinforcing-fillers. The degree of hard segment aggregation or domain formation depends not only upon the weight ratio of the polyurethane hard

26 segment to the polyether or polyester soft segment, but also on the choice of glycol, the

type and the molecular weight of the polyester or polyether diol, and also upon the manufacturing process and the reaction conditions. Hard segments from MDI and linear aliphatic glycols have the general structure depicted in Figure 2-11.

Segmented thermoplastic elastomers exhibit structural heterogeneity on the molecular domain, and in some cases, on a large scale involving spherulitic texture.

O O

-O-C-NH- CH2 NH-C-O-(CH2)n-

X

Figure 2-11 Chemical unit of a typical TPU hard segment.

Several morphological models have been published in the past in order to explain most of the features observed on this type of materials. An early model by Estes58 considered the hard-segment domains as an interconnecting network. Both phases were considered to be continuous and interpenetrating. Early X-ray diffraction studies performed by Bonart et al. 59 identified short-range order associated with hydrogen

bonding. Wilkes and Yusek 60 studied domain formation in polyesterurethanes using X- ray techniques. They found that the domains were lamellar in shape with an average separation of 100 to 250 Å. They concluded that the hard segments act as crosslinks, inhibiting stress-relaxation and inducing stress-crystallization of the soft segments which results in higher tensile strength. Finally, Blackwell et al. 61,62 were able to interpret wide

27 angle X-ray data of heat-set polyurethanes which defined the chain conformation and

packing of hard segments in crystalline polyurethane elastomers. The models were based

on the structure of MDI-butanediol hard-segment analogs with chain packing. Planar zig-

zag –CH2-CH2- sections connect successive diisocyanate units. The chains are linked

together in stacks through C=O---H-N hydrogen bonds which involve half of the urethane

groups.

2.3.3. Shape memory polyurethanes

Shape-memory polymers (SMP) are stimuli-responsive. They have the capability

of changing their shape upon application of an external stimulus. A change in shape can

be caused by increasing the temperature of the surroundings around the material or by

increasing locally the temperature by generating heat with the application of an electrical

voltage. The shape-memory effect is not related to a specific material property of single

polymers; it rather results from a combination of the polymer structure and the polymer

morphology together with the applied processing and deformation.63 The deformation and recovery of a shape is shown schematically in Figure 2-12. First, the polymer is conventionally processed to receive its permanent shape. Then, the polymer is deformed and the temporary shape is fixed. The original shape is then recovered by application of heat.

Segmented polyurethanes with shape memory behavior were discovered at

Mitsubishi Co. in 1988.63,64 However, it was since 1996 that their investigation turned out

to be more systematic.63-72

28 Permanent shape Temporary shape Permanent shape

Deformation Recovery

Heat Heat or electricity

Figure 2-12 Shape deformation and recovery of a SMP.

For this type of TPU, the hard segments act as cross-links and are responsible for

the permanent shape. The polymer can be processed above the melting temperature of the

hard segment domains. The soft segment phase serves as a molecular switch and enables

the fixation of the temporary shape. The transition temperature for the fixation of the

switching segments can either be the glass transition temperature, Tg, or the melting

temperature, Tm. Soft blocks for Tg-controlled shape recovery that have been used are

poly(tetramethylene oxide) glycol69,70, poly(ethylene adipate) 66, and

63 poly(tetrahydrofuran) among others. Soft blocks used to provide Tm-control on the

shape memory behavior are mostly based on poly(ε-caprolactone) (PCL) 65,67,72,73. The

most common system for this case is PCL/MDI/BDO in which PCL diols with number-

average molecular weight (Mn) between 2000 and 8000 form the switching segments.

The switching temperature for the shape-memory effect can oscillate from 44 to 55 °C

depending on the composition of the soft phase (between 50 and 90 wt. %) and the

molecular weight of the PCL diols. The crystallization of PCL in TPU is restricted

because of the connectivity between the soft and the hard phase. As a matter of fact, the 29 values for PCL crystallinity may vary between 10 and 40 % when compared to 100 %

PCL material.67

2.4. Polymer-carbon nanofiber composites

Reinforcement of polymers is one of the most important applications of all types

of carbon fibers. Thermosetting resins and thermoplastic polymers are commonly used as

matrices. Epoxy resins have been the primary polymeric matrix material used in carbon-

based polymer composites.

A prerequisite for a good polymer-carbon nanofiber composite is to have an

adequate interfacial adhesion between the inorganic and the organic material. This occurs

when there is a good wetting of the nanofiber by the molten thermoplastic polymer or the

liquid precursor of a thermosetting resin. The extent of interfacial contact depends on the

wetting behavior (contact angle and viscosity) of the fibrous composite. Consequently

physical bonding between fiber and matrix is very critical for producing a high

performance composite.33

Most of the composites have been prepared using VGCF produced by the fixed

catalyst method. This form of fiber can grow to have a diameter in the range of PAN and pitch-derived carbon fibers. Moreover, it can be oriented and compressed into a mold,

with fiber volumes comparable to composites reinforced with PAN and pitch-derived

fibers. Less research work on organic matrix composites reinforced with VGCF based on

the floating catalyst method has been made.18

30 One of the most important companies in United States related to the production

and commercialization of VGCF is Applied Sciences, Inc, (ASI). In the past few years,

ASI has used the fixed catalyst method to fabricate long VGCF (known as Pyrograf I®)

This type of fiber can be used in composites for thermal management, high power electronic devices, space power system radiator fins, and high performance applications such as plasma facing components in experimental nuclear fusion reactors. These composites include carbon-carbon composites, polymer matrix composites, and metal matrix composites.

Recently, ASI started marketing short nano-VGCF, Pyrograf III®, produced by

the floating catalyst method. It is produced in the vapor phase by decomposing either

methane, ethane, other aliphatic hydrocarbons, or coal gas in the presence of a metal

catalyst, hydrogen sulfide and ammonia. The additional processing can include pyrolytic

stripping to remove tars and other hydrocarbons from the surface of the fiber.74

2.4.1. General properties

The mechanical and the electrical properties of the polymeric matrices may be

enhanced by the incorporation of carbon nanofibers. Thermal conductivity and thermal

degradation of polymers also are properties that can be modified with the addition the

graphitic materials. These properties will be reviewed in the next sections.

31 2.4.1.1. Tensile properties

A major stimulus for the development of any low-cost carbon fibers is for their

potential applications in the automotive industry. A very high degree of graphitic

perfection in the fibers and by inference, a high modulus of elasticity has been

determined by x-ray diffraction for selected preparations of floating catalyst VGCF even

without subjecting the fiber to any post-growth heat treatment. Based on the presumed

high modulus, VGCF can be used to produce thermoplastic- and thermoset-matrix composites with elastic moduli comparable or exceeding that of aluminum, provided that preferential orientation in two dimensions can be obtained.

The tensile strength of an ideal composite (all fibers are aligned in one direction and submitted to an uniaxial tension), σC, can be calculated knowing the fiber strength, σf, the strength of the matrix at ultimate fiber deformation, σmε, and the volume fraction of

1 the fiber, Vf, by using the rule of mixture:

σ C = σ f V f +σ mε (1−V f ) (1)

The premise that discontinuous short fibers such as floating catalyst VGCF can

provide structural reinforcements can be supported by the Cox model 75, and later

76 extended by Baxter. This theoretical model predicts that modulus of a composite, Ec,

can be determined from the fiber and matrix moduli, Ef and Em respectively, and the fiber

volume fraction, Vf, by a variation of the rule of mixtures:

32 ⎛ l ⎞ EC = EmVm + E f V f g⎜ ⎟ f (θ ) (2) ⎝ d ⎠

where Vm is the matrix volume fraction, and functions, f and g, can take values between 0 and 1. The function g is small for particles having a low aspect ratio, but increases rapidly as the aspect ratio increases. The function f is dependent upon orientation of the fibers, θ, and is greatest for uniaxial alignment. These findings imply that if floating catalyst fibers –which have a very high aspect ratio – can be restricted in orientation in two dimensions, the resulting composite could be several times stiffer than glass-reinforced composites.

2.4.1.2. Thermal management

A significant portion of the development work conducted on VGCF composites has been motivated by the potential of these composites for high performance thermal management applications, such as electronic heat sinks, plasma facing materials, and radiator fins. Both the fixed catalyst and the floating catalyst VGCF have the potential to be economically important for thermal management or higher temperature composites.

Composites fabricated with fixed catalyst VGCF can be designed with fibers oriented in preferred directions to produce desired combinations of thermal conductivity and coefficient of thermal expansion. Composites prepared with the smaller floating catalyst fiber are most likely to be used for applications where near-isotropic orientation is favored. Such isotropic properties would be acceptable in carbon/carbon composites

33 for pistons, brake pads, and heat sink applications and the low cost of fiber synthesis could permit these price-sensitive applications to be developed economically. A random orientation of fibers will give a balance of thermal properties in all axes, which can be important in brake and electronic heat sink applications.

2.4.1.3. Electrical conductivity

The electrical conductivity of a composite is generally characterized by its dependence on the filler volume fraction. At low filler loadings, the conductivity of the composite is very close to that of the pure, electrically insulating polymer matrix. At some critical loading (i.e., percolation threshold), the conductivity increases by several orders of magnitude with very little increase in the filler amount. After this region of drastic increase, the conductivity levels off and approaches that of the filler material. It is at the percolation threshold that enough filler has been added so that it begins to form a continuous conductive network through the composite (Figure 2-13).77

There are various models to predict the electrical conductivity behavior of composites based on numerous factors. Physical properties of both the filler and the polymer will influence the composite and include structural properties, interfacial properties, and constituent conductivity. Properties of the filler that play a role in determining the conductivity of the composite are the filler type, shape, and orientation within the matrix.

Different forms of carbon generally have different microstructures and, therefore, will affect electrical conductivity in different ways. For spherical particles, smaller particles lower the percolation threshold. For fillers with an aspect ratio greater than unity, larger

34 aspect ratios and a broader range of them can lower the percolation threshold.77 Most of the models are of the statistical percolation type. These models predict the conductivity based on the probability of particle contacts within the composite.

One of basic statistical models follows a power-law equation as follows,

S σ = σ 0 (V −VC ) (3)

where σ is the conductivity of the mixture; σ0 is the conductivity of the filler; V is the volume fraction of the filler; VC is the percolation volume fraction; and the critical exponent, s, dependent on the dimension of the lattice. More accurate statistical models include polymer gelation, and a general effective media equation which considers resistivities of both components.

0

-4

Percolation -8 concentration

-12

Log Conductivity (S/cm) -16

-20 Vc Carbon Fiber Volume Fraction

Figure 2-13 Dependence of the electrical conductivity on the filler content.

35 Thermodynamic models consider factors like filler and polymer surface energies and polymer melt viscosity among others. At all points above the percolation threshold, the conductivity of a composite has been found to be, 77

k ⎛ φ − φ ⎞ ⎜ C ⎟ logσ = logσ C + (logσ m − logσ C )⎜ ⎟ (4) ⎝ F − φC ⎠ and,

KφC k = 0.75 ; K = A − Bγ pf (5) ()φ − φC

where σC is the conductivity at the percolation threshold; σm, the conductivity at F, the maximum packing fraction; φ, the volume fraction; (l/d), the aspect ratio; φc, the percolation threshold; γpf , the interfacial tension; and A and B, constants. The value k is dependent upon the filler volume fraction, percolation threshold, and interfacial tension as calculated by,77

0.5 γ pf = γ P + γ f − 2(γ pγ f ) (6)

where γpf is the interfacial tension; γp, the surface energy of the polymer; and γf, the surface energy of the filler. This model works for polymers filled only with carbon black. To apply this model for carbon fibers, the maximum packing fraction, F, can be used as,77

36 5 F = (7) 75 + ()l / d 10 + ()l / d

The geometrical percolation model, on the other hand, was intended to predict the conductivity of sintered mixtures of conducting and insulating powders. The main parameters used in determining the conductivity are the diameters of the nonsintered particles or the edge length of the sintered particles. Equations in this model use the diameter of the particles, the probability for the occurrence of long bands of conductive particles, and the arrangement of the conductive particles on the surface of the insulating particles.

Finally, the structure-oriented models are based on the physical construction of the final composite. The electrical conductivity of composite materials is often affected by structural properties such as the filler aspect ratio and filler orientation. These properties are typically a result of the processing techniques employed to make the composite. For example, injection molding a composite will cause an alignment of fillers within a polymer due the flow through the nozzle and the mold. Alignment of the fillers can result in different conductivity results depending on the direction of measurement.

Extrusion and injection-molding processes can also degrade fibers, thereby shortening their lengths.

37 2.4.2. Preparation and characterization

The dispersion of highly entangled materials such as CNF or CNT in polymeric viscous materials is a critical deterrent in transferring all the qualities of the carbon materials to the organic polymer. However, a perfect CNF or CNT dispersion it is not a requirement for all the applications. The level of dispersion can be modulated by the preparation method, and the interface between the polymer and CNF/CNT. Following, we review the principal routes of fabrication of polymer-CNF and polymer-CNT composites.

2.4.2.1. General polymeric systems

One of the first attempts to use the sub-micron floating catalyst form of VGCF to prepare polymer composites was done by Dasch and collaborators.78 They reported the fabrication of thermoplastic composites reinforced with randomly oriented VGCF, up to

30 % of volume fraction, having diameters of 0.08 μm and lengths of 2.5 μm. All the composites exhibited similar flexural strength near 70 MPa, in accordance with Baxter’s theory 76 for 3D short fiber reinforced composites. Also, flexure modulus increased with fiber volume fraction in agreement with calculations based on Cox’s theory for random

3D short fiber reinforcements.8

Discontinuous carbon nanofibers were also combined with continuous carbon microfibers in order to produce epoxy composites.79 In this case, flexural properties were evaluated by Dynamic Mechanical Analysis (DMA). An improvement in tan δ was

38 observed with the addition of discontinuous nanofibers (0.6 vol. %, and diameter of 0.1-

0.2 μm) to an epoxy composite having already 56.5 vol. % of continuous carbon fibers

(diameter of 7 μm).

Carneiro et al. 80 prepared polycarbonate (PC) and VGCF composites by co- rotating twin-screw extruder and then injection molded to evaluate their properties.

Tensile properties of those nanofiber composites were very similar to those PC composites containing carbon black (CB). Moreover, the PC-CNF composites showed a decreasing tendency of their impact properties with CNF content. This fact was attributed to the lack of good adhesion between the untreated fiber and the polymer matrix. They also used oxygen plasma treated CNFs, with an oxygen concentration of 12.6% (in contrast to a 1% for the untreated carbon nanofibers). PC composites made with plasma- treated CNFs did not have great influence on the tensile behavior. Although impact properties improved, still they were below the performance of the resin. The authors concluded in this case that the difference in surface energy among PC and CNFs played an important role on the results observed. Recently Higgins and Brittain81 reported the preparation of PC-CNF composites by in situ polymerization method. They estimated a 9 wt. % in CNF concentration for the percolation threshold. These authors attributed this to the much less electrical conductivity of CNF when compared to CNT. They also observed less uniform dispersion in their composites compared to high-shear methods.

Tibbets 82 obtained low values of tensile strength with as-received and untreated

CNFs combined separately with nylon 6,6 and polypropylene (PP). It was argued that such poor tensile properties were due to the presence of large clumps of fibers (about 500

μm) originated from the reactor. In order to reduce the size of the clumps, the author 39 employed ball milling on the fibers before mixing them with the thermoplastics into a mini injection molding apparatus. By ball-milling, the initial fiber clumps were reduced, and the polymer composites showed an improvement in the mechanical properties.

Additional surface etching of the ball milled-fibers helped in the mechanical properties, unlike the clumped nanofibers. However, they also observed that the fibers aspect ratio reduced dramatically by ball milling.

The difficulty of infiltration of a viscous polymer matrix into carbon nanofiber clumps can be explained using previous models of molten metals being mixed with fibrous materials. The pressure P required to infiltrate a porous material or a clump of fibers, is due to the pressure to overcome capillary effects, Pγ, and the pressure required to overcome viscous drag, Pν:

P = Pγ + Pν (11)

The effective capillary pressure is given by,

Pγ = -Sf εLA cosθ (12)

where Sf is the surface area of the interface of the clump per unit volume of composite, εLA is the surface energy of the polymer, and θ is the contact angle. Once this pressure is supplied, the gradient dPV /dx generated across the clump by the polymer at velocity υ will be given by Darcy’s law in one-dimensional form:

40 dP μυ V = − (13) dx K

where μ is the viscosity of the polymer, and K is the permeability of the clump.

Solving this equation, the infiltration depth D is,

2KPτ D = 0 (14) μ(1−V )

where τ is the duration of the infiltration at an applied pressure P0, and fiber volume V.

For fibrous materials, the permeability K can be represented by the following equation,82

2r 2 2 ⎛ 4V ⎞ ⎜ ⎟ K = ⎜1− ⎟ (15) 9V ⎝ π ⎠

The dependence on r2 means that K for submicron fibers is only 10-4 times the value for conventional 10 μm-diameter fibers. This extremely small value for K leads to very slow infiltration rates in nanofibers.

Lozano and Barrera 83 combined polypropylene (PP) with VGCF in a Banbury- type mixer, and determined thermal and mechanical properties of the composites. One of the main goals in this paper was to produce uniformly dispersed nanofibers in a polymer matrix with the absence of agglomerates and porosity. They used Pyrograf III® carbon

41 nanofibers without and with purification. The purification step was conceived to remove non-nanofiber material (amorphous carbon for example) while opening up the nanofiber network for easy deagglomeration of the nanofibers. They observed a remarkable improvement in the final decomposition temperature analyzed by Thermal Gravimetric

Analysis (TGA) for the composite having 30 wt. % of CNF. This behavior was due to restrictions on the mobility of the polymer chains due to the nano-sized VGCF. From the

DMA data, the storage modulus, E’ (at room temperature), of the composites was superior to the resin alone. However, the E’ difference between 2 and 20 wt.% composites was not as drastic as the change in fiber composition. Furthermore, the high temperature E’ only started to increase significantly when 60 wt. % PP-CNF composite was analyzed. Tensile properties of the composites did not show any improvement when compared with that of PP.

In a subsequent paper, Lozano and collaborators 84 focused on rheological and electrical characterization. To prepare the composites, they employed similar method as in their first paper. The high shear nature of the Banbury mixer, led to composites with a homogeneous dispersion of the nanofibers, and an electrical percolation threshold between 9 and 18 wt. %. They observed that at low concentration of CNFs, a good dispersion of the fibers was obtained, while at higher carbon nanofiber contents, the presence of agglomerates becomes more important. It appeared that as the concentration of VGCF was increased, there was a composition at which fiber interactions tended to alter their distribution toward more interconnected nanofiber dispersion. By being distributed, these aggregates and agglomerated phases can form a conductive three- dimensional network throughout the whole sample, as it was noted by Sandler et al.85 in

42 the case of carbon nanotubes dispersed in epoxy resin. These authors determined a percolation threshold for their materials between 0.0225 and 0.04 wt.%.

In recent years, other polymer systems such as acrylonitrile-butadiene-styrene 86, poly(ethylene terephthalate) 87, polyvinylesters 88, polyester 89 and liquid crystal polymers90 have been explored to produce composites with carbon nanofibers.

2.4.2.2. Poly(methyl methacrylate)-carbon nanofiber composites

One of the first research works dealing with poly(methyl methacrylate)-carbon nanofiber (PMMA-CNF) composites was performed by Wu et al.91. They noted a 50 % reduction of the percolation threshold after addition of 1 wt.% of high density polyethylene (HDPE) to the PMMA-CNF composites. This phenomenon was interpreted as a result of the architecture of self-assembled conductive network constructed by selective adsorption of HDPE in VGCF/PMMA composites.

Zeng et al.92 fabricated PMMA-CNF composites by means of a counter rotating twin-screw extruder, pelletized, and fed into a spinning system in order to obtain composite fibers. It was observed that the tensile modulus of the composites did not improve significantly using high content of CNF (10 wt. %), but a 50 % increase was observed with half the carbon concentration. By applying a modified equation of the Cox model 75,76 for the composite tensile modulus, they determined that the fibers were randomly distributed in the matrix.

Jimenez and Jana were able to fabricate PMMA-CNF composites with a low percolation threshold by using a 2-D chaotic mixer.93,94 These authors used as-received

43 CNFs and mixed them with PMMA in a chaotic mixer and compared the data with same materials prepared in a commercial mixer. It was determined that the low shear mixing in the chaotic mixing device helped build up of carbon nanofiber networks from the dispersed fibers and the agglomerates. These helped to produce a percolation threshold of

2 wt. % CNF in a chaotic mixer compared to a percolation threshold of 6 wt.% for composites prepared in a commercial mixer. In the latter mixer, fibers experienced much damage. Schueler et al.95 stated that the presence of agglomerates was responsible for low percolation thresholds for carbon black in epoxy, which could not be explained in terms of a percolation theory but based on a theory for colloids.

Research regarding PMMA and CNT has been focused on the dispersion and alignment of the CNTs in the polymer matrix. Several approaches have been used in order to reach such level of dispersion and alignment; in-situ polymerization,96,97 solvent casting,98-100 melt mixing,101-104 and coagulation.105-107 Du et al.105-107 reported a new preparation route, called coagulation method, whereby CNTs and PMMA were combined in a suspension, thus preventing the nanotubes agglomeration. PMMA-SWNT composites prepared in such a way were found to be well dispersed and devoid of agglomerates. In this case, the value of tensile modulus increased from 3 GPa for PMMA to about 6 GPa with just the addition of 2 wt. % of SWNT. However, these authors observed an increase on the percolation threshold using aligned SWNT compared to composites made with unaligned SWNT. In this case, the electrical conductivity dropped from 10-4 to 10-10 S/cm for a PMMA composite with 2 wt. % loading of SWNT. They argued that the decrease in electrical conductivity was due to fewer contacts between the

44 nanotubes when aligned – thus aligned composites required more nanotubes to reach the percolation threshold.

2.4.2.3. Thermoplastic polyurethane-carbon nanofiber composites

Scientific reports on thermoplastic polyurethane-carbon composites can be found in conjunction with carbon fibers108-111, carbon black112,113, carbon nanofibers114, multi- walled nanotubes115-122, and single-walled nanotubes 123-125. The onset of thermal degradation of TPU has been reported to increase up to 17 °C for a composite having 30

108 wt. % of carbon fibers. The authors did not observe a major change in the Tg of the polyester-based TPU, even at high loadings of 10 μm-diameter carbon fibers. On the other hand, the tensile stress at break of TPU experienced a dramatic drop due to addition of carbon fibers, and only after addition of 30 wt. % of fibers that the stress at break of the composites reached that of the resin alone. In the case of carbon black (CB), Segal et al.112 prepared TPU-CB composites by melt mixing in a Brabender Plasticorder and observed that the Tg of TPU decreased with CB content. They argued that CB particles interacted more with the polar groups of the hard domains which caused an increased in the segmental mobility of the soft segments. In terms of electrical conductivity, Segal et al.112 found that these materials percolated at 3 wt. % and explained it based on the two- phase morphology present in TPU. These authors concluded that at low CB content, carbon particles interacted only with hard segment domains, thus building up conductive paths.

45 In the case of nanocarbon materials, most of the recent research has been focused on carbon nanotubes much more than carbon nanofibers. Heat-treated CNFs were solution mixed with TPU having crystallizable soft segments, by Koerner et al.114, showing a percolation threshold of about 0.010 wt. %. Carbon nanofibers used in that work had a post production process of heating to temperatures up to 3000 °C, creating highly electrically conductive carbon nanofibers.126 One important observation made by these authors is that the crystallinity of the soft segments exerted strong influence on the tensile properties of TPU-CNF composites. They observed an increase in the Young’s modulus, decrease of the stress at the apparent yield point, and an irregular behavior of the values for stress and elongation at rupture with the addition of CNF. From differential scanning calorimetry experiments, Koerner et al.114 observed enhanced initial soft- segment crystallinity with the addition of CNF up to 5 wt.% as determined by the first

DSC scan. On the other hand, a second DSC scan demonstrated no enhancement of the crystallinity of soft segments which suggested that the initial increase in crystalline fraction with CNF content was probably derived from the inhomogeneous strain distribution created during solvent removal. Finally, the authors concluded that parameters like nucleation, strain induced crystallization, polymer crystallite orientation, and fiber alignment all played roles in determining the reinforcement effect of CNF in

TPU.

On the other hand, Xia et al.118 fabricated and compared TPU composites of

SWNT and MWNT. These authors used ball milling and a dispersing agent to improve the interaction between the polyol and the nanotubes prior to synthesis of the polyurethanes. They observed that for both types of composites – i.e., TPU-SWNT and

46 TPU-MWNT – Tg decreased with carbon content due to a higher phase separation of soft- segment domains from mixtures with hard segments. They also observed that SWNT and

MWNT have different reinforcement effects on TPU, i.e., SWNT yielded better tensile strength and elongation at break, while MWNT improved the Young’s modulus of TPU.

It is noted from prior studies that polyurethanes degrade through three mechanisms: (i) dissociation of the urethane bond into its starting components, e.g., polyol and isocyanate groups, (ii) breaking of the urethane bond with the formation of primary amine, carbon dioxide and an olefin, and (iii) breaking of the urethane bond into secondary amine and carbon dioxide.127 It is also possible to distinguish two characteristic stages of thermal degradation of TPU in a thermogravimetrical experiment; an initial first stage related to the hard segments and a second one related to the soft segments in TPU.118,127 Xia et al.118 found out that these degradation reactions occurred respectively at about 332 °C and 393 °C. They noted that the incorporation of either

SWNT or MWNT only delayed the second degradation temperature by ~ 5 °C. Therefore, the authors concluded that carbon nanotubes may preferably interact with the soft segment in TPU.

2.4.2.4. Effect of fiber surface modification

Although it has been observed that the mechanical properties of polymer composites containing oxidized CNFs are improved,88,128-130 it is also true that the electrical conductivity of such composites is affected negatively.88 Polymer composites of carbon nanofibers having a more perfect graphitic character show higher conductivity

47 values than composites made with more defective as-grown CNFs.129 Similar findings have been found for epoxy systems containing single-walled nanotubes SWNT131 and multi-walled nanotubes MWNT.132 Tensile properties of epoxy having acid-treated

SWNTs were higher than corresponding composites fabricated with untreated SWNTs.131

Conversely, the electrical conductivity of epoxy/acid-treated MWNT composites decreased compared to those with untreated CNT. Decrease of the electrical conductivity was attributed to damage caused on the MWNT by the oxidation treatment.132

Cortes et al.133 treated carbon nanofibers with nitric acid, and copper, and mixed with polypropylene in a Banbury-type mixer. The authors achieved a percolation threshold at about 5 wt. % CNF, which was less than that previously obtained by the same research group (9-18 wt. %). Nitric acid oxidation and the addition of copper did not produce significant changes in the mechanical and rheological properties of the polymer, although a decrease in tensile strength was observed. Xu et al.88, on the other hand, observed a substantial increment on the electrical percolation composition after they treated VGCFs with boiling nitric acid.

Brandl et al.128,134 treated CNFs with plasma oxygen, and combined them with PP.

They observed an increment of the surface energy of the fibers after the plasma treatment.

Mechanical properties of the composites improved after plasma functionalization.

However, the electrical conductivity of these composites was not investigated.

Seo et al.135 combined epoxy and carbon nanofibers treated with atmospheric plasma, and concluded that the general increment in thermal stability of the polymer was due to the reduction of the mobility of the matrix chains by interactions with the fibers.

48 Also, they found an influence of this treatment on the Izod impact strengths of the plasma-treated carbon filler-reinforced polymer matrix composites.

There are few papers in which surface-treated MWNT were used with

PMMA.96,97,136-138 Among those papers, only a few dealt with oxidized MWNT.96,97,138 In all those cases, PMMA-MWNT nanocomposites were fabricated by an in-situ polymerization procedure. A significant improvement in the tensile strength was observed by Jia et al.96 when acid-treated MWNT were used in comparison with non- treated MWNT. On the other hand, Sung et al.97 observed a dramatic reduction of the electrical conductivity by about 6 orders of magnitude, after performing electrospinning of composites of PMMA with acid-treated MWNT. These authors attributed such reduction to the presence of porosity in the electrospun nanofibers and to the fact that

PMMA chains wrapped perfectly around MWNTs which lowered the electrical conductivity.

PMMA composites with oxidized CNF – labeled as CNFOX – were investigated recently by Jimenez and Jana.139 They observed improvement in the quality of dispersion, thermo-oxidative stability (TOS), and dynamic mechanical properties compared to

PMMA composites of untreated CNF. However, electrical conductivity of PMMA-

CNFOX dropped substantially compared to those of PMMA-CNF.

Concerning TPU composites with surface-treated CNF and CNT, again, most of the existing literature is focused on using MWNT 115-117,119,120,122, and SWNT.124 Kwon et al.116,117 used acid-treated MWNTs with 3.5 atomic % of oxygen on their surface. This mild oxidation of the surface allowed them to remove impurities of the nanotubes surface, and to prevent too much opening of the graphite double bond system. Composites of

49 treated CNT with waterborne TPU exhibited better mechanical and thermal properties compared to those of untreated MWNT. However, electrical conductivity was insensitive to nanofibers treatment. On the other hand, Mondal and Hu120 studied the thermal stability and oxidative behaviors of a TPU system containing functionalized MWNT.

They observed a reduction in thermal stability of the TPU polymer at lower MWNT contents (0.25 and 0.50 wt. %), while thermal and oxidative stabilities improved with higher MWNT contents. These authors contend that improved thermal conductivity at higher CNT loading retarded the degradation of TPU.

Recently, Xiong et al.121 performed air and acid oxidation on MWNTs in order to generate carboxylic groups on the CNT surfaces. Later on, they refluxed carbon nanotubes with SOCl2 before reacting with ethylendiamine to produce MWNT-amide derivative. They noticed using Atomic Force Microscope (AFM) that after such chemical treatments, the nanotubes were clean but were very short in length in comparison with the untreated MWNTs. In spite of the surface modification, the authors still could observe agglomerates of MWNT by means of SEM and TEM. They also observed improvements in thermal, and tensile properties in presence of MWNT.

2.4.2.5. Shape memory TPU nanocomposites

Research on shape memory effects in TPU-carbon composites has received recent attention because of the dual effect of these materials – carbon materials provide mechanical reinforcement, while increase electrical conductivity – making them potential candidates for resistivity heatup and actuators of shape memory effects. Therefore, the

50 recovery of the shape can be triggered by either thermal or electrical stimuli. Combining carbon materials with crystallizable soft segment-TPU for shape memory applications require two very important considerations: (1) to preserve or enhance the soft segment crystallinity and (2) to preserve stable hard-segment domains at temperatures above the melting point of the soft segment regions.67,140

Li et al.141 utilized carbon black as filler of a PCL7000/MDI/BDO-based TPU and found that the PCL crystallinity decreased up to 10 wt.% of CB, and then kept almost unchanged at higher CB contents. The authors stated that the final PCL crystallinity obtained in those composites was enough for shape memory applications. The existence of a plateau in the storage modulus after melting of the PCL crystallites indicated the presence of physical cross-links, i.e. hard domains. The length of such plateaus increases with increasing the hard-segment content in a specimen.140 Furthermore, Li et al.141 determined that the magnitude of E’ at 80 °C increased monotonically with CB content.

This behavior was attributed to carbon black serving as an effective filler to strengthen the TPU system. On the other hand, the authors observed that the presence of CB in TPU affected the shape memory properties of pure TPU. A decrease of the speed of shape recovery (recovery rate) occurred when adding CB to the polymer. One explanation given by the authors is that CB increased the bulk viscosity of the material after the melting of PCL crystals, thus increasing the relaxation time of the PCL chains in TPU.

They finally arrived at the conclusion that TPU-CB composites do not have good shape memory behavior in comparison with TPU alone.

Mondal and Hu142 prepared shape memory TPU based on MDI, BDO and a polyol mixture of poly(tetramethylene glycol), (PTMG) and poly(ethylene glycol), (PEG)

51 and reinforced it with aniline-functionalized MWNTs. Soft segments melting temperature for neat TPU was around 22.2 °C. A slight rise in the heat of fusion for a composite having 0.25 wt. % of the functionalized MWNT was observed. However, higher CNT contents produced lower enthalpies of fusion, even lower than that of TPU. It was said that MWNT have a dual effect on the soft segment crystallinity; as a nucleating agent and as filler. The nucleation effect was more noticeable at very low contents of MWNT. By adding higher amounts of carbon nanotubes, the motion of the soft segment chains were restricted, thereby suppressing crystal formation. Similar results were found by Wu and

Chen 143 in neat PCL reinforced with MWNT. These authors found that the activation energy for PCL crystal formation decreased when using 0.25 wt. % of MWNT, and then increased for higher filler contents.

As it was previously mentioned, TPU-CNF or TPU-CNT may have a dual approach to trigger the shape memory effects. Because of high electrical and thermal conductivities of these fillers, the shape memory behavior can be activated in these TPU composites either thermally or electrically. In order to obtain a total or partial recovery of the shape, these materials need to possess a certain electrical resistance low enough to allow the flow of electrical current, but high enough to generate Joule heating to melt the soft segment crystals. As the melting temperature of PCL crystals range from 40-60 °C, this material can be selected to form the soft segments in materials with shape memory effect induced by Joule heating. Cho et al.144 produced TPU-MWNT based on a

PCL3000/MDI/BDO system, with a molar ratio of 1/6/5 to generate a system with 40 wt.% hard segments. The carbon nanotubes used were acid-treated. An application of 60

V to a 5 wt.% composite generated a rise of temperature by 35 °C in eight seconds. The

52 authors calculated the energy-conversion efficiency to be around 10 % which is similar to the value obtained for shape memory alloys. Finally they recorded the electric-field- triggered shape recovery of the 5 wt.% composite, and determined that acid treated

MWNT yielded better performance than untreated carbon nanotubes. Yoo et al.145 worked on the same TPU-MWNT system, however they focused on the effect of the preparation procedure on the shape memory behavior. In a Method 1, the prepolymer was prepared from a reaction of MDI and PCL, followed by the chain-extension step. In

Method 2, MWNT was first dispersed in PCL, and then MDI, and the chain extender were added into this mixture. Finally, in Method 3 the readymade polyurethane and

MWNT were solution mixed in dimethylformamide. They fixed MWNT content at 3 wt.%, and observed that the best tensile, electrical and shape recovery properties were obtained by following Method 2. Thus, they obtained high electrical conductivity, 0.28

S/cm, sufficient to make the sample heated above the PCL crystal melting (36-42°C).

Instead, the electrical conductivity obtained with the other two methods (~10-2 S/cm) was not enough to heat the sample above the transition of TPU.

As it could be seen from this literature review, most of the research has been focused on polymer composites containing carbon nanotubes, either SWNT or MWNT.

Some research has been done on polymer composites with CNF, fewer on oxidized or chemical treated CNF, and none using chaotic mixing to investigate the effect of the CNF surface chemistry on dispersion and properties. No comprehensive research work has been done to date using chaotic mixing to fabricate PMMA-CNF and TPU-CNF composites. The present work provides a complete characterization of polymer-CNF composites prepared by chaotic mixing in terms of such aspects as morphological,

53 characterization and relationship between mechanical, thermal, chemical, and electrical properties with the degree of dispersion.

54 CHAPTER III

EXPERIMENTAL

3.1. Materials

The experiments were conducted using poly (methyl methacrylate), and two thermoplastic polyurethanes. Carbon nanofibers with different surface chemistry were mixed with those polymers to assess the degree of dispersion and its effects on the electrical conductivity.

3.1.1. Poly (methyl methacrylate)

Poly(methyl methacrylate) (PMMA) is commercially the most important member of a range of acrylic polymers.146 PMMA is an atactic, amorphous, and transparent polymer with a glass transition temperature of 110 °C. The chemical structure of the repeat unit of PMMA is shown in Figure 3-1. Acrylics have outstanding outdoor resistance.147

55 CH3 CH C 2 n C O O CH 3

Figure 3-1 Chemical structure of PMMA.

Cast PMMA sheet has excellent optical properties and is more resistant to impact than glass. Acrylics have low water absorption, good electrical resistivity, and fair tensile strength.

Acrylics are available as compounds for extrusion, injection molding, blow molding, and casting. Typical applications include: outdoor signs, glazing, aircraft canopies, washbasins, lighting applications, knobs, handles, safety shields, and machine covers.

PMMA Perspex CP-80 was obtained in the form of pellets kindly supplied by

Lucite International (Cordova, TN) with a melt flow index of 1.8 (ASTM D 1238) and density of 1.19 g/mL (ASTM D 792). PMMA CP-80 is ideal for extrusion applications where stiffer polymer melt is required for processes such as sheet extrusion, or for thicker section rods or profiles. PMMA pellets were grinded, screened into powder particles with sizes no larger than 2 mm, and dried at 90 °C in vacuum for 24 hours. Table 3-1 shows some physical and mechanical properties of this specific grade of PMMA.

56 3.1.2. Thermoplastic polyurethanes

Two kinds of thermoplastic polyurethanes (TPU) were synthesized in this work

(Table 3-2). These systems were coded according to the hard segment content; TPU23 with 23 wt.%, and TPU33 with 33 wt.% hard segments.

Table 3-1 Physical and mechanical properties of PMMA.148

Property PMMA (Perspex CP-80)

Melt flow index, g/10 min 2.2 (230°C/3.8 kg)

Processing temperature, °C 210-260

Density, g/cm3 1.19

Tensile strength at break, MPa 69

Elongation at break, % 4.6

Tensile modulus, GPa 3.4

Table 3-2 Chemical composition of TPU systems.

Hard segments TPU Components Molar Ratio (wt. %)

TPU23 PPG-MDI-BDO 1-2-1 23

TPU33 PCL diol-MDI-BDO 1-6-5 33

57 TPU23 was synthesized to contain soft segment from polyether-based polypropylene glycol (PPG) (Figure 3-2) with 2000 molecular weight (ARCOL® 2000) supplied by Bayer (Pittsburgh, PA), and a hard segment based on 4,4’-diphenylmethane diisocyanate, MDI (see Figure 2-10, p.26 ) (MONDUR M® flakes) also from Bayer, and

1,4-butanediol, BDO from Avocado (Ward Hill, MA) as a chain extender. Dibutyltin dilaurate DABCO 120® from Air Products (Allentown, PA) was used as the catalyst.

CH3

H O O O n O H O

CH3 CH CH3 3

Figure 3-2 Chemical structure of PPG.

The second grade of TPU, TPU33, contained soft segments from diol-terminated poly(ε-caprolactone) (PCL diol) from Solvay (Bruxelles, Belgium) trademark CAPA®

2403D with a melting temperature range of 55-60°C, 4000 molecular weight and with a chemical structure depicted in Figure 3-3.

O O

O *(CH2)6* O H O n O n H

Figure 3-3 Chemical structure of PCL diol.

58 3.1.3. Carbon nanofibers

Carbon nanofibers (CNF) and oxidized carbon nanofibers (CNFOX) grades PR-

24-PS and PR-24-PS-OX respectively were provided by Applied Sciences Inc.

(Cedarville, OH) both with a density of 1.95 g/mL, fiber diameters around 200 nm, and average length around 65 μm. Oxidized nanofibers were obtained by air oxidation of

CNF at 400-500 °C149, although a detailed account of the surface treatment was not disclosed by the supplier.150 Several properties of PR-24-PS can be found in Table 3-3.

A third type of carbon nanofiber (CNFOL) was produced in our laboratory starting from CNFOX. CNFOL stands for a polyol-grafted carbon nanofiber and was prepared based on the following procedure:53

Table 3-3 Physical and mechanical properties of CNF.74

Pyrograph III® PR-24-PS Property

Ultimate Strength, GPa 7.0

Tensile Modulus, GPa 600

Diameter, μm 0.060-0.2

Length, μm 100

Density, g/cm3 1.95

Electrical Resistivity, μΩ-cm 55

59 (i) Analysis of the carboxyl groups content in CNFOX:

Carboxyl group content was determined according to the Rivin method.151 An amount of 0.10g of CNFOX and 50.0 mL of sodium bicarbonate 0.10 mol/L (EMD

Chemicals, New York) were charged into a 100 mL Erlenmeyer flask. The mixture was stirred with a magnetic stirrer for 4h, and filtered. A volume of 25.0 mL of the filtrate and

25.0 mL of hydrochloric acid 0.10 mol/L (Alfa Aesar, Ward Hill, MA) were charged into a 300 mL Erlenmeyer flask, and then the mixture was boiled to remove dissolved CO2 for

20 min. After cooling, it was titrated with sodium hydroxide 0.05 mol/L (Merck,

Germany) using phenolphthalein as indicator. Carboxyl content was calculated from the amount of sodium bicarbonate consumed by CNFOX.

(ii) Grafting reaction of PPG onto CNFOX

Grafting of PPG onto CNFOX having carboxyl groups was performed by the direct condensation of carboxyl groups on VGCF with terminal hydroxyl groups in PPG in the presence of N,N’-dicyclohexylcarbodiimide, DCC (99 % purity, Alfa Aesar, Ward

Hill, MA). CNFOX (0.10g), PPG (2 equivalents to carboxyl groups), DCC (1 equivalent to carboxyl groups), and 20.0 mL of tetrahydrofuran, THF (EMD Chemicals, New York) were charged into a 100 mL flask. This mixture was stirred with a magnetic stirrer at

60°C for 48 h. The product obtained from the grafting reaction was dispersed in methanol

(Fisher Scientific, Hampton, NH) which is a good solvent for PPG. The dispersion was centrifuged at 2000 rpm for 30 minutes until CNFOL precipitated completely.

60 Precipitated CNFOL was again dispersed in methanol, and the dispersion centrifuged.

This was repeated one more time in order to remove free PPG.

(iii) Percentage of grafting

Percentage of polymer grafting is calculated by the following equation:

Grafting (%) = [Polymer grafted (g) / VGCF used (g)] X 100 (16)

The amount of polymer grafted onto the VGCF surface was determined by weight loss when the polymer-grafted VGCF was heated from room temperature to 500 °C under nitrogen at 10 °C/min by using a thermal gravimetric analyzer.

3.2. Composites preparation procedures

A list of composite systems that were prepared is shown in Table 3-4 according to their composition and processing technique. Labels will start with TPU system, type of carbon nanofiber and weight percentage of the carbon material, e.g. TPU23-CNF2 indicates a TPU23 system mixed with 2 wt. % of untreated carbon nanofibers.

61 Table 3-4 Nomenclature of the composite systems.

Composite Polymer Preparation Filler content Filler system matrix method (wt.%)

PMMA-CNF PMMA CNF Chaotic mixer 1 0,0.5, 1, 2, 3, 4, 6, 8, 10

PMMA-CNF PMMA CNF Plasticorder 0, 0.5, 1, 2, 3, 4, 6, 8, 10

PMMA-CNFOX PMMA CNFOX Chaotic mixer 1 0, 2, 4, 6, 10

TPU23-CNF TPU23 CNF Chaotic mixer 2 0, 0.5, 1, 3

TPU23-CNFOX TPU23 CNFOX Chaotic mixer 2 0, 0.5, 1, 3

TPU23-CNFOL TPU23 CNFOL Chaotic mixer 2 0, 0.5, 1, 3

TPU33-CNF TPU33 CNF Chaotic mixer 2 0, 1, 3, 5, 7

TPU33-CNF TPU33 CNF Plasticorder 0, 1, 3, 5, 7

TPU33-CNFOX TPU33 CNFOX Chaotic mixer 2 0, 1, 3, 5, 7

3.2.1. PMMA-carbon nanofiber composites

PMMA-CNF composites were prepared separately in a two-dimensional chaotic mixer (chaotic mixer 1 or CM1 hereafter), and a commercial internal mixer, Brabender

Plasticorder (C. W. Brabender, Germany). PMMA-CNFOX composites were prepared only in the CM1. Figure 3-4 shows images of the CM1 set-up and its mixing chamber, while Figure 3-5 shows a picture of the mixing head of the Brabender Plasticorder.

Sketches of both types of mixing chambers can be found in Figure 3-6.

62 (a)

(b)

50 cm

5 cm

Figure 3-4 (a) Set-up of chaotic mixer 1 set-up and (b) chaotic mixing chamber. Scale bar is given to provide an idea about the dimension of the equipment.

5 cm

Figure 3-5 Mixing head of a Brabender Plasticorder. Scale bar is given to provide an idea about the dimension of the equipment.

The chaotic mixer chamber has a volume of 33 cm3 with a uniform mixing gap

(H) of 12.5 mm and rotor radius of 12.5 mm (Figure 3-6(a)). The mixing chamber of the

Brabender Plasticorder (Figure 3-6(b)) has a minimum mixing gap of 2.16 mm and internal volume of 60 cm3.

63 (a)

H H H H = 12.7 mm R R R = 12.5 mm

h (b)

H H H h = 2.16 mm R R R = 10.8 mm H = 13 mm

Figure 3-6 Diagrams showing geometry and dimensions of two mixing heads: (a)Chaotic mixer, and (b)Brabender Plasticorder mixer.

Co-rotating motion of each rotor in CM1 was delivered through servo motors controlled by a computer. The speeds of co-rotating rotors (uA, uB) were varied time- periodically in sinusoidal fashion with a phase lag of 90°, uA=U (1+cos 2πt/T); uB=U(1- cos 2πt/T), where U is the maximum rotor speed, ~0.144 m/s, t is time, and T is time- period. This resulted in a peak shear rate of 8.3 s-1 and 4.1 s-1 at the rotors surface and mixing chamber walls respectively. Then, the time-averaged shear rate of the mixer was

6.2 s-1.25 The angular displacement per period of each rotor, θ was 1440° - i.e., each rotor describes four revolutions around its axis in one period. Rotors reached peak speeds of

64 100 rpm. Desired amounts of dried and grinded PMMA and dried CNF or CNFOX were pre-mixed in a plastic bag before pouring into the chaotic mixer, which was preheated at

170°C. PMMA and CNF or CNFOX were mixed at 230 °C for a total period of 4 minutes unless otherwise stated. The mixing chamber was separated from the mixer set up at the end of mixing and cooled in ice-water mixture for 30 minutes to freeze the morphology. The mixed material was retrieved in the form of an 8-shaped block. A detailed description of the flow patterns, the effects of θ on mixing effectiveness and morphology development, and operation of CM1 used in this study can be found elsewhere.23-25

PMMA-CNF mixtures were also produced in Brabender Plasticorder at 230°C for a total mixing time of 5 minutes. The rotors turned at 60 rpm in a counter-rotating fashion with a speed ratio of 1.1. This resulted in a shear rate of 97s-1 at the narrow mixing gap,

(h), although the volume average mean shear rate in the mixer was 4.6s-1. A portion of the mixed material from internal mixer was grinded into small pieces, poured into the chaotic mixer chamber, and molded into an 8-shaped block. These 8-shaped specimens were used in measurement of electrical conductivity in the same fashion as for materials produced in chaotic mixer.

Also, a mini-compounder with co-rotating screws (Haake mini-Lab) with a capacity of 5 g was used to fabricate some of the composites at 230 °C, 100 rpm, and with 5 minutes of mixing.

65 3.2.2. TPU-carbon nanofiber composites

Both systems, TPU23 and TPU33 were synthesized by a two-step polymerization procedure. Stoichiometric amounts of MDI and polyol already detailed previously in

Table 3-2, were reacted in a setup sketched in Figure 3-7, at 80°C for 5 ½ hours (TPU23) and 2 ½ hours (TPU33) to produce the respective prepolymer.

Thermometer Motor

Propeller Stirrer

N2 gas

Oil Bath Hot Plate

Figure 3-7 Sketch of the setup to prepare a TPU prepolymer.

In the case of TPU23 system, a typical experiment required 80 g of the prepolymer hand-mixed with certain amount of carbon nanofiber (CNF, CNFOX or

CNFOL), and then combined with 2.88 g of BDO containing catalyst (7.9x10-4 mol/L).

The mixture was poured into the chaotic mixer 2 (or CM2) (Figure 3-8) preheated at

80 °C, and mixed for 90 minutes.

CM2 has a 70 cc chamber of 15 cm length, 7.5 cm width, and 9 cm depth, and rotors with a radius of 12.5 mm, and a mixing gap of 7.62 mm. Rotors co-rotating in a sinusoidal fashion at peak speed of 65 rpm were put into motion by two independent 66 servo motors, and waveform generation was similar as that of CM1 previously discussed.

This generated a peak shear rate of 9.5 s-1 and 5.4 s-1 at the rotors surface and at the mixing chamber walls. Correspondingly, time-averaged shear rates were 4.8 s-1 and 2.7 s-

1 producing a mean shear rate of 3.8 s-1. Details about CM2 can also be found elsewhere.152,153 On the other hand, a typical preparation procedure for TPU33 used 80 g of molten prepolymer, hand-mixed with a proper amount of CNF or CNFOX, and then mixed with 6.5 g of BDO without catalyst (only blank TPU33 required catalyst).

5cm

Figure 3-8 Mixing chamber of chaotic mixer 2.

The mixture was also poured in CM2 preheated at 110 °C, and mixed for 10 minutes. In both situations, TPU composite materials were collected out of the chamber, and compression molded; TPU23 composites were compression molded at 165 °C and

TPU33 at 220 °C. Pure TPU and composite systems were then kept at room temperature for further analysis.

TPU33-CNF composites were also prepared in a Brabender Plasticorder by reacting the prepolymer with BDO and catalyst at 110 °C for 1 min at 100 rpm, and then adding CNF and mixing at 140 °C, and 100 rpm for 5 minutes.

67 3.3. Characterization techniques

Carbon nanofibers were analyzed in terms of their morphology, surface chemistry, and thermal properties. PMMA and TPU composites were analyzed based on fibers dispersion, thermal, static and dynamic mechanical, and electrical properties.

Additionally, status of the extent of hydrogen bonding in TPU was investigated.

3.3.1. X-ray photoelectron spectroscopy

Specimens of CNF, CNFOX and CNFOL were subjected to XPS, in order to determine the chemical composition of the nanofibers surface. XPS was performed under high vacuum conditions with a pressure of 1x10-8 Torr, an aluminum anode, and scans were taken with 1eV resolution.

3.3.2. Scanning electron microscopy

Scanning electron microscopy (SEM), was used to reveal the state of the fiber agglomerates, fiber-matrix interface in the planes of fractured specimens, and aspect ratio of carbon nanofibers in the final composites. Samples were coated with silver using a

K575x sputter coater from Emitech (Kent, England) under argon gas atmosphere. Surface morphology was observed using a SEM S-2150 from Hitachi (Ibaraki, Japan) at 20kv.

68 3.3.3. Transmission electron microscopy and ultramicrotoming

Thin sections of about 200 nm were produced by ultramicrotome (Microstar,

Huntsville, TX) with a glass knife. Slices of PMMA composites were sectioned at room temperature. Sections of TPU composites were cut normal to the flow direction, at -80 °C and collected on a copper grid. In addition, film casting of TPU solutions in dimethylformamide on a copper grid was carried out. TPU solutions were produced under quiescent conditions to prevent additional damage to the carbon nanofibers.

Transmission electron microscopy (TEM) images were taken using Tecnai-12 TEM, from FEI (Hillsboro, OR) with an operating voltage of 120 kV.

3.3.4. Optical microscopy and image analysis

Thin slices of ~ 40-70 μm thickness were cut out of samples of PMMA composites along the flow direction by using an Isomet low-speed saw by Buehler (Lake

Bluff, IL). Micrographs of the slices were observed through a Leitz Laborlux 12 Pol S

(Oberkochen, Germany) optical microscope in transmission mode, and images were acquired with a Diagnostic Instruments Inc. 11.2 Color Mosaic digital camera (Sterling

Heights, MI). Images were analyzed using Image Tool Version 3.00 software in order to assess the quality of the dispersion and fiber alignment.

69 3.3.5. Electrical conductivity

The 8-shaped PMMA composite blocks obtained from the chaotic mixer 1 (Figure

3-9(a)) were cut into halves (Figure 3-9(b)) along the flow direction and used for measurement of volume electrical conductivity along the flow direction. The volume electrical conductivity was measured after applying a layer of conductive silver paste to the edges of the specimen to ensure good contact with the electrodes. The electrodes were connected to a Keithley 487 picoammeter-voltage source (Cleveland, OH), from which electrical resistance, and consequently volume resistivity and conductivity were calculated according to ASTM D 257. A maximum standard deviation of 10% was incurred in these measurements. The value of transverse volume conductivity, i.e., across the thickness of the 8-shaped specimen, was also measured by sandwiching the specimen between two heavy stainless steel plates (Figure 3-9(c)). Measured resistance, R, was used combined with the specimen cross-sectional area, A, and length, L, to estimate volume resistivity, ρ,

RA ρ = (17) L

Volume conductivity, σ, was calculated as the inverse of resistivity,

1 σ = (18) ρ

70 In the case of TPU composites, circular compressed sheets with 8 cm diameter and 2 mm thickness were placed in a Keithley 8009 Resistivity Test Fixture connected to the Keithley 487 picoammeter-voltage source in order to measure the surface and volume resistivity.

The following protocol based on ASTM D 257 method was followed in the measurement of electrical resistivity: (i) for insulating materials a voltage of 500 V was applied, and (ii) for conductive materials a voltage of 1 V was used. In both cases, the reading was taken after 60 seconds of applying the electrical voltage.

A (b)

A (a)

Ammeter (c) A Voltage source

Figure 3-9 Composite specimen and electrode set up for measurement of volume electrical conductivity in PMMA composites (a) 8-shaped specimen, (b) one half of 8- shaped specimen for conductivity measurement along flow direction, and (c) conductivity along the thickness direction. 71 3.3.6. Differential scanning calorimetry

Thermal properties of TPU composites such as soft segment glass transition temperature, (Tg,ss), soft segment melting, (Tm,ss), soft segment crystallization (Tc,ss), hard segment melting point (Tm,HS), and their respective enthalpies were determined using a

TA Instruments 2920 Modulated DSC (New Castle, DE) at a heating rate of 10 °C per minute from -100 up to 230 °C under nitrogen gas atmosphere. Double heating scans were also performed on specimens weighing 3.8 ± 0.6 mg from -100 to 100 °C at

10 °C/min each scan in order to determine TPU crystallinity. Percent of crystallinity was determined by comparing the enthalpy of fusion of PCL in TPU33 to 136 J/g which is the enthalpy of fusion for 100 % crystalline PCL.141 Cooling scans were also performed from

100 °C up to -30 °C at 10 °C/min.

3.3.7. Thermogravimetrical analysis

Thermal and thermo-oxidative stability of TPU and PMMA composites were investigated in a TGA 2050 device from TA Instruments (New Castle, DE). Specimens weighing ca. 6 mg were subjected to thermal scan from room temperature up to 1000 °C, at a scan rate of 20 °C/min. The temperature at the onset of thermal degradation or 5 % mass loss, T1, was estimated from a mass versus temperature plot. The temperature at the maximum mass loss rate, T2, was obtained from the plot of the first derivative of mass loss versus temperature. Finally, T3 was used to characterize the end of a thermal degradation.

72 3.3.8. Tensile test

Tensile properties of PMMA and TPU composites were analyzed by means of

Instron 5567 machine (Norwood, MA) with a crosshead speed of 1 and 100 mm/min respectively, and an effective length of 20 mm. Specimens were in a rectangular form.

On the other hand, tensile properties at 60 °C of TPU33 composites were analyzed in an

Instron 4204 tensile tester with heating capability, with a crosshead speed of 20 mm/min, and 20 mm gauge length.

3.3.9. Dynamic mechanical analysis

Dynamic mechanical analysis (DMA) was performed with a Pyris Diamond DMA, from Perkin Elmer-Seiko Instruments (Boston, MA). Analysis was performed in tensile mode at 4 °C/min under nitrogen atmosphere, and 1 Hz frequency. PMMA composites were heated from room temperature up to 140 °C, while TPU composites were scanned from -100 to 120 °C.

3.3.10. Fourier-transform infrared spectroscopy

Infrared spectroscopy was used to characterize the most important functional groups and to determine the extent of hydrogen bonding in TPU composites. The amount of carbonyl groups participating in hydrogen bonding can be described by the carbonyl hydrogen bonding index, R, given by,154

73 AHCO R A −1 A −1 (19) = = 1700cm 1730cm ACO

where A −1 and A −1 are the absorbances of deconvoluted peaks 1700cm 1730cm

-1 corresponding to hydrogen-bonded, AHCO appearing at about 1700 cm and free carbonyl

-1 group, ACO at about 1730 cm . The extent of hard segments linking with hard segments

(degree of phase separation, DPS) and the degree of hard segments linking with soft segments or carbon nanofibers (degree of phase mixing, DPM) can be determined from equations (20) and (21) respectively.154

R DPS = (20) R +1

DPM = 1− DPS (21)

Extent of hydrogen bonding in the hard domains in TPU33 was also estimated by tracking the ratio of absorbances of the hydrogen-bonded amine group, ANH appearing at

-1 -1 about 3330 cm to that of the aliphatic group, ACH at 2860 and 2960 cm .

Films of TPU were cast on KBr disc from a dimethylformamide solution, and FT-

IR spectra were recorded from 4000 to 800 cm-1 in a Perkin Elmer 16 PC FT-IR

(Wellesley, MA) in transmission mode, and 4 cm-1 resolution. Thermal-FTIR was also recorded by placing a KBr disc inside of a heating stage. Infrared spectra were taken at

60, 120, and 180 °C for TPU33 composites.

74 3.3.11. Gel permeation chromatography

The molecular weight, MW and molecular weight distribution, MWD of PMMA and TPU samples were determined in a Waters Model 515 GPC device equipped with a

2414 refractometer. Tetrahydrofuran (UV-HPLC quality) was used as the mobile phase at a rate of 1.0 ml/min through two Plegel 3 μm MIXED-E columns from Polymer

Laboratories. The column temperature was maintained at 35oC. Molecular weight identification was possible after calibration of the system with polystyrene standards. All samples were prepared at 0.1% (w/v) in THF.

3.3.12. Shape memory properties

In order to investigate the shape memory behavior of TPU33 and its composites with carbon nanofibers, rectangular strips of 50x5x0.5mm3 were cut out of the compressed TPU33 materials. Strips with 20 mm gauge length (L0) were clamped on an

Instron 4204 Tensile Tester adapted with a heating chamber, and the temperature was slowly raised to 60°C. When the temperature reached this value, the upper grip started to move at 20 mm/min to reach a predetermined length (L1) and then stopped. The specimen was cooled down with an electric fan and reached a temperature of 40 °C, which is below the melting point of PCL crystals, in approximately 30 seconds. Additional 9 minutes of cooling were needed to reach the room temperature (25 °C), after which the specimen was taken out of the tensile grips. The retained or fixed length, L2, was measured after

75 removal from the grips. Percentage of shape retention or fixity was determined by the following equation:

L − L Shape retention or fixity (%) = 2 0 x100 (22) L1 − L0

A set of strained specimens were placed in a Pyris Diamond DMA apparatus and the recovery force was scanned from room temperature to 100 °C at 4 °C/min, under constant length condition (L mode). Another set of specimens were placed in an oven at

80 °C for 10 minutes in order to induce shape recovery. The recovery length, L3, was measured and the shape recovery percentage calculated according to the following equation:

L − L Shape recovery (%) = 1 3 x100 (23) L1 − L0

Shape memory actions triggered by Joule heating were investigated for TPU33-

CNF composites. For this purpose, the composites were made electrically conductive. In typical shape memory demonstrations, rectangular films of composites were fixed between two alligator clips. While voltage was passed through the specimen, the surface temperature was recorded with an accuracy of ± 5 °C by means of an IR sensor gun. The accuracy of the IR sensor gun was previously checked against a standard thermocouple by simultaneously measuring the temperature of the same composite specimen.

Composites which generate enough heating on the surface by the passage of electricity were chosen for the shape memory demonstration. The initial deformed state 76 was produced as follows: the rectangular specimens were placed in an oven at 80 °C for 5 minutes, and deformed by bending at an angle inside the oven followed by immediate freezing in ice-water mixture to keep the deformed shape. The deformed specimen was then clamped between the alligator clips, and a desired voltage was applied. The change in the deformed shape was recorded by means of a video camera.

77 CHAPTER IV

PREPARATION AND CHARACTERIZATION OF CARBON NANOFIBERS

Two types of carbon nanofibers, CNF and CNFOX were used as-received, while

CNFOL was prepared in our laboratory starting with CNFOX. Morphology of the nanofibers and composites was analyzed by SEM, and the surface chemistry of the nanofibers was inspected by XPS.

4.1. Morphology of CNF and CNFOX

Figures 4-1(a) and 4-1(b) show scanning electron micrographs of as-received

CNF and CNFOX respectively. The as-received fibers were highly entangled – some fibers were available as bundles and knotted together by layers of amorphous carbon shown as the encircled white spots in Figure 4-1. These knots potentially acted as deterrent to dispersion during polymer-CNF mixing.80,83 These fibers also differed in the extent of their hydrophilicity. For example, when equal amounts of nanofibers were added to the same volume of deionized water and mechanically mixed, a stable dispersion was obtained in case of the oxidized carbon nanofibers (Figure 4-2).

78 (a) (b)

Figure 4-1 SEM images of (a) CNF and (b) CNFOX.

This indicates that CNFOX possesses a higher degree of hydrophilicity than CNF due to the presence of oxygen-containing functional groups such as alcohol, carbonyl, and carboxyl on the surface. This is discussed in detail in the section below.

Figure 4-2 Dispersion in water of CNF and CNFOX.

79 4.2. Surface chemistry of CNF and CNFOX

Figure 4-3 shows the XPS spectra of CNF and CNFOX. In CNF, only C(1s) produced a signal with significant intensity, while in the case of CNFOX, C(1s) and

O(1s) generated significant peaks. A very small peak due to O(1s) at about 530 eV still can be observed for CNF.

C(1s)

O(1s)

CNFOX Intensity (a.u.) Intensity

CNF

0 1000 800 600 400 200 0 Binding Energy (eV)

Figure 4-3 XPS spectra of both types of carbon nanofibers.

By determining the area under the peaks assigned to C(1s) and O(1s), it is possible to establish the oxygen/carbon ratio on the surface of the carbon nanofibers (see

Table 4-1). It is observed that CNFOX shows the presence of a substantial amount of oxygen compared to CNF. This fact accounts for better dispersability of CNFOX in water as seen in Figure 4-2.

80 Figure 4-4 shows the deconvolutions of the C(1s) peaks from high-resolution XPS of CNF (Fig. 4-4(a)) and CNFOX (Fig. 4-4(b)). The increase of oxygen-containing functional groups – mostly C=O groups – on the CNFOX surface is clear from the change in asymmetry of the C(1s) peak toward higher values of binding energies corresponding to non-graphitic C(1s).

Table 4-1 Oxygen/carbon ratio present on the carbon nanofibers.

Ratio of area Carbon Nanofiber under peaks O(1s)/C(1s)

CNF 0.015

CNFOX 0.13

25000 25000 (a) 20000 (b) 20000 CNF Graphitic CNFOX carbon Graphitic 15000 carbon 15000 Counts 10000 10000 Counts COOH, C=O COOH, COOR COOR 5000 5000

0 0 300.0 295.0 290.0 285.0 280.0 300.0 295.0 290.0 285.0 280.0 BE / eV BE / eV

Figure 4-4 Narrow XPS spectra of the C(1s) region in (a) CNF, and (b) CNFOX.

81 4.3. Preparation and characterization of CNFOL

A previous knowledge of the concentration of carboxylic groups on CNFOX was required before reacting this fiber with PPG 2000 to produce CNFOL. By using the Rivin method151, the amount of COOH groups in CNFOX was estimated in 0.4 ± 0.1 mmol per gram of carbon nanofiber. After the reaction and washing procedures were completed, the material was evaluated by TGA in order to evaluate the success of the reaction and to determine the grafting percentage. Figure 4-5 provides the weight loss values of CNF,

CNFOX and CNFOL at 500 °C. The difference in weight loss between CNFOL and

CNFOX was 3.3% which can be attributed to the amount of PPG 2000 attached to

CNFOX, i.e., the grafting percentage.

100 CNF 90

CNFOX 80 Weight (%) Weight CNFOL 70

60 50 200 350 500 650 800 Temperature ( °C)

Figure 4-5 Weight loss for CNF, CNFOX and CNFOL.

In view of the grafting percentage, mass of CNFOL measured in the TGA, and the amount of carboxylic groups, the amount of carboxylic groups used for grafting was 82 calculated as 4 %. This is a reasonable value considering that Wei et al.53 observed 1 % of the total COOH groups grafted per gram of CNF. These results indicate that the

“grafting to” procedure in fact produces very low grafting yields at low reaction rates due to slow diffusion of the polymer chains to the surface of the carbon nanofibers in order to react with the COOH groups.

25000 Graphitic carbon 20000 C 1s 15000 C-OR O 1s

10000 Counts COOH, C=O COOR 5000

0 300.0 295.0 290.0 285.0 280.0 BE / eV Counts (Arb. Units)

Oxygen/carbon ratio = 0.18

1200 1000 800 600 400 200 0 Binding Energy (eV)

Figure 4-6 Full and narrow (inset) XPS spectra of CNFOL.

In the inset of Figure 4-6, a high resolution XPS of the C(1s) region in CNFOL is shown. A comparison of this plot with that of Figure 4-4 reveals the appearance of a new peak corresponding to the C-O-R bond which comes from the incorporation of PPG 2000

83 to the carbon nanofiber surface. The occurrence of such peak together with their weight loss analysis indicates that the grafting reaction of PPG with CNFOX was carried out successfully.

As it was discussed in Chapter II, section 2.2.5, normally chemical treatments of the carbon nanofibers are expected to produce some damage or shortening of the fibers.

In order to have an insight of the resultant morphology of the CNFOL fibers, SEM and

TEM images were taken (Figure 4-7). The TEM image in Figure 4-7(b) was obtained from a film that was solution cast in order to avoid chopping of the fibers during microtoming.

(a) (b)

Figure 4-7 Morphology of CNFOL as seen by (a) SEM, and (b) TEM.

As it can be seen from Figure 4-7(a), CNFOL fibers also show amorphous carbon

(inside of white circles in the Figure) that was not removed after the chemical treatment.

Although fibers seem to be less entangled than that of CNF and CNFOX (Figure 4-1), they still keep reasonable lengths as observed in the SEM image in Figure 4-7(a). In addition, no appreciable breakage of the nanofibers could be observed from the TEM

84 image in Figure 4-7(b). From this Figure, an aspect ratio of about 45 can be estimated in the case of long fibers.

4.4. Summary

From the results presented in this Chapter, it can be inferred that the present carbon nanofiber materials are highly entangled, and show significant amounts of amorphous carbon regions. Such morphology indicates that attempts to disperse these fibers in a viscous material such as polymer will be difficult. One alternative is then to utilize carbon nanofibers with polar functional groups that can improve the affinity towards polar polymers such as PMMA, polyurethanes or polyamides. Oxidation of the carbon nanofiber surface is known to increase the amount of atomic oxygen in the graphitic structure in the form of organic functional groups. The presence of such functional groups as –COOH allowed us to treat nanofibers with polyol in an effort to obtain CNFOL. This also showed that by controlling the surface chemistry of these carbon materials one can insert a desired functionality, but at the same time face the risk of affecting the properties such as mechanical or electrical due to extensive treatments.

85 CHAPTER V

PMMA-CARBON NANOFIBER COMPOSITES

This section discusses the preparation and properties of PMMA-carbon nanofiber composites and the effect of composition, processing technique, mixing time and surface chemistry of CNF on the electrical, thermal and mechanical properties.

5.1. CNF composition and processing technique

The values of volume electrical conductivity of the composites were measured both along the flow direction, and across the thickness direction. The objective of such measurement was to identify any differences in electrical conductivity due to anisotropic distribution of carbon nanofibers along and across the flow direction. As can be observed in Figure 5-1, composites prepared in the chaotic mixer and Brabender Plasticorder showed percolation compositions at about 1.5 and 5 wt. % CNF, respectively. These values are lower than that reported by Wu et al.91 who observed conductivity at 7 wt. % of CNF in PMMA (please refer to Chapter II, section 2.4.2.2).

86 1.E+00100 ■ Chaotic mixer

1.E-0210-2 ▲ Plasti-corder (a) Along flow direction Thru thickness (S/cm) 10-4

σ 1.E-04 (b) 1.E-0610-6 (d)

1.E-0810-8 (c)

-10 Volume conductivity, 1.E-1010

1.E-1210-12 012345678910 Carbon Nanofiber content (wt %)

Figure 5-1 Volume electrical conductivity of PMMA-CNF composites prepared in chaotic mixer and Brabender Plasticorder.

Figure 5-1 also shows the values of electrical conductivity measured through the thickness of a specimen. The values of through the thickness conductivity (Curve d,

Figure 5-1) of materials prepared in Brabender Plasticorder were practically the same as those of longitudinal conductivity (Curve c, Figure 5-1). This result is reasonable due to isotropic nature of mixing in Brabender Plasticorder mixer. However, a significant difference is observed between the conductivities along the flow direction (Curve a,

Figure 5-1) and through the thickness in materials prepared by chaotic mixing (Curve b,

Figure 5-1). Similar observations were found in PMMA-CNT films that showed higher electrical conductivity along the flow direction than perpendicular to it.102 In the present case, the anisotropy originated from the two-dimensional nature of the mixing flow 25, which expedited the mixing in the flow direction, but not through the thickness.

87 In order to elucidate the possible reasons for the observed trends in electrical conductivity, a multi-scale morphological analysis using optical, scanning, and transmission electron microscopy techniques was performed. Figure 5-2 shows optical micrographs of thin films of composites with carbon nanofiber contents less than or around percolation threshold in each case. Fiber agglomerates in the form of “striations” parallel to the flow direction, shown by arrows, can be seen in composites mixed in the chaotic mixer (Figures 5-2(a) and 5-2(b)). Fiber agglomerates are also seen in Brabender- mixed composites, but there was no preferred orientation is observed (Figures 5-2(c) and

5-2(d)). In addition, smaller particle agglomerates were found compared to chaotically mixed materials as revealed by the histogram of particle size distribution shown in Figure

5-3. Particle sizes ranged between 1 and 5 μm for a composite with 4 wt. % CNF prepared in Brabender Plasticorder compared to 10 to 20 μm fiber agglomerates in chaotically mixed materials. In view of this, it can be inferred that chaotic mixing produced much less dispersive mixing and more anisotropic distribution of fibers in

PMMA matrix. This type of anisotropic morphology favored the formation of conductive networks by using small quantities of fibers, thus lowering the value of percolation threshold. Similar results have been observed in chaotically mixed polystyrene-carbon black composites17,21, and in epoxy-CNT composites prepared by intense stirring of the epoxy mixture.85

Figure 5-4 indicates that composites prepared in chaotic mixer did not undergo much breakage and contained much longer fibers than those prepared in Brabender

Plasticorder mixer. Fiber length distributions were generated by computing length of individual fibers from several SEM images of the same sample specimen. 88 100 μm 100 μm

(a) Chaotic 1 wt. % CNF (b) Chaotic 2 wt. % CNF

100 μm 50 μm

(c) Brabender Plasticorder 4 wt. % CNF (d) Brabender Plasticorder 6 wt. % CNF

Figure 5-2 Optical micrographs of PMMA-CNF composites in the lower and upper limit of their electrical percolation threshold according the processing technique. Arrow indicates flow direction.

The presence of large size fiber agglomerates observed in chaotically mixed materials can be attributed to mild shear conditions prevalent in the chaotic mixer. Such mild shear was the result of a wide gap size between the rotors, and periodically varying speed of the rotors.

89 0.9 1 wt. % chaotic 0.8 2 wt. % chaotic 0.7 4 wt. % Brabender

0.6

0.5

0.4 Frequency 0.3

0.2

0.1

0 1-5 10-15 20-25 30-35 40-45 50-55 60-65 70-75 80-85 90-95 100- 200- 150 500 Size range (μm)

Figure 5-3 Particle size distribution for PMMA-CNF composites prepared by two processing techniques.

70

Chaotic Plasti-corder 60

50

40

30

20

10

0 1357910-50 Fiber length (μm)

Figure 5-4 Fiber length distribution of fibers extracted from a PMMA composite with 4 wt. % of CNF prepared in chaotic and Brabender Plasti-corder.

90 Note, however, that the mean value of shear rate in Brabender Plasticorder was maintained at 4.6 s-1, which is in the same neighborhood of the mean shear rate in the chaotic mixer.

Chaotic mixing also provided alignment of fibers as evident from optical microscope image in Figure 5-5. It is seen that long carbon nanofibers were pulled out of agglomerates, and oriented along the flow direction. These unwound fibers provided electrically conductive paths between neighboring agglomerates, which additionally explains why chaotically mixed materials showed much lower percolation threshold in

Figure 5-1. Cooper et al.101 saw similar alignment of particles in PMMA-CNF and

PMMA-CNT composites prepared by a multi-step method of power-mixing, sonication, batch mixing, and finally single-screw extrusion.

Flow direction

Figure 5-5 Optical micrograph showing conductive networks in 4 wt. % CNF composite prepared in the chaotic mixer. Arrow indicates flow direction.

Thermo-oxidative stability, TOS, of the composites was studied using TGA under static air environment. The values of derivative of weight loss as function of temperature are plotted in Figure 5-6 for representative composite materials. It is clear that the

91 presence of CNF increased the thermal stability of the composites. The temperature at peak shifted to higher values and the degradation peak became wider with the increase of fiber content. This indicates that more PMMA chains were affected by the presence of carbon nanofibers. The values of onset and end temperatures of thermal degradation, T1,

T3 respectively, and the temperature at the maximum mass loss rate T2 are shown in Table

5-1

3.0

2.5 PMMA 1 wt. %

) 2.0

6 wt. % 1.5

10 wt. % Deriv. Weight (%/°C 1.0 3 wt. %

0.5

0.0 250 275 300 325 350 375 400 425 Temperature (°C)

Figure 5-6 Thermo-oxidative stability of PMMA-CNF composites prepared by chaotic mixing.

It is apparent that the value of T1 did not change much with CNF content; similar observations were made in the case of PMMA-SWNT105 and PMMA-CNF composites.92

An increase in the values of T2 and T3 indicates delay of thermo-oxidative decomposition of the polymer. The acquired thermal stability may be due to the interaction of the CNF

92 with the free macroradicals generated during thermal degradation of PMMA, as it has

155 been observed in the presence of fullerene C60.

Table 5-1 Thermo-oxidative degradation of PMMA-CNF composites prepared in the chaotic mixer.

T (°C) PMMA 1 wt. % 3 wt. % 6 wt. % 10 wt. %

T1 306 305 304 309 304

T2 324 320 330 344 348

T3 355 359 374 385 387

The dynamic mechanical properties of composites are presented in Figures 5-7 and 5-8. Values of the storage modulus of composites prepared in chaotic mixer shows dramatic increase over PMMA at low CNF content, <1wt. % (Figure 5-7(a)). However, a slight decrease is observed for CNF greater than 2 wt. %. For composites prepared in

Brabender Plasticorder, (Figure 5-7(b)), the increase of E’ over PMMA was modest and almost independent of the loading of CNF. From tan δ curves in Figure 5-8, it is seen that the glass transition temperature decreased when CNF was added, using either chaotic

(Figure 5-8(a)) or Brabender Plasticorder mixer (Figure 5-8(b)). This indicates the lack of a good interaction between the fibers and the polymer matrix. The maximum values of tan δ, found at glass transition temperature, were much smaller in case of composites prepared in chaotic mixer (Figure 5-9). Finegan 130 reported that higher values of E’ and smaller values of the loss tangent originate from higher aspect ratio of fibers.

93 3.0 PMMA 2.5 0.5 wt. % 1 wt. % 2 wt. % 2.0 3 wt. % 4 wt. % 1.5 E'(GPa) 1.0

0.5 (a)

0.0 20 40 60 80 100 120 140 Temperature (°C)

3.0 PMMA 2.5 0.5 wt. % 1 wt. % 2 wt. % 2.0 3 wt. % 4 wt. % 1.5 E'(GPa) 1.0

0.5 (b) 0.0 20 40 60 80 100 120 140 Temperature (°C)

Figure 5-7 Storage modulus of PMMA-CNF composites prepared in (a) chaotic mixer, and (b) Brabender Plasticorder.

This along with the fiber length distribution in Figure 5-4 and the optical image in

Figure 5-5 may explain the higher values of high E’ observed in composites with 0.5 and

1 wt. % CNF prepared in the chaotic mixer. The values of tensile properties showed similar trends. Table 5-2 shows that the values of elongation at break are lower than that of PMMA, due to the presence of carbon nanofibers. The values of stress at break (Figure

5-10(a)) decreases, while the Young’s modulus slightly decreases for materials fabricated in the Brabender Plasticorder and remains constant for those mixed in the chaotic device

(Figure 5-10(b)).

94 1.6 PMMA (a) 0.5 wt. % 1.2 1 wt. % 2 wt. % 3 wt. % δ 0.8 4 wt. % tan

0.4

0.0 100 110 120 130 140 Temperature (°C) 1.6 PMMA (b) 0.5 wt. % 1.2 1 wt. % 2 wt. %

δ 3 wt. % 0.8 4 wt. % tan

0.4

0.0 100 110 120 130 140 Temperature (°C)

Figure 5-8 Variation of tan δ with temperature for PMMA-CNF composites prepared in (a) chaotic mixer, and (b) Brabender Plasticorder. 1.8

1.6 Brabender Plasti-corder δ 1.4 tan Chaotic 1.2

1 01234 CNF wt. %

Figure 5-9 Maximum values of loss tangent (at glass transition temperature) of PMMA- CNF composites prepared in chaotic mixer, and Brabender Plasticorder.

95 Table 5-2 Elongation at break (%) of PMMA-CNF composites.

Composite (wt. %) 0 0.5 1 2 3 Chaotic mixer 5.5 ± 1.3 3.1 ± 1.6 4.0 ± 0.6 3.1 ± 1.4 2.9 ± 0.3 Plasticorder 5.2 ± 1.5 2.6 ± 1.5 3.7 ± 0.9 3.2 ± 0.3 4.0 ± 1.1

50 (a)

40

Brabender Plasticorder

Strees (MPa) atStrees break 30

Chaotic mixer

20 0123 CNF content (wt. %) 2.0 (b)

1.5 Chaotic mixer

1.0 Brabender Plasticorder Young's Modulus (GPa) Modulus Young's

0.5 0123 CNF content (wt. %)

Figure 5-10 Tensile properties of PMMA-CNF composites (a) Stress at break, and (b) Young’s modulus.

96 By using a modified Cox’s model 75,76, the Young’s modulus of a PMMA-CNF composite can be estimated by a combination with the rule of mixtures according to Eq.

(23).

⎛ tanh β ⎞ Ec = ()1−V Em + q⎜1− ⎟VE f (23) ⎝ β ⎠

where Ec is the composite modulus, V is the volume fraction of the fibers in the composite, q is the orientation factor, Ef is the fiber axial modulus and β is a factor defined by Eq. (24).

l E β = m (24) d ()1+ν E f ln()π 4V

where l and d are the length and diameter of the fibers respectively, Em and ν is the Young’s modulus and Poisson’s ratio of the polymer matrix. Using Em as 1.31 GPa,

92 Ef as 100 GPa , l as 65 μm, d as 150 nm, and the respective densities as presented in

Tables 3-1 and 3-3 in Chapter III. Thus, this model predicts an increase of 43% in tensile modulus (~1.87 GPa) for 1 wt. % (V = 0.6 vol %) of fully dispersed and aligned CNF.

This theoretical value contrasts with the present experimental results, and indicates the high level of carbon nanofiber agglomeration in the composites produced either in the chaotic or conventional mixer. Another factor that influenced the mechanical properties of the composites was the presence of bubbles which were very difficult to remove.

97 5.2. Effect of chaotic mixing time

It is apparent from Figures 5-1 and 5-2 that electrical conductivity depended greatly on the state of dispersion of CNF. In view of this, we also evaluated the effects of mixing time on evolution of CNF structures, especially with compositions around percolation threshold and well above the percolation threshold. The results are compared with a case when no mixing was provided, e.g., premixed PMMA and CNF mixtures were allowed to melt in the chaotic chamber in the absence of mixing and then solidified.

This represents a case with no fiber alignment, and the state of the initial condition of the fibers agglomerates in the polymer before dispersion was carried out in the chaotic mixer.

Figure 5-11 presents how volume electrical conductivity changed with mixing time in the chaotic device. Composites with no mixing showed the highest conductivity, while the conductivity value decreased with the mixing time. This effect is more dramatic for composites with CNF content around the percolation threshold, e.g., 2 wt. % (see Figure

5-1). It can be seen that a composite with 2 wt. % CNF is electrically conductive (with σ

~ 10-5 S/cm) up to 4 minutes of mixing, but it turned into an insulator after 10 minutes of mixing. One can argue that nanofiber structures underwent breakage with continued mixing, which resulted in reduced conductivity. However, composites of 6 and 10 wt. %

CNF showed slight or no reduction in the electrical conductivity, respectively, with longer mixing times. This was due to the fact that those compositions are far above the percolation, and this indicates that the presence of agglomerates is mainly responsible for the electrical conductivity in the composites. In view of this, optical micrographs of 2 wt. % CNF composites at different mixing times were compared as seen in Figure 5-12.

98 The initial mixture (0 min) shows large fiber clumps with sizes as large as 500 μm.

A more uniform distribution of fibers is seen after 10 minutes of mixing, which resembles the one obtained with Brabender-Plasticorder (see Figure 5-2(d)). Dynamic mechanical properties of these materials were also analyzed, and the values of storage modulus as function of mixing time is presented in Figure 5-13. It is evident that some mixing of CNF in PMMA was necessary to obtain good mechanical properties.

1.E+00100 ( 1.E-01-1 m) 10 10 wt. % 1.E-0210-2 6 wt. % 1.E-0310-3 1.E-0410-4 1.E-0510-5 1.E-0610-6 2 wt. % -7 Volume conductivity (S/c conductivity Volume 1.E-0710 -8 1.E-0810 0246810 Mixing time, (min)

Figure 5-11 Effect of mixing time on the electrical volume conductivity along flow direction for composites prepared in the chaotic mixer.

The storage modulus at room temperature increased over that of the initial mixture after both 4 and 10 minutes mixing. In light of Figures 5-11, 5-12, and 5-13, one can infer the following: (a) at or near percolation threshold, both electrical conductivity and mechanical properties change with CNF dispersion, but in opposite manner; (b) at

CNF content well above percolation threshold, electrical conductivity is not very sensitive to the size of the fiber agglomerates, as enough “conductive paths” are created 99 due to large volume fraction of fibers, and (c) the chaotic flow exerted alignment and stretching of the fiber agglomerates in the flow direction.

0 min 4 min

500 μm 100 μm

10 min

100 μm

Figure 5-12 Effect of mixing time on morphological characteristics of PMMA-CNF 2 wt. % prepared in the chaotic mixer.

2.0 PMMA 4 min 1.6

10 min 1.2

0 min E'(GPa) 0.8

0.4

0.0 20 40 60 80 100 120 140 Temperature (°C)

Figure 5-13 Effect of mixing time on the storage modulus of PMMA-CNF 2 wt. % prepared in the chaotic mixer.

100 5.3. Effect of surface chemistry in CNF

Optical microscopy revealed that nanofiber composites of CNFOX showed much better dispersion as presented in Figure 5-14. It is worth noting that mixing conditions were identical, and equal amount of carbon nanofibers to those used in the case of CNF were used. Carbon agglomerates were of much reduced size for PMMA-CNFOX (Figure

5-15). Close inspection of Figure 5-14 shows that a large fraction of agglomerates in the case of CNFOX were smaller than 15 μm. In addition, some agglomerates with size larger than 100 μm could be seen in the case of CNF, while they were absent in the case of CNFOX. This is a ramification of the improvement in better dispersion, owing to better fiber – polymer interactions.

The evidence of better polymer-fiber interface is also manifested with good wetting between the fibers and the matrix. It is seen in Figure 5-16 (circled areas) that carbon nanofibers were more wetted by the polymer in PMMA-CNFOX composites, which in turn produced better fiber dispersion in the matrix.

(a) (b)

100 μm 100 μm

Figure 5-14 Optical micrographs of PMMA 2 wt. % composites with (a) CNF and (b) CNFOX. Arrows indicate flow direction.

101 0.4 CNF CNFOx 0.3

0.2 Frequency

0.1

0 5-10 15-20 25-30 35-40 45-50 55-60 65-70 75-80 85-90 95-100 150- 200 Size range (μm)

Figure 5-15 Particle size distribution histogram for PMMA 2 wt.% composites with CNF and CNFOX.

(a) (b)

Figure 5-16 SEM images of the fractured surface of PMMA composites having 2 wt. % of (a) CNF and (b) CNFOX. Some nanofibers are circled for comparison.

102 TOS behavior of PMMA-CNFOX composites were also evaluated and compared to PMMA-CNF composites. Figure 5-17 shows the first derivative of mass loss with temperature for PMMA-CNFOX composites.

3.0 PMMA 2.5 1 wt.%

) 2 wt.% 2.0 6 wt.% 10 wt.% 1.5

1.0 Deriv. Weight(%/°C

0.5

0.0 250 275 300 325 350 375 400 425 Temperature (°C)

Figure 5-17 Thermo-oxidative stability of composites of PMMA-CNFOX.

It is evident that CNFOX also improved TOS of the polymer matrix like it was seen previously in the case of untreated CNF (Figure 5-6 and Table 5-1). However, T2 was shifted remarkably to higher values with 6 and 10 wt. % of CNFOX as it can be seen in Figure 5-17 and Table 5-3. An increase in T2 of about 70 °C with respect to that of unfilled PMMA could be observed at 6 wt. % of CNFOX content. On the other hand, the onset of oxidative decomposition, T1, barely changed in the presence of 6 wt. % of CNF, while the presence of the same amount of CNFOX increased that value by 15 °C.

103 Table 5-3 Thermo-oxidative degradation of PMMA and composites prepared in chaotic mixer with different CNFOX contents.

T (°C) PMMA 1 wt.% 2 wt. % 6 wt. % 10 wt. %

T1 306 293 293 323 313

T2 324 324 344 397 392

T3 355 393 395 406 404

This indicates a much better interaction between functionalized carbon nanofibers and polymer chains. Functional groups such as those carrying atomic oxygen on the carbon nanofiber surface may have interacted via dipole-dipole forces with carbonyl groups present in the PMMA macromolecules, therefore restricting their motion upon

156 heating, as it has been reported for PMMA filled with TiO2 or Fe2O3. Additionally, enhancement of TOS in CNFOX composites can be explained on the basis of the carbon nanofibers forming a barrier, thus hindering the transport of degradation products from solid to the gas phase. This argument has been used to explain similar behavior in polymer composites of nanoclay157 and carbon nanotubes.158

Figures 5-18(a) and 5-18(b) show the storage modulus and tan δ of PMMA-

CNFOX composites respectively. These composites showed an increase of E’ at room and elevated temperatures for most of CNFOX compositions. Similar increase was observed in the case of Tg values extracted from the tan δ plot in Figure 5-18(b). This clearly indicates a favorable interaction between fibers and the polymer matrix. Such additional restriction in the chains mobility influenced the TOS values, as it was previously discussed.

104 2.0 (a) 1.6

1.2 PMMA 0.5 wt.% E'(GPa) 0.8 1 wt.% 2 wt.% 0.4 4 wt.%

0.0 20 40 60 80 100 120 140 Temperature (°C)

(b) 1.6 PMMA 0.5 wt.% 1 wt.% 1.2 2 wt.%

δ 4 wt.%

tan 0.8

0.4

0.0 100 110 120 130 140 Temperature (°C)

Figure 5-18 Dynamic mechanical properties of PMMA-CNFOX composites prepared by chaotic mixing. (a) Storage modulus and (b) tan δ.

Values of E’, and tan δ of PMMA composites containing 4 wt. % of CNF and

CNFOX are compared in Figure 5-19. As it can be seen, both fibers enhanced E’ at low temperatures, however enhancement is more prominent in the case of PMMA-CNFOX composites. The storage modulus at temperatures around the Tg of PMMA (120-130 °C), showed reinforcement only in the case of functionalized nanofibers, as is seen in Figure

5-19(a). Figure 5-19(b) shows the variation of tan δ with temperature for the two types of 105 composites. As it can be seen, Tg, determined from the maximum of the tan δ curve, only increased for the PMMA-CNFOX composite. This can be attributed to a much stronger polymer-fiber interaction which in turn was responsible for hindering the macromolecules mobility and therefore increasing E’ and Tg, as it was discussed previously and as it has been described for epoxy-vapor-grown graphite nanofibers.135

2 (a) 1.6 PMMA-CNFOX 4 wt.%

1.2 PMMA

0.8

0.4 PMMA-CNF 4 wt.% 0 20 40 60 80 100 120 140 Temperature (°C)

1.6 PMMA (b)

1.2 δ

tan 0.8 PMMA-CNF 4 wt.% 0.4 PMMA-CNFOX 4 wt.% 0 90 100 110 120 130 140 Temperature (°C)

Figure 5-19 Dynamic mechanical properties of PMMA with 4 wt.% of CNF and CNFOX prepared in a chaotic mixer (a) Storage modulus, E’, and (b) tan δ.

Composites of PMMA and CNFOX, although showed better dispersion, they possessed poor electrical conductivity as compared to composites with CNF (Figure 5-

20). It is seen that conductive composites were obtained with just 2 wt. % of CNF, while

106 materials made with CNFOX were still insulators at that composition. Approximately 6 wt. % of CNFOX was needed to make these composites conductive enough to have potential in electro-static discharge applications. This dramatic decrease in electrical conductivity seen with CNFOX can be attributed to improvement of fiber dispersion and agglomerates reduction.

1.E+01101

1.E-0110-1

-3 CNF

(S/cm) 1.E-0310 σ 1.E-0510-5

-7 1.E-0710

-9 1.E-0910 CNFOX -11 Volume conductivity, 1.E-1110

-13 1.E-1310 0246810 Carbon nanofiber content (wt %)

Figure 5-20 Volume electrical conductivity of PMMA-CNF and PMMA-CNFOX materials prepared by chaotic mixing.

It has been stated previously93,94, that the dispersion of agglomerates and its distribution by chaotic mixing play an important role in obtaining highly conductive materials. It has also been determined in this work that the electrical conductivity in

PMMA-CNF composites is strongly affected by parameters like mixing time and shear force (Sections 5.1 and .5.2 in this Chapter). Under the current processing conditions, the most likely reason why the electrical conductivity dropped in composites of CNFOX is two-fold: (i) dispersion was enhanced i.e., more fibers are surrounded by polymer,

107 thereby limiting the formation of conductive networks, and (ii) reduction of number and size of the nanofiber agglomerates, thereby increasing the agglomerate-agglomerate distance. Similar findings were reported for oxidized carbon nanofibers and vinyl ester polymer – composites by Xu et al.88 These authors determined that even with up to 15 wt. % of nitric acid-treated carbon nanofibers the composite materials exhibited electrical no electrical conductivity. They concluded that an oxygen-rich surface and the resin acted as electronic barrier among the carbon nanofibers. Additionally, as it was mentioned in earlier (Chapter II, Section 2.2.5) carbon nanofibers lose graphitic character upon oxidation, and therefore a reduction of their intrinsic electrical conductivity can be expected. An attempt to measure the bulk electrical conductivity of both CNF and

CNFOX fibers using a makeshift conductivity tester was carried out in our laboratory.

The results shown in Table 5-4 represent the average of 7 measurements of the bulk conductivity of CNF and CNFOX powder.

Table 5-4 Bulk electrical conductivity of CNF and CNFOX.

Carbon nanofiber Bulk conductivity (S/cm)

CNF 0.21 ± 0.03

CNFOX 0.033 ± 0.007

It was found that CNF is approximately 7 times more conductive than CNFOX.

Therefore, the dramatic reduction in electrical conductivity in PMMA-CNFOX composites seen in Figure 5-20 can be due to a combined effect of less graphitic carbon

108 nanofibers, and a much better dispersion of nanofibers during chaotic mixing. However, it is acknowledged that the dispersion of the nanofibers and the disintegration of their agglomerates is probably the main cause for the reduction of electrical conductivity in

PMMA-CNFOX composites.

5.4. Summary

This chapter showed that chaotic mixing can be used effectively in the production of electrically conductive polymer composites with less carbon nanofiber content than with a commercial internal mixer. It was also shown that conductive networks are formed early due to orientation of fibers pulled out of the agglomerates. In addition, dispersed fibers were much less damaged due to low shear conditions. A sketch as presented in Figure 5-21 depicts how agglomerates were formed in PMMA-CNF composites during mixing in chaotic or Brabender Plasticorder mixers. It is noted that

CNFs agglomerates were made of highly entangled fibers with high aspect ratios (Figure

5-21(a)). Due to low shear force in the chaotic mixer, long fibers were dispersed and aligned while others were pulled out of the bundles. These in turn, made contact with neighboring agglomerates (Figure 5-21(b)), and formed conductive networks. Meanwhile, carbon nanofibers in materials prepared in Brabender Plasticorder experienced better dispersion. This delayed electrical percolation to higher fiber volume fractions in comparison to materials processes in the chaotic mixer.

109 (a)

BUNDLES OF FIBERS

(b) (c) CHAOTIC MIXING INTERNAL MIXING (partial unwinding)

Figure 5-21 Sketch of the carbon nanofibers morphology (a) as-received, (b) after chaotic mixing, and (c) after Brabender plasticorder mixing.

The study also showed that electrical conductivity is sensitive to the time of mixing, especially at nanofiber loadings around percolation thresholds. On the other hand, surface treatment of carbon nanofibers resulted in polymer composites with improved thermal and thermo-mechanical properties but with high electrical resistivity. The fiber interface was more compatible with the polymer due to the added functionality (-CHO, -

OH, -COOH, -COOR) on the oxidized nanofibers. Thermal and mechanical properties improved in the case of oxidized fibers due to an improvement in the fiber-polymer interface and fiber dispersion. Better dispersion in turn led to increased electrical resistivity of those composites when compared to materials having untreated carbon nanofibers. Chaotic mixing provides additional versatility in the production of PMMA- carbon nanofiber composites.

110 CHAPTER VI

TPU-CARBON NANOFIBER COMPOSITES

In order to investigate the influence of carbon nanofibers on TPU polymers, two completely different systems were chosen: one with amorphous soft segments, TPU23, and another with crystallizable (at room temperature) soft segments, TPU33. This

Chapter covers the results of the investigation of these two TPU formulations and their composites of three different types of carbon nanofibers.

6.1. TPU23-carbon nanofiber composites

The effect of fiber surface chemistry on dispersion was thoroughly evaluated by comparing composites of CNF with those of CNFOX and CNFOL. TPU23 prepolymer,

BDO and carbon nanofibers were poured in the chaotic mixer 2 to fill up a volume of 70 cm3 and the mixing torque was recorded. Figure 6-1 shows the values of mixing torque for TPU23 with 0.5 wt. % of CNF, CNFOX, and CNFOL. A sudden increase in the value of the torque is seen with carbon nanofibers in about 5 minutes of mixing. This sudden increase, however, is not seen for TPU23 alone. It can be noticed that the torque

111 7 CNFOL 0.5 wt. % 6

5

4 CNFOX 0.5 wt. %

3 Torque (N-m)

2

1 CNF 0.5 wt. % TPU23

0 0 500 1000 1500 2000 2500 3000 3500 4000 4500 5000 5500 6000 Time (sec)

Figure 6-1 Mixing torque of TPU23 with different types of carbon nanofibers.

value is higher when treated fibers are used, e.g., CNFOX and CNFOL. This is a ramification of the affinity existing between TPU23 and the treated fibers. Such affinity was most manifested when the polyol-treated CNF was used.

6.1.1. Morphological analysis

TEM images of thin sections of TPU23 composites with 3 wt.% of CNF, CNFOX and CNFOL are shown in Figure 6-2. Good dispersion and alignment of the carbon nanofibers can be observed for both types of composites.

112 (b) (a)

Flow 2 μm Flow 2 μm

(c)

1 μm Flow

Figure 6-2 TEM images of TPU23 composites with 3 wt. % of (a) CNF, (b) CNFOX, and (c) CNFOL. Arrows indicate the flow direction. Scale bars are 2 μm in (a) and (b) and 1 μm for (c).

In each case, good dispersion was obtained due to long mixing time during preparation of these composites in the chaotic mixer. One important feature of 2D chaotic mixing is the possibility of fiber alignment, which can be seen in Figure 6-2 as well. In

Chapter V, alignment of CNF in PMMA along the flow direction was observed. In the present case, flow-induced alignment of the dispersed fibers is also seen. Figure 6-2(c) shows that CNFOL fibers were irregular in shape and were more damaged after the oxidation, chemical treatment, and the mixing step.

Any attempts to use TEM images in Figure 6-2 to interpret the effect of processing on nanofiber aspect ratio, would give us wrong information. First, it is likely

113 that nanofibers were cut during sectioning due to their pronounced curved structure159 and that only portions of the fibers can be inspected in TEM. Second, only small portions of mixed materials are captured by TEM micrographs. To circumvent this, thin films were cast to provide specimens with more complete appearance of nanofibers, thereby allowing the observation of the real length of fibers after processing. Figure 6-3 shows

TEM micrographs of such cast thin films of CNFOX 3 wt.% composite. It is seen that after processing, nanofibers were of much reduced length compared to the original length.

It is seen that after processing, nanofibers with typical lengths of 4 μm represents a 96% length reduction from the original fiber size (~100μm).

Figure 6-3 TEM image of a cast film of TPU-CNFOX composite with 3 wt.% nanofibers.

6.1.2. Thermal degradation

The thermal stability of TPU23 composites was evaluated by TGA. Figures 6-4(a) and Figure 6-4(b) show the mass loss and its first derivative versus temperature respectively. It can be seen that T1 (~278 °C) did not change much while T2 of TPU

114 increased with the addition of merely 0.5 wt. % of CNF, CNFOX or CNFOL (from

346 °C to 380 °C).

100 2.0 (a) TPU TPU (b) CNF 0.5 wt.% CNF 0.5 wt% 80 CNFOX 0.5 wt.% 1.5 CNFOX 0.5 wt.% CNFOL 0.5 wt.% CNFOL 0.5 wt.% 60 1.0 40 Weight (%) Weight 0.5

20 Deriv. Weight (%/°C)

0 0.0 20 120 220 320 420 520 620 200 250 300 350 400 450 500 Temperature (°C) Temperature (°C)

Figure 6-4 Thermogravimetric behavior in N2 gas of TPU composites with 0.5 wt. % of CNF, CNFOX and CNFOL (a) Mass loss (b) First derivative of mass loss.

Thermal degradation of TPUs is a complicated process as hard and soft segments respond differently to elevated temperatures. It has been found that thermal degradation of TPU occurs in two stages.118,127 In Stage I, decomposition of the hard segments occurs, involving dissociation of urethane to the original polyol and isocyanate. In Stage II, depolycondensation and polyol degradation in the soft segments take place. Therefore,

Stage I can be associated with the hard segments and T1, while Stage II can be related to the soft segments and T2. As it can be noticed from Figure 6-4, T2 was more affected by the presence of carbon nanofibers, which suggests that the nanofibers interacted more with the soft segment chains in TPU. Similar results were observed for TPU-MWNT composites.118 As in the case of PMMA-CNF composites (Chapter V), improvement in the thermal stability of TPU may be explained on the basis of stabilization of CNF- bonded macroradicals.155

115 6.1.3. Molecular weight and molecular weight distribution

Figures 6-5(a) and 6-5(b) show the weight average molecular weight (Mw) and the polydispersity index (Mw/Mn) for TPU23 and selected composites.

160 14 (a) (b) 140 12

120 10 (in thousand Daltons) thousand (in

W 100

M 8 80 6 60 4 40 Polydispersity index, index, Mw/Mn Polydispersity

20 2

0 0 TPU23 TPU23-CNF05 TPU23-CNFOX 05 TPU23-CNFOL05 Weight averageWeight molecular weight, TPU23 TPU23-CNF05 TPU23-CNFOX 05 TPU23-CNFOL05

Figure 6-5 (a) Molecular weight and (b) molecular weight distribution of TPU23 and its corresponding 0.5 wt. % composites.

Neat TPU23 showed higher MW than the corresponding 0.5 wt.% composites of different types of carbon nanofibers. It is clear that the presence of very small quantities of carbon nanofibers affected the chain growth in TPU. Shorter macromolecular chains generated narrower polydispersity index values for the composites, except in the case of

0.5 wt. % CNFOL, as shown in Figure 6-5(b).

6.1.4. Thermo-mechanical behavior

Figures 6-6, 6-7, and 6-8 show plots for E’, E” and tan δ of TPU23-CNF, TPU23-

CNFOX, and TPU23-CNFOL composites respectively.

116 1.E+10 (a) 0 wt.% 4.6 0.5 wt.% 4.4 1 wt.% 1.E+09 4.2

E'(GPa) 3 wt.% 4.0

3.8 -90 -85 -80 -75 -70 -65 -60 1.E+08 Temperature (°C) E'(Pa)

1.E+07

1.E+06 -90 -70 -50 -30 -10 10 30 50 Temperature (°C)

1.E+09 (b) 0 wt.% 0.5 wt.% 1.E+08 1 wt.% 3 wt.%

1.E+07 E"(Pa)

1.E+06

1.E+05 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

1.2 (c) 0 wt.% 0.5 wt.% 1 wt.% 3 wt.% 0.8 δ tan tan

0.4

0.0 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

Figure 6-6 Thermo-mechanical properties of TPU23-CNF composites: (a) storage modulus, (b) loss modulus and (c) tan δ.

117 1.E+10 (a) 0 wt.%

5.2 0.5 wt.%

1.E+09 4.8 1 wt.%

4.4 E'(GPa) 3 wt.% 4.0

1.E+08 3.6 -90 -85 -80 -75 -70 -65 -60 Temperature (°C)

E'(Pa) 1.E+07

1.E+06

1.E+05 -90 -70 -50 -30 -10 10 30 50 Temperature (°C)

1.E+09 (b) 0 wt.% 0.5 wt.% 1.E+08 1 wt.% 3 wt.%

1.E+07 E"(Pa)

1.E+06

1.E+05 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

1.2 (c) 0 wt.% 0.5 wt.% 1 wt.% 3 wt.% 0.8 δ tan

0.4

0.0 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

Figure 6-7 Thermo-mechanical properties of TPU23-CNFOX composites: (a) storage modulus, (b) loss modulus and (c) tan δ.

118 TPU-CNFOX and TPU-CNFOL composites show a higher increase in the values of E’ in the glassy plateau (-90 °C ≤ T ≤ -60 °C) compared to TPU-CNF composites, as it can be noticed from the insets in Figures 6-6(a), 6-7(a), and 6-8(a).

1.E+10 (a) 0 wt.% 5.6 0.5 wt.% 5.2 1 wt.% 1.E+09 4.8

E'(GPa) 3 wt.% 4.4

4.0

3.6 -90 -85 -80 -75 -70 -65 -60 1.E+08 Temperature (°C) E'(Pa)

1.E+07

1.E+06 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

1.E+09 (b) 0 wt.% 0.5 wt.% 1.E+08 1 wt.% 3 wt.%

1.E+07 E"(Pa)

1.E+06

1.E+05 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

1.2 0 wt.% (c) 0.5 wt.% 1 wt.% 3 wt.% 0.8 δ tan tan

0.4

0.0 -100 -80 -60 -40 -20 0 20 40 Temperature (°C)

Figure 6-8 Thermo-mechanical properties of TPU23-CNFOL composites: (a) storage modulus, (b) loss modulus and (c) tan δ.

119 It can also be observed that the values of E’ around the glass transition are higher for composites with CNFOX and CNFOL. The increase in E’ with fiber content comes from the fact that carbon nanofibers reinforced TPU and then increased its stiffness at lower temperatures. This has been observed for TPU with SWNT and MWNT.118 A more favorable interaction of CNFOX and CNFOL with TPU soft segment chains can be revealed from the high values of E’ around the glass transition in Figures 6-7(a) and 6-

8(a). On the other hand, a drastic reduction of E’ at room temperature for TPU-CNFOX 3 wt. % can be noticed in Figure 6-7(a). Values of E’ and E” at room temperature for TPU-

CNF, TPU-CNFOX, and TPU-CNFOX composites are presented in Table 6-1. It can be observed from this Table that while E’ and E” of TPU23 increased with CNF content at

25 °C, they decreased with 3 wt. % CNFOX. On the other hand, CNFOL showed a similar behavior to that of CNF composites.

Table 6-1 Storage, E’, and loss, E”, modulus at room temperature for TPU23 composites prepared in the chaotic mixer.

CNF E’ (MPa) at 25 °C E” (MPa) at 25 °C

content CNF CNFOX CNFOL CNF CNFOX CNFOL (wt. %)

0 3.3 3.3 3.3 0.43 0.43 0.43

0.5 4.1 3.3 4.2 0.49 0.47 0.57

1 3.7 4.3 3.9 0.52 0.53 0.59

3 5.5 0.9 3.2 0.61 0.13 0.86

120 A shift to higher values of the soft segment glass transition temperature, Tg,SS can be appreciated from E” and tan δ plots in Figures 6-7(b), 6-7(c), and 6-8(b), 6-8(c) for

CNFOX and CNFOL composites respectively. Values of Tg,SS are also presented in Table

6-2. This phenomenon has been reported for TPU composites of acid treated MWNT in contrast to untreated MWNT.116 This reflects a reduction of the extent of phase separation.

This may be the reason for the decrease in storage modulus value at room temperature for

3 wt.% CNFOX (Table 6-1). A reduction of the tan δ value and gradual increase of Tg,SS values after 0.5 wt. % with CNFOX and CNFOL in Figures 6-7(b,c) and 6-8(b,c), may be explained in terms of increased carbon nanofibers agglomerate size at high nanofiber concentration, as it was pointed out by Kwon and Kim.116

Table 6-2 Glass transition temperature of soft segments, Tg,SS from E” and tan δ plots of TPU23 composites prepared in the chaotic mixer.

CNF Tg,SS (°C) from E” plot Tg,SS (°C) from tan δ plot

content CNF CNFOX CNFOL CNF CNFOX CNFOL (wt. %)

0 -53.6 -53.6 -53.6 -40.9 -40.9 -40.9

0.5 -53.6 -50.0 -45.9 -39.9 -27.9 -31.6

1 -55.2 -53.6 -52.8 -42.0 -32.6 -34.7

3 -52.8 -55.9 -52.8 -38.9 -34.7 -42.0

121 6.1.5. Electrical conductivity

Although the degree of nanofiber dispersion increased due to fiber treatment,

TPU23 composites did not show any improvement in electrical conductivity at least up to

3 wt. % loading. As it was observed from TEM micrographs for these materials (Figures

6-2 and 6-3), a high level of dispersion and reduction in the aspect ratio did not allow overlapping of the fibers and therefore establishment of conductive networks. A specimen of TPU23 with 3 wt. % CNF was prepared with much shorter mixing time – that is, 10 minutes instead of 90 min – in order to obtain higher electrical conductivity.

However, this material showed similar low values of electrical conductivity (10-14 S/cm).

In view of this, higher nanofiber content is required in order to obtain electrically conductive TPU23 composite materials.

6.1.6. Tensile properties

Tensile properties of TPU23-CNF, TPU23-CNFOX, and TPU23-CNFOL composites were investigated and the results are presented in Figure 6-9. It can be seen from Figure 6-9(a) that the stress at break did not change by the presence of CNF, while with CNFOX, a maximum value of stress was obtained with 1 wt.% CNFOX, and the 3 wt. % composite exhibited lower stress even compared to unfilled TPU. CNFOL yielded lower tensile stress values. The values of strain at break increased up to 1 wt. % nanofibers content for both CNF and CNFOX (Figure 6-9(b)), which was not the case for

CNFOL composites. A possible explanation for the lower tensile strength and strain

122 observed in the case of TPU-CNFOL composites may be based on the fibers morphology observed in Figure 6-2(c), in which they seemed to have been damaged after several chemical treatments.

14 (a) 12

10

8

6

Stress at break (MPa) 4 CNF 2 CNFOX CNFOL 0 0123 Carbon Nanofiber content (wt. %)

2600 (b) 2200

1800

1400

1000 Strain at breakStrain (%) CNF 600 CNFOX CNFOL 200 0123 Carbon Nanofiber content (wt. %)

2 CNF (c) CNFOX 1.6 CNFOL

1.2

0.8

Young's Modulus (MPa) 0.4

0 0123 Carbon Nanofiber content (wt. %)

Figure 6-9 Tensile properties TPU composites of CNF, CNFOX, and CNFOL (a) Stress at break (b) strain at break, and (c) Young’s modulus.

123 On the other hand, and as with most TPU-CNF and TPU-CNT nanocomposites reported in literature114-116,118,119,121,123,160, the value of Young’s modulus for both types of composites increased with fiber content (Figure 6-9(c)), and no significant difference can be observed between CNF, CNFOX or CNFOL. In summary, the presence of CNF shows typical reinforcement behavior, while the effect of CNFOX is a more complex one, which needs further investigation in terms of FTIR and DSC data. CNFOL increased the modulus of TPU but decreased the tensile strength and strain due to a more damaged graphitic structure.

6.1.7. Degree of phase separation

The degree of phase separation in segmented polyurethanes can be estimated from infrared spectroscopy. A typical infrared spectrum of a TPU composite is shown in

Figure 6-10, with the principal functional groups assigned to their respective peaks.57 The carbonyl peak was deconvoluted by curve fitting into overlapping peaks as presented in

Figure 6-11. A similar procedure described by Xia et al.118 was followed in order to obtain values for R and DPS (Chapter III, Section 3.3.10). Figure 6-12 shows that the degree of phase separation decreased with the carbon nanofiber content. The values of R and DPS for CNFOX and CNFOL composites are higher than those of CNF composites.

The additional hydrogen bonding in TPU23-CNFOX, and TPU23-CNFOL composites may have come from chemical interactions of the urethane groups with functional groups on CNFOX and CNFOL fibers. Several researchers have reported results on phase separation in TPU-montmorillonite composites.154,161,162 Tien and Wei154 found that

124 hydroxyl groups on silicate layers could form hydrogen bonding either with hard and soft segments in TPU. They concluded that although silicate layers reinforced TPU, they also reduced hydrogen bonding in the hard segment region. Pattanayak161-163 and Gorrasi164 reported a reduction in hydrogen bonding in TPU in the presence of nanoclay.

ν(C=C) benzene ring

(C=O) ν(N-H) ν free N- H bonded

ν(N-H) bonded N-H

ν (C=O) free Transmittance (%) Transmittance

νa(CH2) νs(CH2) δ(N-H) + ν(C-N) ν = stretching (a = asymmetric; s = symmetric) ν(CH2-O-CH2) δ = bending aliphatic ether

3800 3300 2800 2300 1800 1300 800 Wavenumber (cm-1)

Figure 6-10 Typical infrared spectra of a TPU23 carbon nanofiber composite with principal peak assignments (3 wt. % CNFOL).

ν free C=O ~1730 cm-1 ν bonded C=O ~1700 cm-1

Figure 6-11 Deconvolution of the carbonyl peak into free and hydrogen-bonded. 125 2.5 (a)

2

1.5

1

Hydrogen bonding index, R 0.5

0 Blank CNF1 CNFOX1 CNFOL1 CNF3 CNFOX3 CNFOL3 Material

80 (b) 70

60

50

40

30

20

Degree of Phase Separation, (%)DPS 10

0 Blank CNF1 CNFOX1 CNFOL1 CNF3 CNFOX3 CNFOL3 Material

Figure 6-12 (a) Hydrogen bonding index and (b) degree of phase separation of TPU23 composites.

Chen160 reported segmented polyurethane composites with SWNTs, and observed good mechanical properties with 0.5 wt. % nanotube content. However, further addition of SWNT led to reduction in tensile strength and elongation at break. These authors attributed it to several possible reasons such as altered microphase morphology of TPU, excessive polymer-filler interaction, and inhibition of morphological changes during mechanical deformation of TPU. The same arguments can be applied in our study to explain the reduction of tensile strength and degree of phase separation values presented

126 previously in Figure 6-9(a) and Figure 6-12(b) respectively for TPU composites having 3 wt. % of CNFOX and CNFOL. At this point, we require more information on the behavior of the hard segment domains that can be obtained from the investigation of the thermal transitions of these composites.

6.1.8. Thermal transitions

DSC traces of TPU23 and its composites are shown in Figure 6-13.

Neat TPU23

CNFOX 0.5 wt.%

CNFOL 0.5 wt.% CNF 0.5 wt.% Heat Flow (mW) CNF 3 wt.%

CNFOX 3 wt.%

CNFOL 3 wt.%

-100 -50 0 50 100 150 200 250 Temperature (°C)

Figure 6-13 Thermal transitions of TPU composites.

MDI-BDO-based TPU materials are known to show multiple endotherms and significant research work has been done to discover the morphological sources behind

127 165-170 those endotherms. The first endotherm, TI, appears at approximately 20-30 °C above the annealing temperature – normally room temperature – and it has been ascribed to a short-range order in the hard domains. A second endotherm, TII observed between

120-200 °C has been attributed either to a long-range ordering in the hard domains or to a microphase mixing occurring between the hard and soft domains. Finally, melting of the hard segment microcrystalline regions produces an endotherm at TIII above 200 °C.

Besides hard segment transitions, soft segments provide a glass transition Tg, SS occurring at -40 °C. Soft segment crystallization is absent in TPU23 due to the presence of a methyl group in PPG polyol which probably prevents crystalline packing of the polyol chains with a 2000 molecular weight.

Two morphological sources have been proposed to explain the first endotherm TI;

(i) disordering of short-range hard segments165 and (ii) hard segment glass transition.166

More recently Chen et al.168 investigated this endotherm and determined that its source could be due to an enthalpy relaxation resulting from physical aging of the amorphous hard segment. From Figure 6-13, it is seen that TI is easily noticeable in composites with

3 wt. % of carbon nanofibers, although the change in slope is more prominent for

CNFOX and CNFOL composites. In view of this, we can infer that high content of functionalized nanofibers may affect the hard segment regions and promote the disorder of such phase. Due to the presence of oxygen-containing functional groups, CNFOX and

CNFOL nanofibers can build up hydrogen bonds with the urethane groups and alter the hard domain morphology. A sketch representing this scenario can be seen in Figure 6-14.

Additional hydrogen bonding in CNFOX and CNFOL composites was already estimated in this work (Figure 6-12). Consequently, TII is practically absent and TIII decreased in

128 TPU-CNFOX3, for example, which reflects the increase of short-range order in this sample upon reduction of the long-order transitions. Although it is conceivable that tensile strength should have improved with CNFOX content, e.g., at 3 wt. %, the disruption of the hard segment domains in TPU was more prevalent, therefore reducing significantly the mechanical strength of the polymer.

Figure 6-14 Sketch of the morphology of hard domains in TPU23 composites with surface treated CNF.

129 6.2. TPU33-carbon nanofiber composites

TPU33 differs from TPU23 in that TPU33 has a crystallizable soft segment phase which plays a more significant role in determining the final properties of TPU. This section covers a thorough characterization of TPU33 system and its composites with CNF and CNFOX prepared in chaotic mixer. One special attribute of TPU33 is the shape memory properties triggered by the action of the melting at about 50 °C of the PCL-based crystalline soft segments. This special feature of TPU33 will be discussed in detail in

Chapter VII. In order to find out the advantages and disadvantages of using chaotic mixing protocol to produce these materials, TPU33-CNF composites were prepared also in Brabender Plasticorder. It was expected that the two-phase morphology of TPU, and the presence of a crystallizable soft segment, would have strong impact on materials prepared by chaotic mixing.

6.2.1. Morphological analysis

Figure 6-15 shows optical micrographs of compressed films of TPU33-CNF1

(Figure 6-15(a)) and TPU33-CNFOX1 (Figure 6-15(b)). It is apparent that like in the case of PMMA-CNFOX composites (Chapter V, Section 5.3), the presence of functional groups on the surface of CNFOX also improved dispersion of carbon nanofibers agglomerates. As it was mentioned in the Experimental Section (Chapter III), TPU33 composites were prepared with a much shorter mixing time (10 min) in comparison to 90 minutes of mixing time for TPU23 composites.

130 (a) (b)

Figure 6-15 Optical micrographs of TPU33 composites prepared in chaotic mixer with 1 wt. % of (a) CNF, and (b) CNFOX.

Such shorter mixing time was used to obtain electrically conductive materials, as it was learned previously (Chapter V, Section 5.2) that electrically conductive composites can be prepared with shorter chaotic mixing times when PMMA was mixed with CNF.

Although the dispersion quality improved in the case of TPU33-CNFOX composites, some agglomerates are still observed.

SEM images of TPU33 composites with 5 wt. % of CNF and CNFOX prepared in the chaotic mixer are shown in Figures 6-16(a) and 6-16(b). In addition, Figure 6-16(c) shows the fractured surface of TPU33-CNF5 prepared in the Brabender Plasticorder.

Figure 6-16 shows different morphologies as function of the type of CNF and the method of preparation. In the case of TPU33-CNF5 prepared by chaotic mixer (Figure 6-16a), aggregates of dry fibers can be seen which corresponds to poor interactions at the fiber- polymer interface. Note that the shear rate operating in the chaotic mixer was low (~4 s-1).

Fibers are seen more dispersed in the case of TPU33-CNF5 composite prepared in the

Brabender Plasticorder mixer. However, in both cases, CNF were not very well wetted by the polymer. Fibers were pulled out of the polymer during the fracture process. On the

131 other hand, Figure 6-16(b) shows no such pulling out of fibers. Only the tips of nanofibers can be observed as white dots in Figure 6-16(b). This indicates good adhesion between CNFOX and TPU33.

(a) (b)

(c)

Figure 6-16 SEM pictures of (a) TPU33-CNF5 prepared in chaotic mixer, (b) TPU33- CNFOX5 prepared in chaotic mixer and (c) TPU33-CNF5 prepared in Brabender Plasticorder.

6.2.2. Thermal degradation

Thermal stability of TPU33-CNF and TPU33-CNFOX composites prepared in the chaotic mixer degradation was investigated, and results are shown in Figure 6-17. T1 and

T2 values for both types of composites showed a similar trend, i.e., reduce at 1 wt. % and

132 then increase with the addition of higher content of CNF. Similar behavior was observed recently by Mondal and Hu120 for a TPU-MWNT system. They observed a decline in the thermal stability of TPU with 0.25 and 0.50 wt. % of functionalized MWNT, and then an increase with higher amounts of MWNT.

325 (a) 320 CNFOX 315

310 CNF 305

Temperature (°C) 300

295

290 01234567 Carbon nanofiber content (wt. %)

390 (b) 385 380 375 370 CNF 365

Temperature (°C) 360 CNFOX 355 350 01234567 Carbon nanofiber content (wt. %)

Figure 6-17 Thermal degradation stability of TPU33-CNF and TPU33-CNFOX composites (a) T1, and (b) T2 values.

This phenomenon was explained based on the dual effect provided by MWNT during heating and degradation of the polymer matrix. MWNT, in the same manner as

CNF and CNFOX, enhances the thermal conductivity of TPU, thus facilitating the

133 thermal degradation at low carbon contents. On the other hand, higher contents of

MWNT, CNF or CNFOX produce a delay in the thermal degradation of TPU33, as it was discussed in this Chapter for TPU23 composites (Section 6.1.2). Similar to TPU23, soft segments in TPU33 were more affected by the presence of CNF or CNFOX.

6.2.3. Thermo-mechanical behavior

The dynamic mechanical properties of TPU33-CNF composites prepared in the chaotic mixer are presented in Figure 6-18. Results for TPU33-CNF composites prepared in the Brabender Plasticorder mixer are shown in Figure 6-19. DMA results for TPU33-

CNFOX composites prepared in the chaotic mixer are presented in Figure 6-20. Several thermal transitions can be determined from these plots. From the values of E’ (Figures 6-

18(a), 6-19(a), and 6-20(a)) it is seen that the glassy plateau modulus is slightly higher for

TPU than those with nanofibers. A drop in the values of E’ versus temperature represents the onset of the glass transition which is more prominent for composites with carbon nanofibers content above 1 wt.%, as seen in Figures 6-18(a), and 6-19(a), and above 3 wt. % in the case of materials with CNFOX, as in Figure 6-20(a). As the rubbery plateau is concerned, TPU33-CNF materials prepared in the Brabender Plasticorder showed the lowest storage modulus levels. On the other hand, all the composites experienced a decline in E’ values between 25 °C and 50 °C due to the melting of PCL crystals. Finally, the composites showed higher E’ values than pristine TPU above 50 °C. The glass transition behavior is seen from plots of E” vs. T shown in Figures 6-18(b), 6-19(b), and

6-20(b), and from plots of tan δ vs. T in Figures 6-18(c), 6-19(c), and 6-20(c). In the case

134 of neat TPU33, Tg transition is seen to be broad and it was not possible to assign a value to the glass transition temperature. This indicates that less amorphous polymer chains were available to participate in the glass transition in unfilled TPU33, indicating a higher degree of crystallinity. Above 1 wt. % CNF, the presence of more amorphous chains is reflected by an increase in the values of tan δ and E”. In contrast, TPU33-CNFOX composites showed the onset of distinct glass transition of soft segment chains until 3 wt. % CNFOX. If we compare the Tg values obtained from E” and tan δ curves, we notice that Tg values were higher for materials prepared in Brabender Plasticorder than in the chaotic mixer, as seen in Table 6-3. This is due to the presence of more hard segments mixed within the amorphous soft segments in the case of the Brabender Plasticorder than those produced in the chaotic mixer.

Table 6-3 Tg values obtained from E” and tan δ plots for TPU33-CNF composites prepared in the chaotic mixer and Brabender Plasticorder.

CNF content Tg (°C) from E” plot Tg (°C) from tan δ plot

(wt. %) Chaotic Brabender Chaotic Brabender

0 ND -59 ND ND

1 -67 -61 ND ND

3 -66 -59 -62 -54

5 -68 -55 -64 -44

7 -66 -53 -62 -42

ND: non-detectable

135 1.E+10 (a) 0 WT.% 1 WT.% 3 WT.% 1.E+09 5 WT.% 7 WT.% E'(Pa)

1.E+08

1.E+07 -100 -75 -50 -25 0 25 50 75 100 Temperature (°C)

1.E+09 (b) 0 WT.% 1 WT.% 1.E+08 3 WT.% 5 WT.% 7 WT.%

1.E+07 E"(Pa)

1.E+06

1.E+05 -100 -75 -50 -25 0 25 50 75 100 Temperature (°C)

0.3 (c) 0 WT.% 1 WT.% 3 WT.% 5 WT.% 0.2 7 WT.% δ tan tan

0.1

0 -120 -95 -70 -45 -20 5 30 55 80 105 130 Temperature (°C)

Figure 6-18 Thermo-mechanical properties of TPU33-CNF composites prepared in chaotic mixer: (a) storage modulus, (b) loss modulus and (c) tan δ.

136 1.E+10 (a) 0 WT.% 1 WT.% 1.E+09 3 WT.% 5 WT.% 7 WT.% 1.E+08 E'(Pa)

1.E+07

1.E+06 -100 -50 0 50 100 Temperature (°C)

1.E+09 (b) 0 WT.% 1 WT.% 1.E+08 3 WT.% 5 WT.% 7 WT.% 1.E+07 E"(Pa)

1.E+06

1.E+05 -100 -50 0 50 100 Temperature (°C)

0.4 (c) 0 WT.% 1 WT.% 0.3 3 WT.% 5 WT.% 7 WT.% δ 0.2 tan

0.1

0 -130 -80 -30 20 70 120 Temperature (°C)

Figure 6-19 Thermo-mechanical properties of TPU33-CNF composites prepared in Brabender mixer: (a) storage modulus, (b) loss modulus and (c) tan δ.

137 1.E+10 (a) 0 WT.% 1 WT.% 3 WT.% 1.E+09 5 WT.% 7 WT.% E'(Pa) 1.E+08

1.E+07 -100 -75 -50 -25 0 25 50 75 100 Temperature (°C)

1.E+09 (b) 0 WT.% 1 WT.% 3 WT.% 5 WT.% 1.E+08 7 WT.% E"(Pa) 1.E+07

1.E+06 -100 -75 -50 -25 0 25 50 75 100 Temperature (°C)

0.35 (c) 0 WT.% 0.3 1 WT.% 3 WT.% 0.25 5 WT.% 7 WT.% 0.2 δ

tan tan 0.15

0.1

0.05

0 -120 -95 -70 -45 -20 5 30 55 80 105 130 Temperature (°C)

Figure 6-20 Thermo-mechanical properties of TPU33-CNFOX composites prepared in chaotic mixer: (a) storage modulus, (b) loss modulus and (c) tan δ.

138 6.2.4. Electrical conductivity

Figures 6-21(a) and 6-21(b) show how the values of volume electrical conductivity and surface resistivity changed with CNF content for composites of TPU33 prepared in the chaotic mixer and Brabender Plasticorder, respectively. Figure 6-21(c) shows the electrical conductivity values for TPU33-CNFOX composites prepared in the chaotic mixer. It is seen that an electrical percolation occurs at 3 wt. % CNF as observed from the plot in Figure 6-21(a) for materials prepared in the chaotic mixer. In contrast, materials prepared in the Brabender Plasticorder did not show electrical conductivity up to 7 wt. % CNF content as seen in Figure 6-21(b). The lack of electrical conductivity in materials prepared in Brabender Plasticorder can be attributed to two effects: (i) The level of dispersion is much better than in the chaotic mixer as it was observed in the case of

PMMA-CNF materials (Chapter V). (ii) The fiber length reduced more due to high shear at the tips of the rotors in the Brabender Plasticorder, as it was also discussed in Chapter

V. These two factors play important roles in determining the final electrical conductivity of polymer-CNF composites. A third important factor in this regard is the surface chemistry of the graphitic structure of the carbon materials. The influence of this factor on the electrical conductivity of TPU33-CNF was investigated and it is shown in Figure

6-21(c). In this case, the electrical percolation occurred at around 4 wt.% CNFOX.

Composites with 5 and 7 wt.% CNFOX showed volume electrical conductivity with the same order of magnitude as their respective CNF counterparts (~10-5 S/cm), although in the latter case, percolation was obtained with 3 wt. % CNF. This result is in stark contrast to what was observed in the case of PMMA-CNFOX composites (Chapter V).

139 1.E+01 1.E+16 (a) 1.E-01 1.E+14

(S/cm) 1.E-03 σ

1.E-05 1.E+12

1.E-07 1.E+10 1.E-09 1.E+08 Volume conductivity, conductivity, Volume

1.E-11 Surface Resistivity (Ohm/square)

1.E-13 1.E+06 02468 Carbon nanofiber content (wt %)

1.E+01 1.E+16 (b) 1.E-01 1.E+14

(S/cm) 1.E-03 σ 1.E+12 1.E-05

1.E-07 1.E+10

1.E-09 1.E+08 Volume conductivity, conductivity, Volume

1.E-11 Resistivity (Ohm/square) Surface

1.E-13 1.E+06 02468 Carbon nanofiber content (wt %)

1.E+01 1.E+16 (c) 1.E-01 1.E+14

(S/cm) 1.E-03 σ

1.E-05 1.E+12

1.E-07 1.E+10 1.E-09 1.E+08 Volume conductivity,

1.E-11 (Ohm/square) Resistivity Surface

1.E-13 1.E+06 02468 Carbon nanofiber content (wt %)

Figure 6-21 Electrical conductivity of TPU33-CNF composites prepared in (a) chaotic mixer and (b) Brabender Plasticorder and (c) TPU-CNFOX composites prepared in chaotic mixer.

This can be attributed to differences in the nature of the polymer. Note that

PMMA is a completely amorphous material, while TPU33 is a semi-crystalline one. It has been known that carbon composites of amorphous and semi-crystalline polymers

140 have different electrical conduction behavior.171,172 Semi-crystalline polymers are known to require smaller amounts of conductive fillers to percolate. In this case, carbon agglomerates are found preferentially in the amorphous region, thus increasing the effective concentration for percolation. A similar observation was made recently in polyamide 6-polypropylene blends prepared in a chaotic mixer.27 It was also observed in the case of TPU23 materials (Section 6.1.5), that 10 minutes of mixing to prepare 3 wt. %

CNF composite did not yield electrical conductivity, e.g., compare conductivity values of

10-14 S/cm for TPU23-CNF3 vs. 10-8 S/cm obtained in TPU33-CNF3.

6.2.5. Tensile properties

The tensile properties of TPU33-CNF and TPU33-CNFOX composites prepared in the chaotic mixer and the Brabender Plasticorder were evaluated at room temperature.

Because this type of materials shows shape memory effect triggered by the melting of

PCL crystals at around 50-60 °C, it is important that the tensile properties are evaluated at such temperatures. At above the crystalline melting temperatures, the mechanical properties rely only on the hard domains and the possible reinforcement provided by the fillers.141

Tensile strength, tensile strain and Young’s modulus of TPU-CNF composites prepared in the chaotic mixer and Brabender Plasticorder, as well as those of TPU33-

CNFOX materials produced in the chaotic mixer, are shown in Figure 6-22. In general, neat TPU33 and CNF composites prepared in Brabender Plasticorder showed higher tensile stress and strain values but lower Young’s modulus than those prepared in the

141 chaotic mixer, as shown in Figures 6-22(a), 6-22(b) and 6-22(c). This behavior indicates that TPU33 materials produced in the Brabender Plasticorder are more elastic while the materials produced in the chaotic mixer are stiffer. It can be seen from Figure 6-22 (c), for example, a difference of more than 75 MPa in the Young’s modulus of neat TPU33 produced in the chaotic mixer compared to the same material produced in the commercial mixer.

Turning now our attention to TPU33-CNF composites fabricated by chaotic mixing, we notice an increase of both tensile stress and strain in comparison to neat

TPU33 also synthesized in the chaotic mixer. A similar behavior is seen for TPU33-

CNFOX materials prepared in the chaotic mixer as shown in Figures 6-22(a) and 6-22(b).

On the other hand, the trend for Young’s modulus shown in Figure 6-22(c) does not indicate any contribution to stiffness by the addition of CNF or CNFOX. On the contrary, the modulus decreased for all compositions of CNF and CNFOX prepared either in the chaotic or commercial mixer. This drop seems to be more dramatic when

CNF was used. This result is in stark contrast to how Young’s modulus varied with CNF content in TPU23 composites, as discussed in Section 6.1.6. The decrease of the modulus in TPU33 corresponds to the trend observed previously with the value of storage modulus at room temperature observed in Section 6.2.3. As TPU33 contains a crystallizable soft segment which melts above room temperature (~50 °C), this morphological feature predominates in determining the modulus of the materials at room temperature. In view of this, it is necessary to melt the PCL soft segment crystals and then measure the tensile properties to determine their effect on the Young’s modulus at room temperature. Kim et al.65 investigated the tensile modulus at room temperature and at 65 °C of a series of

142 45 (a) TPU33-CNF CHAOTIC TPU33-CNF BRABENDER 40 TPU33-CNFOX CHAOTIC

35

30

25 Stress at break (MPa)

20

15 012345678 Carbon nanofiber content (wt. %)

1400 (b) TPU33-CNF CHAOTIC TPU33-CNF BRABENDER 1200 TPU33-CNFOX CHAOTIC

1000

800 Strain at break (%) 600

400 02468 Carbon nanofiber content (wt. %)

225 (c) TPU33-CNF CHAOTIC 200 TPU33-CNF BRABENDER TPU33-CNFOX CHAOTIC 175 150 125 100 75

Young's Modulus Modulus (MPa) Young's 50 25 0 02468 Carbon nanofiber content (wt. %)

Figure 6-22 Tensile properties at room temperature of TPU33-CNF and TPU33-CNFOX composites: (a) stress at break, (b) strain at break, and (c) Young’s modulus.

143 PCL-MDI-BDO-based TPUs with different PCL-diol molecular weights and soft segment contents. They observed an increase of the modulus at room temperature with the increase of the semi-crystalline soft segment content. Conversely, they observed a reduction of the modulus at 65 °C with the increase of the soft segment content. This result indicated that the crystalline soft segments predominately affected the value of the

Young’s modulus of the TPU materials.

Typical stress-strain diagrams obtained at 60 °C for TPU33-CNF and TPU33-

CNFOX materials prepared in the chaotic mixer are shown in Figure 6-23. The values of stress at break and Young’s modulus at 60 °C and 100 % elongation for the composites are shown in Figure 6-24. As PCL crystals are melted, tensile properties rely only on the hard segment domains and the possible reinforcement effect of CNF or CNFOX. From

Figure 6-23, it is observed that neat TPU33 could not be stretched to 100 % at 60 °C. On the contrary, CNF and CNFOX composites were easily stretched and their stress values increased with carbon nanofibers content (Figure 6-24(a)). Neat TPU33 materials prepared in the chaotic mixer and Brabender Plasticorder showed similar Young’s modulus values at 60 °C, as shown in Figure 6-24(b). This is in accordance with their equal hard segment composition. However, composite materials prepared in Brabender

Plasticorder showed a reduction of maximum stress and Young’s modulus with the increase of CNF content (Figure 6-24). As it was learned in Chapter V in the context of composites of PMMA, carbon nanofibers processed in a Brabender Plasticorder suffered more damage and therefore higher reduction of the fiber aspect ratio than in the chaotic mixer. TPU33 composites with CNFOX showed higher values of Young’s modulus than

144 CNF composites which may represent the effect of the chemical treatments on these fibers.

8 7 (a) 6 5

4 0 wt. % 3 1 wt. % Stress (MPa) Stress 3 wt. % 2 5 wt. % 1 7 wt. % 0 0 0.2 0.4 0.6 0.8 1 Strain 8 (b) 7 6 5 4 0 wt. % 3 1 wt. % Stress (MPa) 2 3 wt. % 5 wt. % 1 7 wt. % 0 00.20.40.60.81 Strain

Figure 6-23 Typical stress-strain curves at 60 °C for (a) TPU33-CNF and (b) TPU33- CNFOX composites prepared in a chaotic mixer.

The data provided in Figures 6-22(c) and 6-24(b) indicate that the modulus of

TPU33 materials was largely determined by the presence of PCL crystals. The formation of these crystals was impeded at room temperature by the presence of carbon nanofibers in the system, thus causing a large reduction of the values of Young’s modulus, except in

145 the case of TPU33-CNFOX. However, once the crystals were melted, e.g., at 60 °C, nanofibers with higher aspect ratio caused reinforcement of the polymer matrix and led to increase of tensile strength, tensile modulus, and storage modulus of composites prepared in the chaotic mixer.

9 TPU33-CNF Chaotic (a) 8 TPU33-CNF Brabender 7 TPU33-CNFOX Chaotic 6 5 4 3 2 Maximum Stress (MPa) 1 0 02468 Carbon nanofiber content (wt. %)

60.0 TPU33-CNF Chaotic (b) TPU33-CNF Brabender 50.0 TPU33-CNFOX Chaotic

40.0

30.0

20.0 Young's modulus(MPa) 10.0

0.0 02468 Carbon nanofiber content (wt. %)

Figure 6-24 Tensile properties at 60 °C for TPU33-CNF composites prepared in chaotic mixer and Brabender Plasticorder, and TPU33-CNFOX in chaotic mixer. (a) Maximum stress and (b) Young’s modulus.

146 6.2.6. Hydrogen bonding

Figure 6-25 shows a typical FTIR spectrum of TPU33 materials with the main peaks assigned. Unlike in the case of TPU23 materials (Section 6.1.7), the degree of phase separation, DPS, cannot be estimated based on the hydrogen bonding index, R, obtained from the FTIR analysis of the materials. We need to remember that the presence of carbonyl groups in the soft segments in TPU33 entitles us only to utilize R (AHCO/ACO) values as a measure of the status of total hydrogen-bonded carbonyl groups in TPU. In addition, the status of hydrogen-bonded NH groups was determined from the values of

ANH/ACH ratio. This quantity was used to observe the influence of carbon nanofiber composition, mixing method, nanofiber surface chemistry and temperature on the behavior of NH groups, and consequently on the hard domains.

ν(C=C) (N-H) ν benzene ring free N-H νs(CH2) ν(N-H) bonded N-H

νa(CH2) ν (C=O) H bonded

Transmitance (%) δ(N-H) + ν(C-N)

ν (C=O) ν = stretching (a = asymmetric; s = symmetric) free δ = bending

3800 3300 2800 2300 1800 1300 800 Wavenumber (cm-1)

Figure 6-25 Typical FTIR spectra of TPU33-CNF5 prepared in Brabender Plasticorder with principal peak assignments. 147 Amine groups in polyurethanes might show up to four overlapping symmetric peaks at 3440, 3320, 3270, and 3180 cm-1. These peaks have been assigned to free NH, hydrogen-bonded NH, trans, and cis NH conformations, respectively.173 The proton on the amine group in the hard segments can interact with other functional groups, namely carbonyl groups either in the hard or soft segments, and even with π-electrons of the

174 aromatic rings. Table 6-4 show ANH/ACH ratios for TPU33 according carbon nanofiber content and type, and according to the processing method used to produce the materials.

It can be seen from Table 6-4 that hydrogen-bonded NH groups in TPU33 slightly decreased with CNF content in materials prepared by mixing in Brabender Plasticorder.

This quantity did not change appreciably for TPU33 composites prepared in the chaotic mixer, and showed fluctuation in the case of TPU33-CNFOX materials.

Table 6-4 Ratio of the area under the peak of hydrogen-bonded NH to aliphatic CH, ANH/ACH of TPU33-CNF and TPU33-CNFOX composites (±0.05).

Composition Chaotic mixer Brabender mixer Chaotic mixer

(wt. %) (TPU33-CNF) (TPU33-CNF) (TPU33-CNFOX)

0 0.40 0.45 0.40

1 0.40 0.41 0.34

3 0.42 0.35 0.40

5 0.38 0.35 0.35

7 0.41 0.38 N/D

148 This result is opposite of what was found for TPU23 materials in which hard domains were more affected by the presence of CNF, which was more prominent in the case of CNFOX and CNFOL due to their higher interaction with TPU23. The insensitivity of values of ANH/ACH in TPU33 can be attributed to a higher content of hard segments (33 wt. %) than in TPU23 (23 wt. %).

The effect of CNF content and mixing method on the AHCO/ACO values at room temperature of TPU33 was investigated (Figure 6-26(a)). Similarly, the effect of oxidized nanofibers on AHCO/ACO at room temperature is compared in Figure 6-26(b).

0.8 (a) Brabender Chaotic 0.6 CO CO

/A 0.4 HCO A

0.2

0 0 wt. % 1 wt. % 3 wt. % 5 wt. % 7 wt. % CNF content 1 (b) CNF CNFOX 0.8

CO 0.6 /A HCO

A 0.4

0.2

0 0 wt. % 1 wt. % 3 wt. % 5 wt. % Carbon nanofiber content

Figure 6-26 Values of AHCO/ACO in TPU33 composites according to (a) mixing method, and (b) fiber oxidation.

149 From Figure 6-26(a), it is observed that in general, the concentration of hydrogen- bonded carbonyl groups in TPU33 decreased with the addition of CNF. Such effect was more obvious in materials prepared in Brabender Plasticorder. On the other hand, the incorporation CNFOX to TPU33 system increased AHCO/ACO values. These extra hydrogen-bonded carbonyls may have come from the interaction of amine groups in the hard segments with carboxylic groups present on the surface of CNFOX, as it was seen in the case of TPU23 with CNFOX and CNFOL (Section 6.1.7).

The behavior of hydrogen-bonded -NH groups at temperatures above 25 °C was analyzed for pure TPU33 and its 7 wt. % composites. Change in the ANH/ACH ratio was determined for every temperature as shown in Table 6-5. It is noticed that the amount of hydrogen-bonded NH groups is reduced with an increase in temperature.

Table 6-5 Effect of the temperature on the ratio of the area under the peak of hydrogen- bonded NH to aliphatic CH, ANH/ACH in neat TPU33 and TPU33-CNF7 (±0.05).

Temperature TPU33 TPU33-CNF 7 TPU33-CNF 7

(°C) (Chaotic mixer) (Chaotic mixer) (Brabender mixer)

25 0.40 0.39 0.41

60 0.42 0.35 0.34

120 0.28 0.30 0.31

180 0.25 0.28 0.22

Such reduction becomes more prominent at temperatures of 120 °C and above, which can be related to the glass transition of the hard segments.173 It is important to note

150 that even at 180 °C there are still some hydrogen-bonded amine groups in neat TPU33 and composites. It was found previously that significant hydrogen bonding can survive at

173 even 200 °C in a TPU. The deviation of ANH/ACH values from room temperature in

Table 6-5 seems to be faster for composites than for the neat polymer. One reason for this behavior might be the ability of the carbon nanofibers to improve the thermal conductivity of the polymer, allowing for more uniform heating and disruption of hydrogen bonds in the composites at 60 °C. Such effect was also reflected in reduction of hydrogen-bonded carbonyl groups, HCO, as it can be observed in Figure 6-27. At 60 °C, practically the same amount of HCO groups still remained as in neat TPU33 (Figure 6-

27a), while in the 7 wt. % composite, hydrogen bonding started to diminish at that temperature (Figure 6-27b).

(a) (b) RT

120 °C Absorbance (a.u.)

RT (a.u.) Absorbance

60 °C

120 °C 180 °C 60 °C 180 °C

1820 1800 1780 1760 1740 1720 1700 1680 1660 1640 1820 1800 1780 1760 1740 1720 1700 1680 1660 1640 Wavenumber (cm-1) Wavenumber (cm-1)

Figure 6-27 FT-IR spectra of C=O stretching region at different temperatures for (a) pristine TPU33 and (b) TPU33-CNF 7 wt. % prepared in the chaotic mixer.

The behavior of AHCO/ACO at various temperatures was also analyzed for pure

TPU33 and its 7 wt. % composites. Change in AHCO/ACO with increasing temperature can be observed in Figure 6-28. From this figure, it is clear that dissociation of hydrogen bonding occurred with increasing the temperature. Pristine TPU33 experienced more

151 significant dissociation only by heating up to 120 °C. Pattanayak and Jana162 observed a decay in the IR absorption of hydrogen-bonded carbonyl peak above 110 °C in TPU and

TPU-nanoclay composites, which is similar to what we found here for TPU33. On the other hand, such levels of dissociation occurred prematurely at 60 °C in the composites with CNF, as it is observed in Figure 6-28.

0.7 RT 0.6

60 °C 0.5 RT

CO 0.4 RT /A 60 °C HCO 0.3 120 °C A 120 °C 60 °C 120 °C 180 °C 0.2 180 °C 180 °C 0.1

0 0 wt. % 7 wt.% chaotic 7 wt.% Brabender

Figure 6-28 Values of AHCO/ACO of neat TPU33 and its 7 wt. % composites at several temperatures.

6.2.7. Thermal transitions

DSC technique was used to investigate thermal transitions on heating and cooling of neat TPU33 and its composites with CNF and CNFOX prepared by two different mixing techniques.

Specimens of TPU33-CNF composites prepared in the chaotic mixer were subjected to a heating scan in the DSC apparatus in order to determine the most important

152 morphological features based on the thermal transitions. Figure 6-29 reveals some of the transitions that can be found in TPU literature166-170,175 and already found and discussed previously for TPU23 materials in Section 6.1.8. These transitions include soft segment glass transition temperature around -50 °C (Region I), hard segment glass transition around 125 °C (Region II), and endotherms for hard segments melting between 175 and

225 °C (Region III). In addition to those transitions, TPU33 shows an endotherm at about

50 °C corresponding to the melting of PCL crystals which form the soft segments. The effect of adding CNF in TPU33 is clearly seen in the PCL melting endotherm. Increasing

CNF content shifted PCL crystals melting temperature to lower values, and reduced the area under the curve which implies a reduction in the crystallinity.

Region I Region II Region III

Figure 6-29 Full DSC heating scan of TPU33-CNF composites prepared in the chaotic mixer.

On the other hand, hard domains seemed not to be affected by the presence of

CNF, as their melting endotherms can be appreciated even at 5 wt. % CNF content

(Region III). This is in agreement with the results presented in Section 6.2.6.1 concerning 153 the behavior of ANH/ACH values, and with what it has been observed in the case of using heat treated CNF.114

In order to fully understand the influence of carbon nanofibers on the soft segments morphology of TPU33, two consecutive heating scans were performed, as described in Chapter III, Section 3.3.6. Figures 6-30(a) and 6-30(b) show first and second scan respectively of TPU33-CNF materials synthesized in the chaotic mixer. Figures 6-

31(a) and 6-31(b) show similar data for materials prepared in the Brabender Plasticorder.

Neat TPU33 prepared in either mixer shows the highest melting temperature in the first

DSC scan as seen in Figures 6-30(a) and 6-31(a) respectively. In a second scan, different populations of spherulites in TPU33-CNF composites rearranged and produced a uniform size distribution after melting and annealing at 100 °C at the end of the first temperature scan (Figures 6-30(b) and 6-31(b)).

(a) (b)

Figure 6-30 (a) First and (b) second heating scan of specimens of TPU33-CNF composites mixed in a chaotic mixer.

154 (a) (b)

Figure 6-31 (a) First and (b) second heating scan of specimens of TPU33-CNF composites mixed in the Brabender mixer.

As it can be seen from Figures 6-30, there is a tendency of the melting temperature of crystalline soft segments to increase after the initial decrease at 1 wt. %

CNF for materials prepared in the chaotic mixer.

PLC crystallinity was evaluated as the ratio of the heat of fusion of each melting endotherm to the heat of fusion of a 100 % crystalline PCL141, and the results are shown in Figure 6-32. The presence of CNF reduced the PCL crystallinity for all composites prepared either in the chaotic mixer or Brabender Plasticorder. Although the overall crystallinity was higher for TPU33 materials made in the chaotic mixer, the composites experienced similar drop in the crystallinity of PCL soft segments with CNF content.

Lower crystallinity in TPU33 materials prepared in the Brabender Plasticorder might be a consequence of the higher degree of phase mixing between hard and soft segments.

Higher values of Tg previously obtained from E” and tan δ plots in DMA (Section 6.2.3) also pointed toward a similar conclusion.

155 30 30 (a) (b) 25 25

20 20 Chaotic 15 Chaotic 15

10 10 Crystallinity (%) Crystallinity Crystallinity (%) Crystallinity Brabender 5 5 Brabender 0 0 0246802468 CNF content (wt. %) CNF content (wt. %)

Figure 6-32 Percentage of crystallinity of TPU33-CNF prepared by two processing techniques, (a) first and (b) second DSC scan.

Accordingly, Tg values were also estimated from DSC thermograms, and the results are shown in Figure 6-33. Values from DSC followed similar trend as those found by DMA, that is, higher Tg values in materials prepared in Brabender Plasticorder. In this case, more interaction and connectivity between the hard and soft segments increased the soft segments Tg, and decreased the ability of the soft segment chains to arrange and crystallize properly.

-48

-50

-52 Brabender

-54 Temperature (°C) -56 Chaotic

-58 012345678 CNF content (wt. %)

Figure 6-33 Glass transition temperature, Tg, of TPU33 materials synthesized in two different mixers (from first DSC scan).

156 The effect of the surface chemistry in carbon nanofibers on PCL crystallinity was also assessed by running double heating scans on TPU33-CNFOX materials prepared in the chaotic mixer. Figures 6-34(a) and 6-34(b) show first and second DSC thermograms for these composites.

(a) (b)

Figure 6-34 (a) First and (b) second heating scan of specimens of TPU33-CNFOX composites mixed in a chaotic mixer.

First DSC scan in Figure 6-34(a) reveals only one crystalline melting endotherm for PCL soft segments in TPU33-CNFOX in contrast to what was observed in Figure 6-

30(a). Also from this figure it can be seen that the melting temperature slightly shifted to lower values for a TPU33 composite with 1 wt. % of CNFOX. There was a more significant reduction on crystalline PCL melting for composites having 3 wt. % CNFOX or more. A second DSC scan (Figure 6-34(b)) revealed a progressive but slight shift to lower values for this transition. This is in contrast to what was observed in a second DSC scan for TPU33-CNF composites (Figure 6-30(b)). As it can be clearly seen, there was a remarkable difference in terms of crystalline behavior of TPU33-CNFOX in comparison to TPU33-CNF composites.

157 30 30 (a) (b) 25 CNFOX 25 20 CNFOX

15 20 CNF 10 Crystallinity (%) Crystallinity Crystallinity (%) Crystallinity 15 CNF 5

0 10 0246802468 CNF content (wt. %) CNF content (wt. %)

Figure 6-35 Percentage of crystallinity of TPU33-CNFOX prepared by chaotic mixing, (a) first and (b) second DSC scan.

A nucleation effect can be observed in the case of 1 wt. % CNFOX content as seen in Figure 6-35(b). Nucleation by the presence of CNF114, SWNT124 and nanoclay176 in TPU have been reported previously in the literature. On the other hand, the nucleating effect of acid-treated MWNT in a PCL matrix has been recently investigated.143 These authors found that 0.25 wt. % MWNT induced a heterogeneous nucleation by reducing the activation energy of crystallization. However, higher amounts of MWNT increased the activation energy due to a restriction on the mobility of the polymer chains during crystallization.

It was seen above that CNF affected the crystallinity behavior of PCL in TPU33.

A reduction in the crystal melting temperature implied less perfect PCL crystallites and as a consequence, smaller and defective spherulites. In order to complete that scenario, crystallization temperature of PCL-soft segments in TPU33-CNF composites was determined by a non-isothermal crystallization procedure. Figure 6-36 shows heating- cooling cycles for TPU-CNF 5 wt. % composites prepared in Brabender Plasticorder

(Figure 6-36(a)) and in the chaotic mixer (Figure 6-36(b)).

158 5 5 (a) (b) Cooling Cooling

0 0 Heat flow (mW) Heat flow (mW) Heating Heating

-5 -5 -10 0 10 20 30 40 50 60 70 80 90 100 110 -10 0 10 20 30 40 50 60 70 80 90 100 110 Temperature (°C) Temperature (°C)

Figure 6-36 Heating and cooling DSC scan of TPU-CNF 5 wt. % prepared in (a) Brabender and (b) chaotic mixer.

Crystallization of PCL segments was more constrained in TPU33 composites fabricated in the conventional mixer, as a lower temperature was needed for soft segment crystallization, as it can be observed in Figure 6-36. Similar trend was observed for neat

TPU33 as well as the rest of the composites according to Figure 6-37.

36

CHAOTIC 32

28 BRABENDER

Temperature (°C) Temperature 24

20 012345678 CNF content (wt. %)

Figure 6-37 Crystallization temperature of TPU33-CNF composites.

159 6.3. Summary

In this Chapter we applied the concept of chaotic mixing to synthesize TPU- carbon nanofiber composites, and compared the results with materials prepared in a commercial mixer. Good dispersion of the nanofibers in TPU was possible by chaotic mixing, and fiber alignment was possible due to the 2D character of the chaotic flow. In general, the presence of carbon nanofibers in either TPU23 or TPU33 provided thermal stability to the matrix. The presence of oxygen-containing functional groups on the surface of the carbon nanofibers improved their dispersion in the TPU matrix. Although

TPU23-CNF materials did not show electrical conductivity in the evaluated range of composition, it is believed that TPU33-CNF and TPU33-CNFOX showed high electrical conductivity because of the different morphologies. PCL crystals in TPU33 may have played an important role in producing electrical composites by pushing away and agglomerating the carbon filler to limited spaces. This hypothesis is based on the following two aspects. Although is true that by using 90 minutes of mixing to produce

TPU23 composites might have produced electrically insulator materials, we also wanted to know how much of this time corresponded to effective mixing time. By recalling

Figure 6-1 which describes the torque increase with time for TPU23 composites, it was observed that the torque increased dramatically in less than 5 minutes. At that time, reaction was complete and continued motion of the rotors did not produce additional mixing. A second aspect to consider is the results obtained for PMMA-CNFOX in

Chapter V. While PMMA-CNFOX materials showed lower electrical conductivity than their CNF counterparts, TPU33-CNFOX had similar electrical behavior to TPU33-CNF.

160 In the first case, surface chemistry played a key role in determining the electrical conductivity properties, while in the second case, a more complex morphology played that key role. This subtle difference between the morphologies of TPU23 and TPU33 could not be detected by processing in the Brabender Plasticorder mixer due to its intensive mixing character. In this regard, chaotic mixing offers a powerful tool to investigate very complex morphologies such as those found in TPU33 composites.

In general, it was found that CNFs interacted with both soft and hard segments in

TPU. Due to the lack of a crystalline soft segment in TPU23 and to its lower hard segments content, the presence of CNF in the system generated disorder in the hard domains. This disorder was critical when using chemically treated CNF due to the higher interaction of these fibers with the polymer matrix. On the other hand, the presence of

CNF in TPU33 materials was more detrimental to the crystallinity of the soft phase.

Crystallizable soft segments in TPU33 controlled the mechanical properties at room temperature. It is evident that CNF reduced the ability of the PCL phase to crystallize. This negative effect on crystallization was mild in the case of TPU33-

CNFOX materials prepared by chaotic mixing, intermediate for TPU33-CNF composites prepared also in the chaotic mixer, and severe in the case of TPU33-CNF produced in the

Brabender Plasticorder. More phase mixing occurred in TPU33 materials prepared in the

Brabender mixer which affected PCL crystallinity. Note that high crystallinity is a requirement for TPU materials utilized in shape memory applications. Next Chapter presents results on shape memory behavior of TPU-CNF composites prepared in the chaotic mixer to capitalize on the high crystallinity of such materials.

161 CHAPTER VII

SHAPE MEMORY PROPERTIES IN TPU COMPOSITES

In this Chapter the shape memory properties of TPU33-CNF and TPU33-CNFOX composites are discussed. The recovery stress, shape fixity, and shape recovery values were evaluated.

7.1. Thermally-induced shape memory effect

Specimens of TPU33-CNF and TPU33-CNFOX were subjected to 50 % strain at

60 °C in the tensile tester and then cooled down in order to determine the percentage of shape retention or fixity. Figure 7-1 shows a plot showing the values of shape fixity for

CNF and CNFOX composites. It is seen that fixity values for composites reduced compared to unfilled TPU. However, this decrease was more severe in the case of

TPU33-CNF composites. Retention of the strain in this type of materials is directly related to the percentage of crystallinity in the soft segments. Composites prepared in the

162 Brabender Plasticorder mixer showed lower crystallinity than those prepared in the chaotic mixer, as was discussed in Chapter VI.

100 TPU33-CNF 80 TPU33-CNFOX

60

Fixity (%) 40

20

0 02468 Carbon Nanofiber content (wt. %)

Figure 7-1 Percent of shape retention or fixity of TPU33-CNF and TPU33-CNFOX composites after 50 % strain at 60 °C.

On the other hand, the values of fixity decreased gradually in the case of TPU33-

CNFOX. This behavior also corresponded to how crystallinity reduced with carbon nanofibers for this type of materials as observed in Chapter VI. Prior studies on TPU systems have established relationships of the magnitude of fixity with the crystallinity in soft segment phases.73

Figure 7-2 shows how stress recovery occurred in TPU33-CNF and TPU33-

CNFOX composite specimens that were initially 100 % elongated at 60 °C. It can be appreciated from this figure that TPU33-CNF 5 and TPU33-CNFOX 1 composites have the highest recovery stresses for each set of materials. TPU33-CNF 5 shows a maximum recovery stress of 3 MPa while TPU33-CNFOX 1 reaches almost 4 MPa in stress recovery. In general, composites with 7 wt. % of carbon nanofibers showed the lowest stress recovery values, mainly due to their low crystallinity and fixity values. 163 3.5 4.5 (a) (b) 3 5 wt. % 4 1 wt. % 3.5 2.5 5 wt. % 3 wt. % 3 2 7 wt. % 2.5 3 wt. % 1.5 1 wt. % 2 Stress (MPa) 0 wt. % (MPa) Stress 1.5 1 7 wt. % 1 0.5 0.5 0 0 20 30 40 50 60 70 80 90 100 20 30 40 50 60 70 80 90 100 Temperature (°C) Temperature (°C)

Figure 7-2 Stress recovery in (a) TPU33-CNF and (b) TPU33-CNFOX composites prepared in a chaotic mixer.

A value of about 2.5 MPa in stress recovery has been previously reported for a neat TPU system consisting of PCL4000/MDI/BDO with 30 wt. % hard segments.65

Mixing commercially available Morthane® with 5 wt. % of heat-treated CNF, Koerner et al.177 obtained a stress recovery value of 1.5 MPa during heating a constrained specimen at 55 °C. Recently, Mondal and Hu142 obtained a maximum of 1 MPa in stress recovery of a PTMG-based TPU having 1.5 wt. % of functionalized MWNT. The recovery ratio was also determined for these materials and it is shown in Figure 7-3. In this case, all

TPU33-CNFOX composites, and TPU33-CNF 1 had 100 % recovery of the original shape. However, we need to consider that the level of fixity in TPU33-CNF 1 was much lower than those in TPU33-CNFOX. TPU33-CNFOX composites showed higher stress recovery, fixity and recovery ratio than their CNF counterparts, due to a combination of higher PCL crystallinity and the reinforcement effect of CNFOX at temperatures above room temperature.

In order to induce recovery electrically of a deformed crystalline material like

TPU33, it is necessary to determine the voltage that can generate enough heat to melt the

164 PCL crystals. For that purpose, the composite should have a low surface resistivity to allow the electrons to flow when a specimen is clamped between two electrodes.

100

90

Recovery ratio (%) TPU33-CNF TPU33-CNFOX 80 02468 Carbon nanofiber content (wt. %)

Figure 7-3 Recovery ratio of TPU33-CNF and TPU33-CNFOX composites.

7.2. Shape memory effect induced by Joule heating

By reflecting on Figure 6-21 in Chapter VI once more, we find that there are only two TPU33 composites with low surface resistivity – 5 and 7 wt. % CNF prepared in the chaotic mixer. Therefore, those two compositions were chosen to investigate the electro- active shape recovery effect. Figure 7-4 presents how the sample temperature changed during application of voltage on these two composites. Joule heating was more important for the composite having 5 wt. % CNF at 300 V compared to 7 wt. % CNF in TPU33. It is seen that at 300 V, surface temperature in TPU33-CNF 5 reached above 50 °C. The heat generated, Q in time t, depends on the resistance, R of the material and the electrical current, I, according to Equation 25.

Q = RI 2t (25)

165 60 40 (a) (b) 150 V 55 300 V 200 V 50 35 100 V 45 200 V 300 V 40 30 150 V 35

Temperature (°C) 30 100 V Temperature (°C) 25 25 20 20 0 306090120 0 306090120 Time (sec) Time (sec)

Figure 7-4 Increase of temperature with applied voltage for (a) TPU33-CNF 5 and (b) TPU33-CNF 7 composites.

Although the Joule heating requires higher electrical resistivity (R), it depends on the square of the electrical current. It was seen in Figure 6-21 that both 5 and 7 wt. %

TPU33-CNF composites had similar surface resistivity. However, TPU33-CNF 5 developed the heat necessary to soften and melt PCL crystals in the polyurethane system.

On the other hand, TPU33-CNFOX composites were tested and none of them produced

Joule heating at any of the applied voltages included in Figure 7-4. These materials showed higher surface resistivity (Figure 6-21c) which did not contribute to the flow of the electrons on the surface. Cho et al.144,145 obtained a temperature increase of up to

35 °C in eight seconds when they applied 60 V to a TPU based on PCL3000/MDI/BDO mixed with 5 wt. % of acid-treated MWNT. Shape recovery of TPU33-CNF5 composite by Joule heating was observed by recording images using a video camera. From the video clips, photo frames were taken at time intervals of 10 seconds, and shown in Figure 7-5.

At the initial position (0 sec) a deformed specimen was placed between two alligator clamps. It can be appreciated from Figure 7-5 that after 20 seconds of applying the electrical voltage, the deformed bar started to move and reduced its angle, θ, with respect

166 to an imaginary vertical line. Such shift or motion continued after 40 seconds, and then remained steady after 60 seconds, as it can be seen from Figure 7-6.

θ θ θ θ θ

0 sec 10 sec 20 sec 40 sec 60 sec

Figure 7-5 Shape recovery of TPU33-CNF 5 triggered by Joule heating.

41 39 37 35 33 (degrees)

θ 31 29 27 25 0204060 Time (sec)

Figure 7-6 Change with time of the deflection angle, θ, in a specimen bar of TPU-CNF 5. Straight lines are used to guide the eye.

167 7.3. Summary

In this Chapter, we explored the shape memory properties of polyurethane nanocomposite materials, whereby their deformed shapes are recovered after heating. For this type of materials, triggering of the shape recovery effect can be achieved by thermal and electrical means. It is clear that PCL crystallinity played a very important role in determining the percentage of fixity, recovery, and stress at recovery during the thermal- actuated procedure. As we learned in Chapter VI, using CNFOX instead of CNF prevented a dramatic drop in PCL crystallinity which resulted in improved shape memory properties compared with neat TPU33 and TPU33-CNF composites. Therefore, it is possible to produce materials by chaotic mixing with higher crystallinity than those produced by conventional internal mixers. This factor turns out to be beneficial for shape memory effects.

On the other hand, TPU33-CNF composites are recommended for shape memory recovery induced by Joule heating due to their lower surface resistivity. In order to maximize this effect, it would be advisable to utilize heat-treated CNF or MWNT which have higher electrical conductivity. In this way, lower amount of carbon filler will be needed and this will reduce the possibility of negatively affecting PCL crystallinity.

168 CHAPTER VIII

OVERALL SUMMARY

The present document covered a research focused on the utilization of a novel polymer mixing technique such as chaotic mixing. The main objective of this research was to explore the concept of chaotic mixing in an attempt to obtain polymer-carbon nanofiber composites with superior properties that could complement or even surpass those properties of materials fabricated in a conventional batch mixer under almost identical conditions. It was learned that chaotic mixing proved to be a useful tool in dealing with highly entangled materials such as carbon nanofibers. The study shows that chaotic mixing filled a gap between low shear solution mixing and high shear intensive melt mixing, and combined the best out of those two approaches. In addition, manipulation of the surface chemistry of the carbon nanofibers worked in tandem with chaotic mixing in terms of facilitating the dispersion of such carbon materials.

Carbon nanofibers used in this project were highly entangled, and filled with amorphous carbon regions that affected the dispersion quality in polar polymers. By using oxidized carbon nanofibers or chemically attaching functional groups on their surface, it was possible to improve their dispersion in PMMA and TPU.

169 Dispersing entangled fibers such as CNF and trying to reduce fiber damage is a compromise in commercial intensive mixers. Chaotic mixing proved to be a technique that allows the user to regulate the level of dispersion in a wider range. In addition, good dispersion can be achieved with limited damage to the fibers length. Thus, chaotic mixers are also recommended in the case of CNT or nanoclay.

Chaotic mixing rendered composites with unique morphologies such as conductive networks of aligned nanofibers obtained for PMMA carbon nanofiber composites that could not be achieved with the commercial internal mixer. Such morphologies resulted in composites with high and directional electrical conductivity, or composites with good fiber dispersion. PMMA composites prepared in the chaotic mixer had much lower percolation composition than those prepared in the Brabender

Plasticorder. It was shown that in chaotic mixing, conductive networks were established due to the contact between dispersed and oriented fibers and their agglomerates. In addition, dispersed fibers were much less damaged due to low shear conditions. The study also showed that the electrical conductivity of materials prepared in the chaotic mixer was sensitive to the time of mixing, especially at nanofiber loadings around percolation thresholds. Shorter mixing times produced materials with high electrical conductivity although their thermo-mechanical properties were low. On the other hand, oxidized carbon nanofibers were more easily dispersed and resulted in PMMA composites with improved thermal and thermo-mechanical properties but with lower electrical conductivity. Such behavior was due to improved dispersion of the fibers which delayed formation of conductive networks until at higher volume fraction than those of untreated nanofibers.

170 In cases where polymerization and mixing takes places in the same processing device, chaotic mixing proved to be superior in producing TPU materials with excellent properties. It was seen that chaotic mixing is an excellent device to carry out TPU reactions, and mixing in situ with materials nanofillers such as CNF. TPU composites obtained by chaotic mixing showed good fiber dispersion, thermal stability, and high crystallinity. Longer aspect ratio fibers and their level of surface oxidation, coupled with chaotic mixing, produced conductive TPU composites with improved shape memory properties. The presence of oxygen-containing functional groups on the surface of the carbon nanofibers improved their dispersion in TPU as was the case with PMMA. In general, it was found that CNFs interacted with both soft and hard segments in TPU. The extent of TPU phase mixing was higher in materials prepared in the Brabender

Plasticorder. Carbon nanofibers improved the tensile stress and modulus of TPU at temperatures above the soft segment melting temperature. This allowed TPU composites to be stretched by 100 % at 60 °C, while neat TPU could not be stretched under the same conditions.

171 REFERENCES

(1) Nielsen, L., Landel, R. Mechanical properties of polymers and composites, 2nd ed.; Marcel Dekker: New York, 1994

(2) Rao, C. N. R., Cheetham, A. K. "Materials science at the nanoscale" In: Nanomaterials handbook, Gogotsi, Y., Ed.; CRC/Taylor & Francis: Boca Raton, 2006

(3) Usuki, A., Kojima, Y., Kawasumi, M., Okada, A., Fukushima, Y., Kurauchi, T., Kamigaito, O." Synthesis of Nylon 6-Clay Hybrid" Journal of Materials Research 1993, 8, 1179-1184.

(4) Kojima, Y., Usuki, A., Kawasumi, M., Okada, A., Fukushima, Y., Kurauchi, T., Kamigaito, O." Mechanical-Properties of Nylon 6-Clay Hybrid" Journal of Materials Research 1993, 8, 1185-1189.

(5) Kojima, Y., Usuki, A., Kawasumi, M., Okada, A., Kurauchi, T., Kamigaito, O." Sorption of Water in Nylon-6 Clay Hybrid" Journal of Applied Polymer Science 1993, 49, 1259-1264.

(6) Kroto, H. W., Heath, J. R., Obrien, S. C., Curl, R. F., Smalley, R. E." C-60 - Buckminsterfullerene" Nature 1985, 318, 162-163.

(7) Iijima, S." Helical Microtubules of Graphitic Carbon" Nature 1991, 354, 56-58.

(8) Lake, M. L., Ting, J.-M. "Vapor grown fiber composites" In: Carbon materials for advanced technologies, Burchell, T. D., Ed.; Pergamon: Netherlands, 1999

(9) Aref, H." Stirring by Chaotic Advection" Journal of Fluid Mechanics 1984, 143, 1- 21.

(10) Ottino, J. M. The kinematics of mixing : stretching, chaos, and transport; Cambridge University Press: Cambridge ; New York, 1989

(11) Ottino, J. M., Muzzio, F. J., Tjahjadi, M., Franjione, J. G., Jana, S. C., Kusch, H. A." Chaos, Symmetry, and Self-Similarity - Exploiting Order and Disorder in Mixing Processes" Science 1992, 257, 754-760.

172 (12) Leong, C. W., Ottino, J. M." Experiments on Mixing Due to Chaotic Advection in a Cavity" Journal of Fluid Mechanics 1989, 209, 463-&.

(13) Jana, S. C., Metcalfe, G., Ottino, J. M." Experimental and Computational Studies of Mixing in Complex Stokes Flows - the Vortex Mixing Flow and Multicellular Cavity Flows" Journal of Fluid Mechanics 1994, 269, 199-246.

(14) Raynal, F., Gence, J. N." Energy saving in chaotic laminar mixing" International Journal of Heat and Mass Transfer 1997, 40, 3267-3273.

(15) Rauwendaal, C. "Mixing in single screw extruders" In: Mixing and compounding of polymers : theory and practice, Agassant, J. F., Ed.; Hanser: OH, 1994

(16) Metcalfe, G., Ottino, J. M." Autocatalytic Processes in Mixing Flows" Physical Review Letters 1994, 72, 2875-2878.

(17) Danescu, R. I., Zumbrunnen, D. A." Creation of conducting networks among particles in polymer melts by chaotic mixing" Journal of Thermoplastic Composite Materials 1998, 11, 299-320.

(18) Jana, S. C., Scott, E., Sundararaj, U. "Single extruder screw for efficient blending of miscible and immiscible polymeric materials". US Patent No. 6132076, 1998

(19) Kim, S. J., Kwon, T. H." Enhancement of mixing performance of single-screw extrusion processes via chaotic flows .1. Basic concepts and experimental study" Advances in Polymer Technology 1996, 15, 41-54.

(20) Chougule, V. A., Zumbrunnen, D. A." In situ assembly using a continuous chaotic advection blending process of electrically conducting networks in carbon black- thermoplastic extrusions" Chemical Engineering Science 2005, 60, 2459-2467.

(21) Danescu, R. I., Zumbrunnen, D. A." Production of electrically conducting plastic composites by three-dimensional chaotic mixing of melts and powder additives" Journal of Vinyl & Additive Technology 2000, 6, 26-33.

(22) Zumbrunnen, D. A., Chhibber, C." Morphology development in polymer blends produced by chaotic mixing at various compositions" Polymer 2002, 43, 3267-3277.

(23) Jana, S. C., San, M." Effects of viscosity ratio and composition on development of morphology in chaotic mixing of polymers" Polymer 2004, 45, 1665-1678.

(24) Sau, M., Jana, S. C." Effect of waveforms on morphology development in chaotic mixing of polymers" Aiche Journal 2004, 50, 2346-2358.

173 (25) Sau, M., Jana, S. C." A study on the effects of chaotic mixer design and operating conditions on morphology development in immiscible polymer systems (vol 44, pg 407, 2004)" Polymer Engineering and Science 2004, 44, 1403-1403.

(26) Dharaiya, D. P., Jana, S. C." Nanoclay-induced morphology development in chaotic mixing of immiscible polymers" Journal of Polymer Science Part B-Polymer Physics 2005, 43, 3638-3651.

(27) Dharaiya, D. P., Jana, S. C., Lyuksyutov, S. F." Production of electrically conductive networks in immiscible polymer blends by chaotic mixing" Polymer Engineering and Science 2006, 46, 19-28.

(28) Zumbrunnen, D. A., Inamdar, S., Kwon, O., Verma, P." Chaotic advection as a means to develop nanoscale structures in viscous melts" Nano Letters 2002, 2, 1143-1148.

(29) Dresselhaus, M. S., Dresselhaus, G., Eklund, P. C., Saito, R., Endo, M. "Introduction to Carbon Materials" In: Carbon nanotubes : preparation and properties, Ebbesen, T. W., Ed.; CRC Press: Boca Raton, Fla., 1997

(30) Mordkovich, V. Z." Carbon nanofibers: A new ultrahigh-strength material for chemical technology" Theoretical Foundations of Chemical Engineering 2003, 37, 429- 438.

(31) Endo, M., Saito, R., Dresselhaus, M. S., Dresselhaus, G. "From Carbon Fibers to Carbon Nanotubes" In: Carbon Nanotubes: Preparation and Properties, Ebbesen, T. W., Ed.; CRC Press: Boca Raton, FL, 1997

(32) Issi, J.-P. "Thermal and Electrical Properties of Carbons: Relationship to Structure" In: Design and Control of Structure of Advanced Carbon Materials for Enhanced Performance, Yardim, M. F., Ed.; Kluwer Academic Publishers: Netherlands, 2001

(33) Ehrburger, P., Vix-Guterl, C. "Surface Properties of Carbons for Advanced Carbon-based Composites" In: Design and Control of Structure of Advanced Carbon Materials for Enhanced Performance, Yardim, M. F., Ed.; Kluwer Academic Publishers, Netherlands: Netherlands, 2001

(34) Ros, T. G., van Dillen, A. J., Geus, J. W., Koningsberger, D. C." Surface oxidation of carbon nanofibres" Chemistry-a European Journal 2002, 8, 1151-1162.

(35) Haiber, S., Ai, X. T., Bubert, H., Heintze, M., Bruser, V., Brandl, W., Marginean, G." Analysis of functional groups on the surface of plasma-treated carbon nanofibers" Analytical and Bioanalytical Chemistry 2003, 375, 875-883.

174 (36) Bubert, H., Ai, X., Haiber, S., Heintze, M., Bruser, V., Pasch, E., Brandl, W., Marginean, G." Basic analytical investigation of plasma-chemically modified carbon fibers" Spectrochimica Acta Part B-Atomic Spectroscopy 2002, 57, 1601-1610.

(37) Yue, Z. R., Jiang, W., Wang, L., Gardner, S. D., Pittman, C. U." Surface characterization of electrochemically oxidized carbon fibers" Carbon 1999, 37, 1785- 1796.

(38) Boehm, H. P." Surface oxides on carbon and their analysis: a critical assessment" Carbon 2002, 40, 145-149.

(39) Peebles, L. H. Carbon fibers : formation, structure, and properties; CRC Press: Boca Raton, FL, 1995

(40) Biniak, S., Szymanski, G., Siedlewski, J., Swiatkowski, A." The characterization of activated carbons with oxygen and nitrogen surface groups" Carbon 1997, 35, 1799-1810.

(41) Moreno-Castilla, C., Lopez-Ramon, M. V., Carrasco-Marin, F." Changes in surface chemistry of activated carbons by wet oxidation" Carbon 2000, 38, 1995-2001.

(42) Darmstadt, H., Summchen, L., Ting, J. M., Roland, U., Kaliaguine, S., Roy, C." Effects of surface treatment on the bulk chemistry and structure of vapor grown carbon fibers" Carbon 1997, 35, 1581-1585.

(43) Wu, Z. H., Pittman, C. U., Gardner, S. D." Nitric-Acid Oxidation of Carbon-Fibers and the Effects of Subsequent Treatment in Refluxing Aqueous Naoh" Carbon 1995, 33, 597-605.

(44) Serp, P., Figueiredo, J. L., Bertrand, P., Issi, J. P." Surface treatments of vapor- grown carbon fibers produced on a substrate" Carbon 1998, 36, 1791-1799.

(45) Figueiredo, J. L., Serp, P., Nysten, B., Issi, J. P." Surface treatments of vapor- grown carbon fibers produced on a substrate - Part II: Atomic force microscopy" Carbon 1999, 37, 1809-1816.

(46) Toebes, M. L., van Heeswijk, E. M. P., Bitter, J. H., van Dillen, A. J., de Jong, K. P." The influence of oxidation on the texture and the number of oxygen-containing surface groups of carbon nanofibers" Carbon 2004, 42, 307-315.

(47) Rakov, E. G. "Chemistry of carbon nanotubes" In: Nanomaterials handbook, Gogotsi, Y., Ed.; CRC/Taylor & Francis: Boca Raton, 2006

(48) Velasco-Santos, C., A.L, M.-H., Castaño, V. M. "Chemical functionalization on carbon nanotubes: Principles and applications" In: Trends in nanotechnology research, Dirote, E. V., Ed.; Nova Science Publishers: New York, 2004

175 (49) Liu, P." Modifications of carbon nanotubes with polymers" European Polymer Journal 2005, 41, 2693-2703.

(50) Jin, Z. X., Sun, X., Xu, G. Q., Goh, S. H., Ji, W." Nonlinear optical properties of some polymer/multi-walled carbon nanotube composites" Chemical Physics Letters 2000, 318, 505-510.

(51) Lin, Y., Rao, A. M., Sadanadan, B., Kenik, E. A., Sun, Y. P." Functionalizing multiple-walled carbon nanotubes with aminopolymers" Journal of Physical Chemistry B 2002, 106, 1294-1298.

(52) Gao, C., Jin, Y. Z., Kong, H., Whitby, R. L. D., Acquah, S. F. A., Chen, G. Y., Qian, H. H., Hartschuh, A., Silva, S. R. P., Henley, S., Fearon, P., Kroto, H. W., Walton, D. R. M." Polyurea-functionalized multiwalled carbon nanotubes: Synthesis, morphology, and Raman spectroscopy" Journal of Physical Chemistry B 2005, 109, 11925-11932.

(53) Wei, G., Saitoh, S., Saitoh, H., Fujiki, K., Yamauchi, T., Tsubokawa, N." Grafting of polymers onto vapor grown carbon fiber surface by ligand-exchange reaction of ferrocene moieties of polymer with polycondensed aromatic rings of the wall-surface" Polymer 2004, 45, 8723-8730.

(54) Woods, G., ICI Polyurethanes (Firm) The ICI Polyurethanes book, 2nd ed.; Published jointly by ICI Polyurethanes and Wiley: Chichester ; New York, 1990

(55) Ravve, A. Principles of polymer chemistry, 2nd ed.; Kluwer Academic/Plenum Publishers: New York, 2000

(56) Randall, D., Lee, S., Woods, G. The polyurethanes book; Distributed by J. Wiley: New York, 2002

(57) Lamba, N. M. K., Woodhouse, K. A., Cooper, S. L., Lelah, M. D. Polyurethanes in biomedical applications; CRC Press: Boca Raton, 1998

(58) Estes, G. M., Seymour, R. W., Cooper, S. L." Infrared studies of segmented polyurethane elastomers. II. Infrared dichroism" Macromolecules 1971, 25, 452-457.

(59) Bonart, R." X-ray investigations concerning the physical structure of crosslinking in segmented elastomers" Journal of Macromolecular Science and Physics 1968, B2, 115-138.

(60) Wilkes, C. E., Yusek, C. S." Investigation of domain structure in urethane elastomers by X-ray and thermal methods" Journal of Macromolecular Science and Physics 1973, B7, 157-175.

176 (61) Blackwell, J. L., Gardner, K. H." Structure of hard segments in polyurethane elastomers" Polymer 1979, 20, 13.

(62) Blackwell, J. L., Lee, C. D." Hard segment domain sizes in MDI/diol polyurethane elastomers" Journal of Polymer Science Polymer Physics Education 1983, 21, 2169.

(63) Lendlein, A., Kelch, S." Shape-memory polymers" Angewandte Chemie- International Edition 2002, 41, 2034-2057.

(64) Lee, B. S., Chun, B. C., Chung, Y. C., Sul, K. I., Cho, J. W." Structure and thermomechanical properties of polyurethane block copolymers with shape memory effect" Macromolecules 2001, 34, 6431-6437.

(65) Kim, B. K., Lee, S. Y., Xu, M." Polyurethanes having shape memory effects" Polymer 1996, 37, 5781-5793.

(66) Takahashi, T., Hayashi, N., Hayashi, S." Structure and properties of shape-memory polyurethane block copolymers" Journal of Applied Polymer Science 1996, 60, 1061- 1069.

(67) Li, F. K., Zhang, X., Hou, J. N., Xu, M., Lu, X. L., Ma, D. Z., Kim, B. K." Studies on thermally stimulated shape memory effect of segmented polyurethanes" Journal of Applied Polymer Science 1997, 64, 1511-1516.

(68) Liang, C., Rogers, C. A., Malafeew, E." Investigation of shape memory polymers and their hybrid composites (Reprinted from Proceedings of the Second Joint Japan/US Conference on Adaptive Structures, Nov, pg 789-802)" Journal of Intelligent Material Systems and Structures 1997, 8, 380-386.

(69) Lin, J. R., Chen, L. W." Study on shape-memory behavior of polyether-based polyurethanes. I. Influence of the hard-segment content" Journal of Applied Polymer Science 1998, 69, 1563-1574.

(70) Lin, J. R., Chen, L. W." Study on shape-memory behavior of polyether-based polyurethanes. II. Influence of soft-segment molecular weight" Journal of Applied Polymer Science 1998, 69, 1575-1586.

(71) Lin, J. R., Chen, L. W." Shape-memorized crosslinked ester-type polyurethane and its mechanical viscoelastic model" Journal of Applied Polymer Science 1999, 73, 1305- 1319.

(72) Ping, P., Wang, W. S., Chen, X. S., Jing, X. B." Poly(epsilon-caprolactone) polyurethane and its shape-memory property" Biomacromolecules 2005, 6, 587-592.

177 (73) Jeong, H. M., Lee, J. B., Lee, S. Y., Kim, B. K." Shape memory polyurethane containing mesogenic moiety" Journal of Materials Science 2000, 35, 279-283.

(74) Applied Sciences Incorporated,Cedarville, OH,http://www.apsci.com

(75) Cox, H. L." The elasticity and strength of paper and other fibrous materials" Brittish Journal of Applied Physics 1952, 3, 72-79.

(76) Baxter, W. J." The Strength of Metal Matrix Composites Reinforced with Randomly Oriented Discontinuous Fibers" Metallurgical Transactions a-Physical Metallurgy and Materials Science 1992, 23, 3045-3053.

(77) Clingerman, M. L., Weber, E. H., King, J. A., Schulz, K. H." Development of an additive equation for predicting the electrical conductivity of carbon-filled composites" Journal of Applied Polymer Science 2003, 88, 2280-2299.

(78) Dasch, C. J., Baxter, W. J., Tibbetts, G. G., "Thermoplastic composites using nanometer-size vapor grown carbon fibers" 21st Biennial Conference on Carbon 1993, 82-83

(79) Hudnut, S. W., Chung, D. D. L." Use of Submicron Diameter Carbon Filaments for Reinforcement between Continuous Carbon-Fiber Layers in a Polymer-Matrix Composite" Carbon 1995, 33, 1627-1631.

(80) Carneiro, O. S., Covas, J. A., Bernardo, C. A., Caldeira, G., Van Hattum, F. W. J., Ting, J. M., Alig, R. L., Lake, M. L." Production and assessment of polycarbonate composites reinforced with vapour-grown carbon fibres" Composites Science and Technology 1998, 58, 401-407.

(81) Higgins, B. A., Brittain, W. J." Polycarbonate carbon nanofiber composites" European Polymer Journal 2005, 41, 889-893.

(82) Tibbetts, G. G., McHugh, J. J." Mechanical properties of vapor-grown carbon fiber composites with thermoplastic matrices" Journal of Materials Research 1999, 14, 2871- 2880.

(83) Lozano, K., Barrera, E. V." Nanofiber-reinforced thermoplastic composites. I. Thermoanalytical and mechanical analyses" Journal of Applied Polymer Science 2001, 79, 125-133.

(84) Lozano, K., Bonilla-Rios, J., Barrera, E. V." A study on nanofiber-reinforced thermoplastic composites (II): Investigation of the mixing rheology and conduction properties" Journal of Applied Polymer Science 2001, 80, 1162-1172.

178 (85) Sandler, J., Shaffer, M. S. P., Prasse, T., Bauhofer, W., Schulte, K., Windle, A. H." Development of a dispersion process for carbon nanotubes in an epoxy matrix and the resulting electrical properties" Polymer 1999, 40, 5967-5971.

(86) Shofner, M. L., Lozano, K., Rodriguez-Macias, F. J., Barrera, E. V." Nanofiber- reinforced polymers prepared by fused deposition modeling" Journal of Applied Polymer Science 2003, 89, 3081-3090.

(87) Ma, H. M., Zeng, J. J., Realff, M. L., Kumar, S., Schiraldi, D. A." Processing, structure, and properties of fibers from polyester/carbon nanofiber composites" Composites Science and Technology 2003, 63, 1617-1628.

(88) Xu, J., Donohoe, J. P., Pittman, C. U." Preparation, electrical and mechanical properties of vapor grown carbon fiber (VGCF)/vinyl ester composites" Composites Part a-Applied Science and Manufacturing 2004, 35, 693-701.

(89) Xu, Y. J., Higgins, B., Brittain, W. J." Bottom-up synthesis of PS-CNF nanocomposites" Polymer 2005, 46, 799-810.

(90) Yang, S. Y., Lozano, K., Lomeli, A., Foltz, H. D., Jones, R." Electromagnetic interference shielding effectiveness of carbon nanofiber/LCP composites" Composites Part a-Applied Science and Manufacturing 2005, 36, 691-697.

(91) Wu, G. Z., Asai, S., Sumita, M." A self-assembled electric conductive network in short carbon fiber filled poly(methyl methacrylate) composites with selective adsorption of polyethylene" Macromolecules 1999, 32, 3534-3536.

(92) Zeng, J. J., Saltysiak, B., Johnson, W. S., Schiraldi, D. A., Kumar, S." Processing and properties of poly(methyl methacrylate)/carbon nano fiber composites" Composites Part B-Engineering 2004, 35, 173-178.

(93) Jimenez, G. A., Jana, S. C., "Preparation of poly(methyl methacrylate) and carbon nanofiber composites by chaotic mixing" 63rd Annual Technical Conference - Society of Plastics Engineers 2005, 1938-1942

(94) Jimenez, G. A., Jana, S. C." Electrically conductive polymer nanocomposites of polymethylmethacrylate and carbon nanofibers prepared by chaotic mixing" Composites Part A: Applied Science and Manufacturing 2006, Accepted for publication.

(95) Schueler, R., Petermann, J., Schulte, K., Wentzel, H. P." Agglomeration and electrical percolation behavior of carbon black dispersed in epoxy resin" Journal of Applied Polymer Science 1997, 63, 1741-1746.

(96) Jia, Z. J., Wang, Z. Y., Xu, C. L., Liang, J., Wei, B. Q., Wu, D. H., Zhu, S. W." Study on poly(methyl methacrylate)/carbon nanotube composites" Materials Science and

179 Engineering a-Structural Materials Properties Microstructure and Processing 1999, 271, 395-400.

(97) Sung, J. H., Kim, H. S., Jin, H. J., Choi, H. J., Chin, I. J." Nanofibrous membranes prepared by multiwalled carbon nanotube/poly(methyl methacrylate) composites" Macromolecules 2004, 37, 9899-9902.

(98) Benoit, J. M., Corraze, B., Lefrant, S., Blau, W. J., Bernier, P., Chauvet, O." Transport properties of PMMA-carbon nanotubes composites" Synthetic Metals 2001, 121, 1215-1216.

(99) de la Chapelle, M. L., Stephan, C., Nguyen, T. P., Lefrant, S., Journet, C., Bernier, P., Munoz, E., Benito, A., Maser, W. K., Martinez, M. T., de la Fuente, G. F., Guillard, T., Flamant, G., Alvarez, L., Laplaze, D." Raman characterization of singlewalled carbon nanotubes and PMMA-nanotubes composites" Synthetic Metals 1999, 103, 2510-2512.

(100) Stephan, C., Nguyen, T. P., de la Chapelle, M. L., Lefrant, S., Journet, C., Bernier, P." Characterization of singlewalled carbon nanotubes-PMMA composites" Synthetic Metals 2000, 108, 139-149.

(101) Cooper, C. A., Ravich, D., Lips, D., Mayer, J., Wagner, H. D." Distribution and alignment of carbon nanotubes and nanofibrils in a polymer matrix" Composites Science and Technology 2002, 62, 1105-1112.

(102) Haggenmueller, R., Gommans, H. H., Rinzler, A. G., Fischer, J. E., Winey, K. I." Aligned single-wall carbon nanotubes in composites by melt processing methods" Chemical Physics Letters 2000, 330, 219-225.

(103) Jin, Z., Pramoda, K. P., Xu, G., Goh, S. H." Dynamic mechanical behavior of melt-processed multi-walled carbon nanotube/poly(methyl methacrylate) composites" Chemical Physics Letters 2001, 337, 43-47.

(104) Gorga, R. E., Cohen, R. E." Toughness enhancements in poly(methyl methacrylate) by addition of oriented multiwall carbon nanotubes" Journal of Polymer Science Part B-Polymer Physics 2004, 42, 2690-2702.

(105) Du, F. M., Fischer, J. E., Winey, K. I." Coagulation method for preparing single- walled carbon nanotube/poly(methyl methacrylate) composites and their modulus, electrical conductivity, and thermal stability" Journal of Polymer Science Part B-Polymer Physics 2003, 41, 3333-3338.

(106) Du, F. M., Scogna, R. C., Zhou, W., Brand, S., Fischer, J. E., Winey, K. I." Nanotube networks in polymer nanocomposites: Rheology and electrical conductivity" Macromolecules 2004, 37, 9048-9055.

180 (107) Kashiwagi, T., Du, F. M., Winey, K. I., Groth, K. A., Shields, J. R., Bellayer, S. P., Kim, H., Douglas, J. F." Flammability properties of polymer nanocomposites with single- walled carbon nanotubes: effects of nanotube dispersion and concentration" Polymer 2005, 46, 471-481.

(108) Correa, R. A., Nunes, R. C. R., Lourenco, V. L." Investigation of the degradation of thermoplastic polyurethane reinforced with short fibres" Polymer Degradation and Stability 1996, 52, 245-251.

(109) Correa, R. A., Nunes, R. C. R., Franco, W. Z." Effect of injections on the orientation of short fibre composites. An optical microscopic analysis" Polymer Testing 1996, 15, 467-475.

(110) Correa, R. A., Nunes, R. C. R., Franco, W. Z." Short fiber reinforced thermoplastic polyurethane elastomer composites" Polymer Composites 1998, 19, 152- 155.

(111) Sanchez-Adsuar, M. S., Linares-Solano, A., Cazorla-Amoros, D., Ibarra-Rueda, L." Influence of the nature and the content of carbon fiber on properties of thermoplastic polyurethane-carbon fiber composites" Journal of Applied Polymer Science 2003, 90, 2676-2683.

(112) Segal, E., Tchoudakov, R., Narkis, M., Siegmann, A." Thermoplastic polyurethane-carbon black compounds: Structure, electrical conductivity and sensing of liquids" Polymer Engineering and Science 2002, 42, 2430-2439.

(113) Xiong, C. X., Zhou, Z. Y., Xu, W., Hu, H. R., Zhang, Y., Dong, L. J." Polyurethane/carbon black composites with high positive temperature coefficient and low critical transformation temperature" Carbon 2005, 43, 1788-1792.

(114) Koerner, H., Liu, W. D., Alexander, M., Mirau, P., Dowty, H., Vaia, R. A." Deformation-morphology correlations in electrically conductive carbon nanotube thermoplastic polyurethane nanocomposites" Polymer 2005, 46, 4405-4420.

(115) Kuan, H. C., Ma, C. C. M., Chang, W. P., Yuen, S. M., Wu, H. H., Lee, T. M." Synthesis, thermal, mechanical and rheological properties of multiwall carbon nano tube/waterborne polyurethane nanocomposite" Composites Science and Technology 2005, 65, 1703-1710.

(116) Kwon, J., Kim, H." Comparison of the properties of waterborne polyurethane/multiwalled carbon nanotube and acid-treated multiwalled carbon nanotube composites prepared by in situ polymerization" Journal of Polymer Science Part a- Polymer Chemistry 2005, 43, 3973-3985.

181 (117) Kwon, J. Y., Kim, H. D." Preparation and properties of acid-treated multiwalled carbon nanotube/waterborne polyurethane nanocomposites" Journal of Applied Polymer Science 2005, 96, 595-604.

(118) Xia, H. S., Song, M." Preparation and characterization of polyurethane-carbon nanotube composites" Soft Matter 2005, 1, 386-394.

(119) Jung, Y. C., Sahoo, N. G., Cho, J. W." Polymeric nanocomposites of polyurethane block copolymers and functionalized multi-walled carbon nanotubes as crosslinkers" Macromolecular Rapid Communications 2006, 27, 126-131.

(120) Mondal, S., Hu, J. L." Thermal degradation study of functionalized MWNT reinforced segmented polyurethane membrane" Journal of Elastomers and Plastics 2006, 38, 261-271.

(121) Xiong, J., Zheng, Z., Qin, X., Li, M., Li, H., Wang, X." The thermal and mechanical properties of a polyurethane/multi-walled carbon nanotube composite" Carbon 2006, 44, 2701-2707.

(122) Xu, M., Zhang, T., Gu, B., Wu, J. L., Chen, Q." Synthesis and properties of novel polyurethane-urea/multiwalled carbon nanotube composites" Macromolecules 2006, 39, 3540-3545.

(123) Chen, W., Tao, X. M." Self-organizing alignment of carbon nanotubes in thermoplastic polyurethane" Macromolecular Rapid Communications 2005, 26, 1763- 1767.

(124) Xia, H. S., Song, M." Preparation and characterisation of polyurethane grafted single-walled carbon nanotubes and derived polyurethane nanocomposites" Journal of Materials Chemistry 2006, 16, 1843-1851.

(125) Xiong, J. W., Zhou, D. S., Zheng, Z., Yang, X. H., Wang, X. L." Fabrication and distribution characteristics of polyurethane/single-walled carbon nanotube composite with anisotropic structure" Polymer 2006, 47, 1763-1766.

(126) Stephan, C., Nguyen, T. P., Lahr, B., Blau, W., Lefrant, S., Chauvet, O." Raman spectroscopy and conductivity measurements on polymer-multiwalled carbon nanotubes composites" Journal of Materials Research 2002, 17, 396-400.

(127) Petrovic, Z. S., Zavargo, Z., Flynn, J. H., Macknight, W. J." Thermal-Degradation of Segmented Polyurethanes" Journal of Applied Polymer Science 1994, 51, 1087-1095.

(128) Brandl, W., Marginean, G., Chirila, V., Warschewski, W." Production and characterisation of vapour grown carbon fiber/polypropylene composites" Carbon 2004, 42, 5-9.

182 (129) Tibbetts, G. G., Finegan, I. C., Kwag, C." Mechanical and electrical properties of vapor-grown carbon fiber thermoplastic composites" Molecular Crystals and Liquid Crystals 2002, 387, 353-357.

(130) Finegan, I. C., Tibbetts, G. G., Glasgow, D. G., Ting, J. M., Lake, M. L." Surface treatments for improving the mechanical properties of carbon nanofiber/thermoplastic composites" Journal of Materials Science 2003, 38, 3485-3490.

(131) Zhu, J., Kim, J. D., Peng, H. Q., Margrave, J. L., Khabashesku, V. N., Barrera, E. V." Improving the dispersion and integration of single-walled carbon nanotubes in epoxy composites through functionalization" Nano Letters 2003, 3, 1107-1113.

(132) Kim, Y. J., Shin, T. S., Choi, H. D., Kwon, J. H., Chung, Y. C., Yoon, H. G." Electrical conductivity of chemically modified multiwalled carbon nanotube/epoxy composites" Carbon 2005, 43, 23-30.

(133) Cortes, P., Lozano, K., Barrera, E. V., Bonilla-Rios, J." Effects of nanofiber treatments on the properties of vapor-grown carbon fiber reinforced polymer composites" Journal of Applied Polymer Science 2003, 89, 2527-2534.

(134) Brandl, W., Marginean, G." Functionalisation of the carbon nanofibres by plasma treatment" Thin Solid Films 2004, 447, 181-186.

(135) Seo, M. K., Park, S. J., Lee, S. K." Influence of atmospheric plasma on physicochemical properties of vapor-grown graphite nanofibers" Journal of Colloid and Interface Science 2005, 285, 306-313.

(136) Hwang, G. L., Shieh, Y. T., Hwang, K. C." Efficient load transfer to polymer- grafted multiwalled carbon nanotubes in polymer composites" Advanced Functional Materials 2004, 14, 487-491.

(137) Gorga, R. E., Lau, K. K. S., Gleason, K. K., Cohen, R. E." The importance of interfacial design at the carbon nanotube/polymer composite interface" Journal of Applied Polymer Science 2006, 102, 1413-1418.

(138) Peng, M., Li, D. S., Chen, Y., Zheng, Q." Electrostatic-assembly of carbon nanotubes (CNTs) and polymer particles in water: a facile approach to improve the dispersion of CNTs in thermoplastics" Macromolecular Rapid Communications 2006, 27, 859-864.

(139) Jimenez, G. A., Jana, S. C., "Polymer composites of modified carbon nanofibers prepared by chaotic mixing" 64th Annual Technical Conference - Society of Plastics Engineers 2006, 352-356

183 (140) Li, F. K., Hou, J. N., Zhu, W., Zhang, X., Xu, M., Luo, X. L., Ma, D. Z., Kim, B. K." Crystallinity and morphology of segmented polyurethanes with different soft- segment length" Journal of Applied Polymer Science 1996, 62, 631-638.

(141) Li, F. K., Qi, L. Y., Yang, J. P., Xu, M., Luo, X. L., Ma, D. Z." Polyurethane/conducting carbon black composites: Structure, electric conductivity, strain recovery behavior; and their relationships" Journal of Applied Polymer Science 2000, 75, 68-77.

(142) Mondal, S., Hu, J. L." Shape memory studies of functionalized MWNT-reinforced polyurethane copolymers" Iranian Polymer Journal 2006, 15, 135-142.

(143) Wu, T. M., Chen, E. C." Crystallization behavior of poly(epsilon- caprolactone)/multiwalled carbon nanotube composites" Journal of Polymer Science Part B-Polymer Physics 2006, 44, 598-606.

(144) Cho, J. W., Kim, J. W., Jung, Y. C., Goo, N. S." Electroactive shape-memory polyurethane composites incorporating carbon nanotubes" Macromolecular Rapid Communications 2005, 26, 412-416.

(145) Yoo, H. J., Jung, Y. C., Sahoo, N. G., Cho, J. W." Polyurethane-carbon nanotube nanocomposites prepared by in-situ polymerization with electroactive shape memory" Journal of Macromolecular Science Part B-Physics 2006, 45, 441-451.

(146) Brydson, J. A. Plastics materials, 2d ed.; Van Nostrand Reinhold: New York,, 1970

(147) Frados, J., Ed. Plastics engineering handbook of the Society of the Plastics Industry, inc, 4th ed.; Van Nostrand Reinhold: New York, 1976

(148) Perpex Datasheet" Lucite International 2002.

(149) Higgins, B. A., "Carbon nanofiber-polymer composites for electronic applications". University of Akron, Akron OH, PhD Thesis, 2006

(150) Hughes, T., Applied Sciences Inc., Cedarville OH, Personal communication

(151) Rivin, D." Use of lithium aluminum hydride in the study of surface chemistry of carbon black" Rubber Chemistry and Technology 1963, 36, 729-739.

(152) Jung, C. D., Jana, S. C., "Synthesis of thermoplastic polyurethanes by chaotic mixing" 62nd Annual Technical Conference - Society of Plastics Engineers 2004, 2814- 2818

184 (153) Jung, C. D., Jana, S. C., "Effect of chaotic mixing on catalyzed thermoplastic polyurethane polymerization" 63rd Annual Technical Conference - Society of Plastics Engineers 2005, 1800-1804

(154) Tien, Y. I., Wei, K. H." Hydrogen bonding and mechanical properties in segmented montmorillonite/polyurethane nanocomposites of different hard segment ratios" Polymer 2001, 42, 3213-3221.

(155) Troitskii, B. B., Troitskaya, L. S., Yakhnov, A. S., Lopatin, M. A., Novikova, M. A." Retardation of thermal degradation of PMMA and PVC by C-60" European Polymer Journal 1997, 33, 1587-1590.

(156) Laachachi, A., Cochez, M., Ferriol, M., Lopez-Cuesta, J. M., Leroy, E." Influence of TiO2 and Fe2O3 fillers on the thermal properties of poly(methyl methacrylate) (PMMA)" Materials Letters 2005, 59, 36-39.

(157) Zanetti, M., Lomakin, S., Camino, G." Polymer layered silicate nanocomposites" Macromolecular Materials and Engineering 2000, 279, 1-9.

(158) Kashiwagi, T., Grulke, E., Hilding, J., Harris, R., Awad, W., Douglas, J." Thermal degradation and flammability properties of poly(propylene)/carbon nanotube composites" Macromolecular Rapid Communications 2002, 23, 761-765.

(159) Potschke, P., Bhattacharyya, A. R., Janke, A." Melt mixing of polycarbonate with multiwalled carbon nanotubes: microscopic studies on the state of dispersion" European Polymer Journal 2004, 40, 137-148.

(160) Chen, W., Tao, X. M." Production and characterization of polymer nanocomposite with aligned single wall carbon nanotubes" Applied Surface Science 2006, 252, 3547- 3552.

(161) Pattanayak, A., Jana, S. C." Synthesis of thermoplastic polyurethane nanocomposites of reactive nanoclay by bulk polymerization methods" Polymer 2005, 46, 3275-3288.

(162) Pattanayak, A., Jana, S. C." Properties of bulk-polymerized thermoplastic polyurethane nanocomposites" Polymer 2005, 46, 3394-3406.

(163) Pattanayak, A., Jana, S. C." Thermoplastic polyurethane nanocomposites of reactive silicate clays: effects of soft segments on properties" Polymer 2005, 46, 5183- 5193.

(164) Gorrasi, G., Tortora, M., Vittoria, V." Synthesis and physical properties of layered silicates/polyurethane nanocomposites" Journal of Polymer Science Part B-Polymer Physics 2005, 43, 2454-2467.

185 (165) Seymour, R. W., Cooper, S. L." Thermal analysis of polyurethane block polymers" Macromolecules 1973, 6, 48-53.

(166) Koberstein, J. T., Russell, T. P." Simultaneous Saxs-Dsc Study of Multiple Endothermic Behavior in Polyether-Based Polyurethane Block Copolymers" Macromolecules 1986, 19, 714-720.

(167) Galambos, A. F., Russell, T. R., Koberstein, J. T." Multiple Melting in Polyurethane Block Copolymers" Abstracts of Papers of the American Chemical Society 1989, 198, 78-Pmse.

(168) Chen, T. K., Shieh, T. S., Chui, J. Y." Studies on the first DSC endotherm of polyurethane hard segment based on 4,4 '-diphenylmethane diisocyanate and 1,4- butanediol" Macromolecules 1998, 31, 1312-1320.

(169) Saiani, A., Daunch, W. A., Verbeke, H., Leenslag, J. W., Higgins, J. S." Origin of multiple melting endotherms in a high hard block content polyurethane. 1. Thermodynamic investigation" Macromolecules 2001, 34, 9059-9068.

(170) Saiani, A., Rochas, C., Eeckhaut, G., Daunch, W. A., Leenslag, J. W., Higgins, J. S." Origin of multiple melting endotherms in a high hard block content polyurethane. 2. Structural investigation" Macromolecules 2004, 37, 1411-1421.

(171) Krupa, I., Chodak, I." Physical properties of thermoplastic/graphite composites" European Polymer Journal 2001, 37, 2159-2168.

(172) Hong, C. M., Kim, J., Jana, S. C." Shear-induced migration of conductive fillers in injection molding" Polymer Engineering and Science 2004, 44, 2101-2109.

(173) Srichatrapimuk, V. W., Cooper, S. L." Infrared thermal analysis of polyurethane block polymers" Journal of Macromolecular Science B: Physics 1978, B15, 267-311.

(174) Wang, F. C., Feve, M., Lam, T. M., Pascault, J. P." Ftir Analysis of Hydrogen- Bonding in Amorphous Linear Aromatic Polyurethanes .1. Influence of Temperature" Journal of Polymer Science Part B-Polymer Physics 1994, 32, 1305-1313.

(175) Seymour, R. W., Allegrezza, A. E. J., Cooper, S. L." Segmental orientation studies of block polymers. I.Hydrogen bonded polyurethanes" Macromolecules 1973, 6, 896-902.

(176) Tortora, M., Gorrasi, G., Vittoria, V., Galli, G., Ritrovati, S., Chiellini, E." Structural characterization and transport properties of organically modified montmorillonite/polyurethane nanocomposites" Polymer 2002, 43, 6147-6157.

186 (177) Koerner, H., Price, G., Pearce, N. A., Alexander, M., Vaia, R. A." Remotely actuated polymer nanocomposites - stress-recovery of carbon-nanotube-filled thermoplastic elastomers" Nature Materials 2004, 3, 115-120.

187