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MULTISCALE FIBER REINFORCED COMPOSITES USING A CARBON NANOFIBER/EPOXY NANOPHASED MATRIX: PROCESSING, PROPERTIES, AND THERMOMECHANICAL BEHAVIOR

by KEITH JAMAHL GREEN

DERRICK R. DEAN, COMMITTEE CHAIR J. BARRY ANDREWS UDAY K. VAIDYA

A THESIS Submitted to the graduate faculty of the University of Alabama at Birmingham, in partial fulfillment of the requirements for the degree of Masters of Science BIRMINGHAM, ALABAMA 2007

MULTISCALE FIBER REINFORCED COMPOSITES USING A CARBON NANOFIBER/EPOXY NANOPHASED MATRIX: PROCESSING, PROPERTIES, AND THERMOMECHANICAL BEHAVIOR KEITH JAMAHL GREEN MATERIALS ENGINEERING ABSTRACT

Fiber-reinforced polymer composites (FRCs) have shown great promise as high strength structural materials due to their high stiffness to weight ratio and their ease in processing. They have found extensive use in aerospace, automotive, construction, and recreational

equipment

material

applications.

Research

in

polymer-based

nanocomposites (PNCs) has shown explosive growth in the past decade with much emphasis focused on PNC using thermosetting polymer matrices.

Modifying

conventional FRCs with a PNC matrix result in the emergence of a hybrid composite material termed multiscale fiber-reinforced composites (M-FRCs). The current research addresses carbon nanofiber (CNF) surface modification, carbon nanofiber/epoxy polymer nanocomposite (CNF/PNC) development and M-FRC fabrication. Vacuum assisted resin infusion molding (VARIM) is used to produce MFRC. The materials used in this research are surface modified and modified CNFs, aerospace grade high temperature epoxy resin, and plain-weave E-glass fiber preforms. The VARIM process used to produce M-FRC is explained in detail. The effects of using a CNF/PNC nanophased polymer matrix is presented and thoroughly investigated.

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Flexural, thermomechanical, and rheological studies are presented on the CNF/PNC nanophased matrix and compared with the neat epoxy resin.

Flexural,

interlaminar shear strength (ILSS) and thermomechanical tests are presented for the 0.1 and 1 wt% M-FRCs and compared with the neat FRC. Experimental results indicate that the glass transition temperatures (Tg) of the CNF/PNC and M-FRC samples were higher than the neat epoxy resin and neat FRC samples, respectively.

Coefficients of thermal expansion (CTE) properties of the

CNF/PNC and M-FRC samples were lower than the neat epoxy resin and neat FRC, respectively.

Flexural studies indicated increases in flexural strength (16-20%) and

flexural modulus (23-26%) for M-FRC samples. The ILSS studies indicated increases in ILSS (8-23%) for M-FRC samples.

The improved Tg and CTE properties in the

CNF/PNC samples are believed to be due to achieving good nanoparticle dispersion. While the increased properties in the M-FRC samples are believed to be due to synergistic interactions between the CNF/PNC nanophased matrix and glass-fiber interactions.

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DEDICATION

I dedicate this academic and professional achievement to my Lord and Savior Jesus Christ. In addition, I dedicate this work to my grandparents Russell and Delores McJett. I could not write enough words to express the depth of love that I have for the both of you. Without your guidance and teaching I definitely would not be where I am today, and I’m eternally grateful to God up above to have both of you in my life.

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ACKNOWLEDGEMENTS

The author would like to take this opportunity to sincerely express much gratitude to those people who provided assistance and support during this academic endeavor. First, I would like to express my appreciation to my advisor Dr. Derrick R. Dean, for his bountiful guidance, and invaluable advice throughout my graduate career. I would like to thank the members of my graduate committee, Dr. Uday K. Vaidya, and Dr. J. Barry Andrews for their technical input and attentiveness that they put forth on my behalf. I especially would like to thank my colleagues in the Polymers Research Group (Mohamed A. Abdalla, Moncy V. Jose, Himani D. Deshpande, Ayesha Swarn, Brian W. Steinert, John Tipton, Roberus S. McIntosh and Larry J. Harris) for their constant encouragement and support during this endeavor. I would truly like to thank Dr. Selvum Pillay and Dr. Haibin Ning for their assistance in performing flexural testing. I sincerely appreciate the assistance of Mohamed A. Abdalla for his time in helping me obtain greatly needed electron microscopy micrographs. The financial support rendered from the NSF-NIRT (DMR 0404278) grant and the ARL sponsored project on Thermoplastic Composites for Body Armor Applications. Finally, I would like to thank my family, and true friends Everett D. Ingram, Peter J. Spears Sr., Harry A. Lewis and Shavon L. Ford. The prayers, love, support, and encouragement that you provided me has sustained me through many difficult times and enabled me to fulfill my true destiny.

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TABLE OF CONTENTS Page ABSTRACT ........................................................................................................................ ii DEDICATION ................................................................................................................... iv ACKNOWLEDGEMENTS .................................................................................................v LIST OF TABLES ........................................................................................................... viii LIST OF FIGURES ........................................................................................................... ix CHAPTER 1 INTRODUCTION .........................................................................................................1 Polymer Matrix Materials ..............................................................................................1 Epoxy Resin .............................................................................................................2 Modified Polymer Matrix Materials ..............................................................................7 Polymer Matrix Nanocomposites ............................................................................7 Fibers..............................................................................................................................8 Fiber-Reinforced Composites ......................................................................................10 Multiscale Fiber-Reinforced Composites ....................................................................11 Applications of Fiber-Reinforced Composites ............................................................12 2 LITERATURE REVIEW ............................................................................................13 3 EXPERIMENTAL .......................................................................................................22 Materials ......................................................................................................................22 Carbon Nanofiber Surface Modification Procedure ....................................................23 Epoxy-Carbon Nanofiber/Polymer Nanocomposite Synthesis ....................................24 Fabrication Procedure: Vacuum Assisted Resin Infusion Molding (VARIM) ............25 Dynamic Mechanical Analysis (DMA) .......................................................................27 Thermomechanical Analysis (TMA) ...........................................................................27 Thermogravimetric Analysis (TGA)............................................................................28 Rheology ......................................................................................................................28 X-ray Photoelectron Spectroscopy (XPS) ...................................................................28 Raman ..........................................................................................................................29

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Mechanical Test Methods ............................................................................................29 Flexural Testing (3-Point Bending) .......................................................................29 Short Beam Test (Interlaminar Shear Strength (ILSS)) .........................................31 Morphology..................................................................................................................32 Optical Microscopy (OM)......................................................................................32 High Resolution Scanning Electron Microscopy (Hi-Res SEM) ..........................32 4 RESULTS AND DISCUSSION ..................................................................................33 Carbon Nanofiber Characterization .............................................................................33 Raman ....................................................................................................................33 Thermogravimetric Analysis .................................................................................35 X-ray Photoelectron Spectroscopy ........................................................................29 Epoxy-Carbon Nanofiber/Polymer Nanocomposite Characterization.........................41 Dynamic Mechanical Analysis ..............................................................................41 Thermomechanical Analysis ..................................................................................48 Epoxy-CNF/PNC Processing Analysis ..................................................................52 Morphology............................................................................................................57 Flexural Testing (3-Point Bending) .......................................................................59 Multiscale Fiber-Reinforced Composite Characterization ..........................................63 Flexural Testing (3-Point Bending) .......................................................................63 Interlaminar Shear Strength Testing ......................................................................68 Dynamic Mechanical Analysis ..............................................................................71 Thermomechanical Analysis ..................................................................................74 5 CONCLUSIONS..........................................................................................................75 LIST OF REFERENCES ...................................................................................................77

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LIST OF TABLES

Table

Page

1 Commonly used polymer matrix materials ......................................................................2 2 Comparison of fiber-reinforcements ................................................................................9 3 Properties of EPON 815C epoxy resin ..........................................................................22 4 Properties of PR-19-HT CNFs .......................................................................................23 5 D-peak/G-peak values from Raman spectrometry for surface unmodified and modified CNFs ........................................................................35 6 Atomic concentration of C1s and O1s obtained from XPS spectra .................................39 7 List of samples and their corresponding codes ..............................................................41 8 Glass transition temperature from the tan delta peaks a) HSM, b) HSU, c) MMM, and d) MMU.....................................................................45 9 CTEs (ppm/oC) of mechanically mixed specimens (a) (1.0 & 0.1% wt/wt unmodified and modified CNF/PNCs) and (b) high shear (1.0 & 0.1% unmodified and modified CNF/PNCs) and neat epoxy resin.......................................................................................................49 10 Flow Rates and Shear Rates .........................................................................................53 11 Flexural strength and flexural modulus mechanical properties ...................................62 12 Flexural strength and flexural modulus mechanical properties ...................................63 13 Shear strength and flexural mechanical properties ......................................................68 14 Storage Modulus and Glass Transition Temperature...................................................73 15 CTEs (ppmC-1) of the neat, 0.1 & 1% samples............................................................74

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LIST OF FIGURES

Figure

Page

1 Idealized chemical structure of an epoxide ring ..............................................................3 2 The cure chemistry mechanisms involving epoxy with a) primary, b) secondary, and c) tertiary amines .............................................................4 3 Chemical structure of diglycidyl ether of bisphenol-A (DGEBA) epoxy molecule ..............................................................................................5 4 Idealized cure chemistry of DGEBA epoxy resin and primary amine curing agent ............................................................................................6 5 a) Mechanical Mixing and b) High Shear Mixing apparatuses ....................................25 6 Setup of VARIM process ..............................................................................................26 7 Experimental setup for 3-point bending test ..................................................................30 8 Raman spectra of a) unmodified CNFs b) modified CNFs, showing peaks at 1363 and 1580, referred to as the D and G bands respectively........................34 9 TGA thermograms of surface unmodified and modified CNFs in nitrogen ..................37 10 TGA thermograms of surface unmodified and modified CNFs in air .........................38 11 XPS scan of the oxygen 1s peak of surface modified CNF .........................................40 12 Storage Modulus vs. Temperature curves of 0.1 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites .......................43 13 Tan delta vs. Temperature curves of the 0.1 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites .......................44 14 Storage modulus vs. Temperature curves of the 1.0 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites .......................46 15 Tan delta vs. Temperature curves of the 1.0 wt%

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a) high shear mixed b) mechanically mixed polymer nanocomposites .......................47 16 CTE graphs of a) 1% samples b) 0.1% samples (pa = parallel), (pd = perpendicular).....................................................................................................51 17 Unmodified CNF/PNCs ...............................................................................................54 18 Modified CNF/PNC samples .......................................................................................54 19 Overlay of All CNF/PNC samples...............................................................................55 20 Representative samples of a) 1% unmodified CNF/PNC b) 1% modified CNF/PNC ...........................................................................................58 21 Flexural stress vs. strain of neat, 0.1 and 1% CNF/PNC samples ...............................60 22 Flexural stress vs. strain of neat, 0.1, 0.5, 1, 2, and 4 % CNF/PNC samples................................................................................................61 23 Flexural stress vs. strain of neat, 0.1 and 1% M-FRC samples ...................................64 24 Hi-Res SEM micrographs of a) 0.1% and b) 1% M-FRC samples .............................66 25 Hi-Res SEM micrographs of 1% M-FRC samples a) 4000X b) 4500X ......................67 26 Shear strength vs. strain of neat, 0.1 and 1% M-FRC samples ....................................69 27 Cross-sectional fracture surface optical micrographs of a) neat, b) 0.1% c) and 1% M-FRC samples................................................................70 28 a) Storage modulus vs. temperature curves and b) tan delta vs. temperature curves of neat, 0.1, and 1% M-FRC samples ..................72

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CHAPTER 1

INTRODUCTION

Polymer Matrix Materials

The matrix component of composite material can be a thermosetting, a thermoplastic or a modified matrix resin. The matrix component in composite materials plays a minor role in tensile load-carrying capacity of a composite structure. While the fibers are principal load-carrying members, the matrix is responsible for: 1. Keeping the fibers in the desired location and orientation. 2. Acting as a load transfer medium transmitting mechanical loads from the matrix to the fibers. 3. Providing lateral support against possible fiber buckling caused by compressive loading. 4. Protecting from an adverse chemical environment. 5. Broadening the transmission of stress concentration from broken to intact fibers. 6. Controlling interfacial bond failure between the fiber and matrix ahead of the crack (and normal to the crack) blunts the crack and absorbing additional energy. The table below shows some commonly used polymer matrix materials.

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Table 1. Commonly used polymer matrix materials Thermosetting resins

Thermoplastic resins

Epoxy

Polyether ether ketone (PEEK)

Phenolic

Polyether imide (PEI)

Polyester

Polyphenylene sulfide (PPS)

Polyimide

Polysulfone

Vinyl ester

In general, thermoplastic polymer composites are more advantageous to process and in addition they show better toughness than their brittle thermoset counterparts (1). On the other hand, thermoplastics have some undesirable properties such as creep resistance. This is due to the non cross-linked nature of many linear thermoplastics.

Epoxy Resin In particular, epoxy resin is a very attractive and popular thermosetting polymer matrix because of its widespread usage (i.e., aerospace industry, automotive industry, marine industry, construction materials & sporting goods, etc.). Epoxy resins represent some of the highest performance resins of those available at this time. In general, they out-perform most other resin types in terms of mechanical properties and environmental degradation resistance which leads to their exclusive use in the aerospace industry. One of the most advantageous properties of epoxy resin is their low shrinkage during curing which minimizes internal stresses.

They have excellent adhesive strength and high

mechanical properties which are also enhanced by high electrical insulation and good

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chemical resistance.

They find uses as adhesives, caulking materials, sealants, and

varnishes, as well as laminating resins for an array of industrial applications. The term ‘epoxy’ refers to the epoxide functional group consisting of an oxygen atom bonded to two adjacent bonded carbon atoms forming a three-member ring. The idealized chemical structure of an epoxide ring is shown in Fig. 1 and is easily identified in complex epoxy molecules.

O H2 C

CH2

Figure 1. Idealized chemical structure of an epoxide ring The three-member epoxide ring is quite reactive toward various reactants or cross-linking agents. The properties of a final thermoset depend not only on the structure of the epoxy resin but also on the type and the amount of curing agent. The speed of cure of an epoxy system is governed by the curing process, the type and concentration of the curing agent, and the chemistry of polymerization. For this reason, epoxy resins can easily and quickly be cured at any temperature from 5 oC to 150 oC (1, 2). For instance, aliphatic amine curing agents can cure at room temperature, while aromatic amine curing agents need high temperatures to cure properly. If the curing agent is a primary or secondary amine, the cure reaction will proceed as an addition reaction where one nitrogen-hydrogen group reacts with one epoxy group (3). Reactions that involve tertiary amines result from the unshared electron pair on the nitrogen.

Since secondary hydroxyl groups are not

generated on the product side of the reaction, the resin homopolymerizes forming

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polyether (3). When epoxy reacts with a primary amine curing agent, a secondary amine and secondary hydroxyl group are formed. When epoxy reacts with a secondary amine curing agent, tertiary amine and secondary hydroxyl groups are formed. In both cases with epoxy resin reacting with primary and secondary amine curing agents, secondary hydroxyls are formed and are believed to catalyze the reaction (3). The cure chemistry mechanisms involving epoxy with primary, secondary, and tertiary amines are shown below in Fig. 2.

O a)

R

NH2

H2C

primary amine

b) R2

CH

R

CH2 B

H

H

N

CH2 C

secondary amine OH secondary hydroxyl

epoxide

H

O NH

H2C

CH2 B

CH

R2

CH2 B

secondary amine epoxide

N

CH2 C

tertiary amine

CH2 B

OH

secondary hydroxyl H

O c) R3

N

tertiary amine

H2C

CH

R3

CH2 B

epoxide

N

CH2 C CH2

O n B

quaternary ammonium polyether Figure 2. The cure chemistry mechanisms involving epoxy with a) primary, b) secondary, and c) tertiary amines; where (B) represents the polymer chain One of the unique qualities of diglycidyl ether of bisphenol-A (DGEBA) epoxy molecule is that it contains two ring groups at its center which enable the molecule to absorb mechanical and thermal stresses better than linear groups and provide tremendous

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stiffness, toughness and heat resistant properties (1, 2, 4). The chemical structure of the DGEBA epoxy molecule is shown in Fig. 3 below.

O CH2 CH CH2

CH3 O

C

O O CH2

CH CH2

CH3 Figure 3. molecule.

Chemical structure of diglycidyl ether of bisphenol-A (DGEBA) epoxy

The position of the epoxide rings in the epoxy molecule is an additional important factor in the determination of epoxy reactivity. The chemistry of this reaction means that there are usually two or more epoxy sites binding to each amine site. Both the number and distance between reactive groups can affect the material’s performance. The idealized cure chemistry of DGEBA epoxy resin and a primary amine curing agent can be seen below in Fig. 4 (4). This reaction is responsible for forming the three-dimensional molecular structure.

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O

O R

O CH2

CH

CH2

CH2

CH CH2

CH2

CH CH2

O

R

NH2 R NH2 R

O CH2

CH

CH2

OH

OH O CH2

CH

R

O

O

R

O

CH2

CH2

CH CH2

CH2

CH CH2

O

R

N R N R

O CH2

CH

CH2

O

R

OH

OH

Figure 4. Idealized cure chemistry of DGEBA epoxy resin and primary amine curing agent.

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Modified-Polymer Matrix Materials

Polymer Nanocomposites A nanocomposite can be defined as a composite material that has one or more of its constituents with nanoscale dimensions (i.e., on the scale of 1 billionth of a meter). Polymer nanocomposite research has attracted international interest since the late 1980s due to their unique property enhancements relative to the virgin polymer.

The

incorporation of nanoscale constituents into polymer matrices leads to novel or modified property enhancements significantly greater than that attainable using conventional fillers or polymer blends. Improvements in mechanical strength, fracture toughness, thermal (i.e., thermal stability, decomposition, and coefficients of thermal expansion (CTE)), and physical (i.e., permeability, optical, dielectric) properties have been observed depending on the type of nanoparticle (i.e., CNT, CNF, layered silicate, etc.). A key feature of polymer nanocomposites is that such enhancements are achievable at very low volume fractions of 2-5% “nanoconstituent.” Where in conventional reinforcing composites the polymer matrix may account for as much as 50% by weight and typical structural reinforcements may be as high as 60% by volume. Such low volume fractions offer a potential for incorporating nanoconstituents into existing polymer resins without significantly modifying existing processing techniques, and increases the tailor-ability options for advanced composite matrices. Hence, polymer nanocomposites provide the opportunity to explore new behaviors and functionalities beyond those found in conventional materials.

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Fibers

The fibers are the principal load-bearing members in FRCs. The proper selection of the type, amount, and orientation of the fibers are very important since it influences the following composite laminate characteristics: 1. Density. 2. Tensile and compressive strength. 3. Fatigue strength and fatigue failure mechanisms. 4. Electrical and thermal conductivity. 5. Cost. The fibers in FRCs have specific roles such as: 1. They support all main loads. 2. Limit deformations. 3. Increase overall mechanical properties (i.e., strength, toughness, and stiffness). 4. Decrease corrosion, creep, and fatigue. The principal fibers in commercial use are various types of glass (i.e., E-glass, S-glass, and C-glass) (1, 2), carbon fiber (i.e., pitch-based, or PAN-based) (1), boron (1, 2), silicon carbide (SiC) (1, 2), ceramic fiber (i.e., alumina-based (Nextel)) (1), polymerbased (i.e., aramid (Kevlar 49, Nomex), and polyethylene (Spectra, Tekmilon) (1). Fibers used in FRCs for structural applications are mostly used in laminate form. This is achieved by stacking a number of fiber preforms and matrix and then consolidating them to a desired thickness. Fiber orientation in each layer and stacking sequence can be

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controlled to obtain desired physical and mechanical properties. Table 2 below shows a comparison of various fiber-reinforcements.

Table 2. Comparison of fiber-reinforcements Fiber

Density, ρ (g/cm3)

Tensile Strength, σ (GPa)

Tensile Modulus, E (GPa)

Poisson’ s ratio, υ

CTE, α (10-6 K-1)

E-glass

2.55

1.5-2.5

70

0.2

5

Carbon

1.75

2.7

250

0.2

-1

Kevlar

1.45

3.6

125

0.35

-2

Boron (W)

2.7

3.8

393

0.2

5

Al2O3

3.95

1.9

379

0.2

7.5

Of all of the principal fibers in use in FRCs, glass-fibers are the most common selection for polymer matrix composite (PMC) materials. A major factor for the use of glass-fibers in commercial applications is due to their relatively low cost when compared to other commercial fibers. They have CTE properties (2-5 (10-6 K-1), which are lower than many metallic fibers or filaments (i.e., steel, aluminum, titanium, nickel, gold, etc.). Two types of glass fibers used in polymer FRCs are E-glass and S-glass. E-glass has the lowest cost of all commercially available fibers, which is the reason for their widespread use in FRC materials. The letter ‘E’ is used because of their original use in electrical applications. S-glass was originally developed for aerospace applications (i.e., aircraft components and missile casings). The letter ‘S’ is used due to their use where high tensile strength properties are desired. At room temperature S-glass is 10-15% stronger than E-glass, and can withstand higher in-use temperatures (1, 2). Another type of glass

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fiber, known as C-glass, is used in chemical applications that require greater resistance to chemical attack than is supplied by E-glass.

Fiber-Reinforced Composites

The history of fiber-reinforced polymer composites date back to work started in the USA in the 1940s (1). It refers to a material that contains high strength and modulus fibers embedded in a matrix with distinct interfaces or interphases between them. The fibers and matrix retain their physical and chemical identities, yet produce a combination of properties unattainable by either constituent. Fiber-reinforced composite materials have shown great promise as high strength materials due to their potential benefits of low density, high strength, high stiffness-to-weight ratio, economic feasibility, good corrosion resistance, low CTE, high thermal damping capacity, excellent durability, design flexibility and excellent in-service applications. They have found use in construction, aerospace, automotive, and recreational equipment material applications.

The main

advantage of most FRC materials is in the weight savings. Incorporating CNFs in the polymer matrix used in making FRCs can have some added benefits such as increased strength and fracture toughness. Carbon nanofibers (carbon fibers with dimensions on the scale of 1 billionth of a meter) are currently being used in various approaches to provide functional (thermal, electrical and magnetic conductivity etc.) and structural property enhancement to polymer based materials including composites (5).

They have exceptionally high thermal and electrical

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conductivity which make them excellent candidates for polymer nanocomposites in applications demanding heat dissipation and thermal shock resistance. They also offer the opportunity to create FRCs with multifunctional properties. Meaning that FRCs would be able to perform mechanically, but also be able to conduct electrically or heat. The potential of carbon based nanocomposites as the matrix material in FRC structural composite materials has not been fully exploited. Carbon nanofiber usage is growing due to their low cost and increased performance in a variety of applications as polymer nanomaterials in the automotive and aerospace industries. Fiber-reinforced composites with multifunctional properties show great potential for use in future automotive, aerospace, and marine applications.

Multiscale Fiber-Reinforced Composites

Synergistically polymer nanocomposites with conventionally filled composites (i.e., glass-fiber, carbon-fiber) have resulted in the multiscale composite materials that possess unique properties. Multiscale composites possess a hierarchal microstructure ranging from nanoscale fibers to micron size fibers. There is a strong need to balance the multiple demands of performance, weight, risk and life-cycle cost in selecting new structural materials. However, multiscale fiber reinforced composite materials (M-FRC) offer a route of creating materials that exhibit multifunctionality (i.e., enhanced thermal stability, lower CTE, high thermal and electrical conductivity, etc.) without sacrificing performance.

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Applications of Fiber-Reinforced Composites

Boeing’s 787 Dreamliner is the first aircraft to have its fuselage made entirely from polymer composite matrix composite materials, and Airbus’ A380 is the largest commercial passenger aircraft ever built and is comprised of 25 weight-percent of composite materials, specifically with polymer composite materials in the fuselage, wings and tail sections.

Both Boeing and Airbus are revolutionizing the aircraft building

industry by introducing polymer composite materials as the primary structural components in their aircraft, making the aircraft more fuel efficient than ever before. Epoxy resins have found use in the marine industry for building boats because of their excellent adhesive properties as a laminating resin and resistance to water degradation (2, 4, 6). Currently epoxy resins are supplanting use of polyester resins as the polymer matrix material in the boating industry because of the environmental degradation caused by water. In the space shuttle, for example, the total weight savings with FRCs is 2688 lb per vehicle (2).

Thermoset polymers, such as epoxy resins, and polyesters have

widespread commercial use for making FRCs due to their ease of processability. They are often used in high performance, high temperature aerospace applications.

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CHAPTER 2

LITERATURE REVIEW

Carbon nanofibers (CNFs) have been recognized as a promising constituent in polymer nanocomposites (PNCs) because they exhibit characteristics similar to their carbon nanotube (CNT) counterparts, such as excellent electrical, thermal, and mechanical properties (7, 8). Carbon nanofibers typically have a larger average diameter (100 ~ 300nm) than CNTs (1 ~ 50nm), and CNFs have a slightly lower modulus of only a few tens of GPa, and are more inexpensive to produce compared to CNTs (8-11). Carbon nanofibers have a one-dimensional morphology consisting of sp2-bonded graphitic carbon oriented along an axis parallel to the basal plane. Carbon nanofibers as a nanofiller in PNCs are quite useful in applications which employ CNTs because CNFs, like their CNT counterparts, exhibit a high aspect ratio (length/diameter, L/D > 100), an interfacial area which can reach 100-1000 m2/g, and chemical inertness (10). Investigations are underway exploring the nature and thermal behavior of PNCs (7-12). One of the major technological issues limiting widespread use of epoxy CNF/PNCs is processability, where CNF dispersion or alignment is crucial to CNF/PNC performance.

For example, one of the problems is that CNFs have a tendency to

aggregate together, resulting in poor dispersion throughout the polymer matrix. Carbon nanofiber aggregation is due to Van der Waals interactions. If CNF-polymer matrix

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adhesion is poor, the CNF/PNC may fail at the interface, leading to mediocre mechanical property enhancements (13). A key in PNC processing, is nanofiller/polymer interfacial chemistry. Recent studies suggests that coupling the dispersion of the nanoconstituent into polymer host and controlling the interfacial chemistry between the polymer host and the nanoconstituent can provide a sufficient means of creating such multi-functional materials (7, 9, 10, 12, 14). In systems employing silicates and CNTs, it has been suggested that modifying their chemical surfaces increases their overall miscibility into polymer hosts and provide increased thermomechanical and bulk properties (7, 9, 10, 12, 14-20). The chemical nature of the surface modifying molecules has proven that they will affect the ease with which CNFs can be homogeneously dispersed in the epoxy resin. Modifying the surfaces of the CNFs plays a vital role in disrupting the aggregating nature of CNFs and simultaneously enhancing the solubility of CNFs within the polymer matrix which can have a profound effect on the load transfer in the nanocomposite. Recent studies have shown that CNF surface modification aids in the dispersion of CNFs into the polymer matrix and leads to increased mechanical properties (9, 10, 14). Rice et al. reported a 34-percent increase in strength and a 130-percent higher modulus with a 4 weight percent CNF loaded epoxy CNF/PNC (14). Thus, good CNF-polymer adhesion can promote interactions between the CNF and polymer matrix which should produce desirable mechanical properties. Toebes et al. reported that the hydrophobic and inert nature of CNFs are unfavorable for some applications, but by introducing the CNFs to an oxidizing acid such as sulfuric/nitric acid (H2SO4/HNO3) can produce CNFs with oxygen-containing surface groups and enhance the CNFs affinity for polar solvents (13). Glasgow and co-workers reported that surface oxygen levels as high as 25 atom % can be

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achieved from using the (H2SO4/HNO3) oxidation surface treatment (12). The oxygen on the surfaces of the CNFs will be in the form of carboxylic acids. Other studies have reported CNF surface treatments that consist of 1:3 (H2SO4/HNO3) (12, 13, 19) , air (12), peracetic acid (12), and cold oxygen plasma treatment (21), with (H2SO4/HNO3) yielding the best results. Recent studies reported that Thermogravimetric analysis (TGA) can be used to quantitatively determine the degree of surface functionalization of fullerenes and single-walled carbon nanotubes (SWNTs) in an argon atmosphere (22, 23). Sayes et al. investigated the functional density dependence of SWNT cytotoxicity in vitro (23). They functionalized SWNTs with three different solvents.

Each solvent provided

functionalization of the SWNT surfaces from the least to the most surface coverage of functional groups. The degrees of functionalization for three samples were 80, 41, and 18 for the least, medium, and most functionalized SWNT respectively. Thermoplastic (TP) CNF/PNCs have been reported for polymers based on systems

employing

isotatic

polypropylene

(iPP),

polycarbonate

(PC),

polyetheretherketone (PEEK), polymethylmethacrylate (PMMA), polyphenylene sulfide (PPS) and nylon as polymer matrices (9, 10, 12, 13, 20). Thermoplastic CNF/PNCs as electrostatic discharge materials are needed for electronic packaging purposes, for garments, work benches, and personnel working in an electronic packaging facility (20). Lozano et al. reported electrical resistivity on iPP-CNF/PNCs which showed that as CNF loading increased from (0 to 40 wt.%), the electrical resistivity decreased 10-fold, suggesting that the CNFs can increase conductivity and increase the electrostatic dissipation properties of TP-CNF/PNCs (20). Other studies have reported thermoset (TS) PNCs using CNFs employing different epoxy resins (low viscosity, high viscosity, room-

15

temperature, and high temperature) and bismaleimides (BMIs) as polymer matrices (9, 10, 12, 14, 16). Both Gauthier et al. (10) and Choi et al. (16) studied epoxy CNF/PNCs. Choi et al. reported strength improvements with 5% CNF loadings (16). Gauthier et al. reported higher strength improvements with strain to break increasing 100% and 160% increase in stress (10). The strain to break significantly increased for Gauthier et al. but when compared to Choi et al. study, the strain to break was relatively moderate. This indicates that the data from both cases studying the mechanical properties of epoxy CNF/PNCs are different from one another due to differences in the epoxy resin composition. The vastly improved enhancement seen in Gauthier et al.’s study indicates strong CNF and polymer matrix interactions are present. Studies by Patton et al. (9) and Choi et al. (16) investigated the effects of using epoxy resins with different viscosities (i.e., high viscosity, low viscosity) in making CNF/PNCs. Patton and co-workers reported that the flexural modulus increased from 2.3 GPa to 8.74 GPa as the CNF loading increased from (0 to 35.5% vol/vol (Vf)) and increased flexural strength up to 20% as the Vf increased to 15% vol/vol and decreased at higher volume fractions for their low viscosity epoxy CNF/PNC (9). The decrease in flexural mechanical properties in their high volume fraction CNF/epoxy samples is due to increased void content and poor dispersion. Choi et al. reported increased glass transition temperatures (Tg) of 26 oC with 20 wt.% CNF loading with low viscosity epoxy resin and a decrease in Tg with 20 % CNF loading with high viscosity resin (16). This suggests that the high loaded CNF/epoxy samples made with high viscosity epoxy resin showed a decrease in Tg presumably due to poor dispersion. In a study investigating thermal degradation of CNF/PNCs, Choi et al. reported that the samples began to decompose at

16

approximately 260oC (first stage), and were completely decomposed at around 640oC, 720oC, 730oC, and 735oC (third stage) for pure epoxy, epoxy/5 wt.%, 10 wt.%, 20 wt.% CNF nanocomposite respectively (16). The thermal conductivity measurements for the epoxy resin and CNFs are inherently different with epoxy resin and CNFs showing typical thermal conductivity measurements at ~0.1 W/mK, and 1950 W/mK respectively (9).

CNFs have low

electrical resistivity ~55 µΩcm (9), while epoxies on the other hand have very high electrical resistivity and are known to be insulators (16). A thermal conductivity study by Chen et al. on epoxy CNF/PNCs showed that as CNF concentration increased from (252% vol/vol) that thermal conductivity increased from zero to 695 W/mK respectively (15). In a study on electrical resistivity of epoxy CNF/PNCs Choi et al. show that as the CNF volume fraction increases; the electrical resistivity decreases in both the low and high viscosity epoxy CNF/PNC systems, with the low viscosity epoxy CNF/PNC samples showing greater sensitivity (16). The use of a low viscosity epoxy resin in making epoxy CNF/PNCs suggests that the CNFs should disperse better throughout the polymer matrix and possibly produce multifunctional polymer nanocomposites with the ability to pass electric current. Most of the data reported on CTE, thermal conductivity, electrical resistivity and thermal properties have been on unmodified CNF reinforced epoxy CNF/PNCs while little data exists on modified CNF reinforced epoxy CNF/CNFs. Abdalla et al. studied the effect of layered silicate (LS) loading on the CTE properties of PMR-15/layered silicate nanocomposites (24).

Their 2.5% PMR/layered silicate sample exhibited a

reduction of 26% in the CTE property relative to their neat sample (24). Chen and co-

17

workers reported as CNF concentration increased from (2-52% vol/vol) the CTE of epoxy CNF/PNCs decreased from 31.29 ppmK-1 to -0.11 ppmK-1 respectively (15). A similar study by Sullivan et al. investigated the CTE properties ceramic nanoparticle/polyimide PNCs, and they noted as the ceramic nanoparticle loading increased from (0 to 6 wt.%) the resultant CTE values decreased, with the 3 wt.% sample showing 18% reduction in CTE property (25). A recent study investigating nanoparticle alignment using uniaxial magnetic fields paid close attention to CTE behavior and showed reduced CTE values in the direction of the maximum layered silicate alignment (26, 27) has led our group to investigate similar behavior with our PNCs. A critical challenge for M-FRC is the integration, control, and exploitation of nanoparticle-enabled properties within a hierarchal structural composite made with commercially processing methods. Nanoparticles can be incorporated in a multitude of different ways within traditional composite materials, (i.e., within a fiber, thin films, part of the polymer matrix, etc.). Many of the limitations of fabricating M-FRC is directly related to achieving a homogenous distribution of the nanoparticles throughout the bulk of the composite (14, 28-35).

For example, in a FRC fabrication process such as

vacuum-assisted resin infusion (VARIM), nanoparticles are sometimes filtered and can lead to a inhomogeneous microstructure (33, 36). In order to take full advantage of the unique mechanical properties of nanoparticles (i.e., CNTs, CNFs, etc.), much effort is needed to improve processing techniques to achieve high-performance M-FRCs with fully integrated and well dispersed nanoparticles.

Control of the dispersion of the

nanocomposite matrix is extremely critical in achieving property enhancements and making the M-FRC a multifunctional material.

18

Bekyarova et al. (28) and Qui et al. (33) used CNTs as the nanoconstituent in fabricating M-FRC. Bekyarova et al. studied the effect of electrophoretically dispersing 0.25 wt% functionalized CNTs into carbon fiber preforms and fabricating M-FRC using the VARIM process and showed 30% enhancement in interlaminar shear strength (ILSS) (28). The 30% increase in ILSS is suggestive of improved matrix dominant properties in the M-FRC. In a study investigating the epoxy/CNT/glass fiber M-FRC, Qui et. al reported a 16% and 27% increase in tensile strength and modulus, respectively for their 1 wt% functionalized CNT M-FRC sample (33). The mechanical property enhancements in their study were due to the use of functionalized CNTs amplifying the dispersion throughout the polymer matrix and strengthening CNT-polymer matrix interactions. Dean et al. (30), Haque et al. (31), Chowdury et al. (29), and Kornman et al. (32) investigated the effects of using LS as the nanoconstituent in fabricating M-FRC. Dean et al. investigated the effect that the LS loading had on mechanical properties of VARIM processed multiscale fiber-reinforced nanocomposites (30). They reported enhancements of 24% and 31% in flexural strength and modulus, respectively for their 2 wt% M-FRC sample. Improvements in flexural strength and modulus in the 2 wt% sample are due to a homogenous dispersion of the LS and low viscosity of the 2 wt% epoxy/LS nanocomposite. Rheological characterization by Dean et al. show that their 2 wt% nanocomposite is well within the range of fabricating FRCs using the VARIM process and is suggestive that its low viscosity can lead to better wetting of the fibers and dispersion of the nanoparticles.

The effect of LS loading on the flexural and

thermomechanical properties of VARIM processed epoxy/LS/carbon fiber M-FRC was studied by Chowdhury et al. (29). Chowdhury and co-workers reported a 25%, 14%, and

19

30% enhancement in flexural strength, modulus and ILSS, respectively. The increased mechanical properties are attributed to enhanced polymer-nanoparticle-fiber interfacial interactions. Haque et. al studied the effect of LS loading on the improvement of mechanical and thermal properties of VARIM processed epoxy/LS/S2-glass M-FRCs (31). They showed by dispersing 1 wt% LS that the resultant epoxy/LS/S2-glass M-FRC exhibited an enhancement of 24%, 14%, and 44% in flexural strength, modulus, and ILSS. It is to be noted that ILSS is the matrix dominant property and such improvements in ILSS of M-FRC is mostly due to the improved property of the epoxy/LS nanocomposite. The mechanical property enhancements in strength, modulus and ILSS are due to increased interfacial areas, bond characteristics and unique phase morphology of epoxy/LS nanocomposites.

Kornman et al. studied the effect of LS loading on

VARIM processed epoxy/LS/glass-fiber M-FRCs (32). They reported improvements of 27% and 6% in flexural strength and modulus for their 10 wt% LS M-FRC sample. The flexural strength and modulus improvements were due to the presence of the LS at the surface of the glass fiber which improved the interfacial bonding between the epoxy/LS nanocomposite matrix and glass-fiber. In a study investigating the dimensional stability of epoxy/CNT/glass-fiber MFRC, Qui et al. (33) reported that their 1 wt% M-FRC sample exhibited a 25% reduction in CTE relative to their neat sample. The CTE reduction was suggestive to having well dispersed the CNTs throughout the M-FRC. Dean et al. (30) and Haque et al. (31) reported very small increases in Tg in their VARIM processed M-FRC samples. On the contrary, Qui et al. (33) reported a 11 °C reduction in Tg of their 1 wt% M-FRC sample relative to the neat FRC sample. The Tg reduction was suggestive that the cross-link

20

topology decreases due to interference of the functional groups on the CNTs during curing. As a result, the presence of the functionalized CNTs disrupts the optimized epoxy resin-curing agent ratio in the curing reaction. Kornmann et al. studied the effect of LS on moisture properties of VARIM processed epoxy/LS/glass-fiber composites (32). They were able to show that their 10 wt% sample exhibited greater sensitivity to moisture uptake at 50°C and is the cause for the apparent decrease in Tg. Studies by Bekyarova et al. (28) and Qui et al. (33) used 4Point probing to investigate the in-plane and out-of plane electrical conductivity in the M-FRC samples. Bekyarova et al. (28) reported a 30% increase in out-of-plane electrical conductivity for their epoxy/MWNT/carbon fiber and a 200% increase in out-of-plane conductivity in the SWNT M-FRC samples. The substantial increase in the SWNT was due to a network of interconnected SWNT bundles forming conducting paths in narrow polymer regions. The differences in the electrical conductivity measurements can be attributed to differences in sample thickness, but may also differ because of the differences between SWNT and MWNT morphologies (28). In the study by Qui et al., a nominal enhancement was exhibited in their functionalized 1 wt% M-FRC sample, but not as significant as of using pristine SWNTs (33). The nominal electrical conductivity enhancement was due to the functionalization disrupting the CNT electronic properties which led to decreased electrical properties. Differences between both studies may also be due to the type of fiber reinforcement used during fabricating M-FRCs. For example, carbon fibers possess excellent electrical conductive properties; while on the contrary, glass-fibers behave as insulators.

21

CHAPTER 3

EXPERIMENTAL

Materials

The materials used in this investigation were, EPON 815C (Modified diglycidyl ether of bisphenol A (DGEBA) Epoxy resin) and EPICURE W curing agent (Aromatic diamine) were purchased from the Miller-Stephenson Chemical Company. The epoxy resin and curing agent were mixed in the ratio of 100:20 and cured at 150 oC for 2-hours. Table 3 below shows the properties of EPON 815C epoxy resin system.

Table 3. Properties of EPON 815C epoxy resin Property

Value

Units

Viscosity

5-7

Poise

Density

1.13

g/cm3

Tensile Strength

60

MPa

Tensile Modulus

2.0

GPa

80-90

ppm/K

Electrical Resistivity

Insulator

µΩ-cm

Thermal Conductivity

0.230

W/m-K

Coefficients of Thermal Expansion

22

The CNFs (heat treated grade PR-19-HT) were kindly provided by (Pyrograph III™ from Applied Sciences Incorporated). Table 4 below shows the properties of PR-19-HT CNFs.

Table 4. Properties of PR-19-HT CNFs Property

Value

Units

100-200

nm

Density

2.1

g/cm3

Tensile Strength

7.0

GPa

Tensile Modulus

600

GPa

Coefficient of Thermal Expansion

-1.0

ppm/K

Electrical Resistivity

55

µΩ-cm

Thermal Conductivity

1950

W/m-K

Fiber Diameter

Carbon Nanofiber Surface Modification Procedure

The CNFs were refluxed in methylene chloride (CH2Cl2 ) at 35oC over a duration of 7 days. Following refluxing they were vacuum filtrated and washed with distilled water to remove any surface residue. Immediately following washing, they were placed into a large volumetric flask containing 3:1 concentrated sulfuric acid/ nitric acid (H2SO4/HNO3), and sonicated in a water bath for three hours at ambient temperature. Next, the CNF’s were vacuum filtrated and washed with excessive amounts of distilled water until no residual acid was present. The resulting CNFs were collected and dried in an oven at 120oC overnight. After drying, the CNFs were ground into powder form.

23

Epoxy-Carbon Nanofiber/Polymer Nanocomposite Synthesis

Epoxy CNF/PNCs were prepared using 2-dispersion methods.

In the first

dispersion method the appropriate amount of CNFs (0.1, or 1.0 wt %) were dispersed into epoxy resin and mixed with a mechanical mixer for 30 min. The appropriate amount of curing agent was added to the epoxy resin/CNF mixture and mixed for an additional 5 minutes. The nanodispersed prepolymer was poured into a steel mold, and degassed under vacuum to remove trapped air bubbles, followed by curing in an oven at 150oC for 2 hours. Fig. 5a below shows the mechanical mixing apparatus. In the second dispersion method (high shearing), the appropriate amount of CNFs (0.1, or 1.0 wt %) were dispersed into the epoxy resin. The epoxy resin/CNF mixture was poured into a cylindrical syringe with a small orifice, followed by high shear mixing, in which the polymer/CNF mixture is sheared under high pressure through the small orifice. The appropriate amount of curing agent was added and additionally sheared for 7-10 times. The high shear mixing apparatus is shown in Fig. 5b. The nanodispersed prepolymer was poured into a steel mold, and degassed under vacuum to remove trapped air bubbles, followed by curing in an oven at 150oC for 2 hours.

24

a)

b)

Figure 5. a) Mechanical Mixing and b) High Shear Mixing apparatuses

Fabrication Procedure: Vacuum Assisted Resin Infusion Molding (VARIM)

Vacuum assisted resin infusion molding (VARIM) is a very attractive, cost effective and environmentally friendly method of processing composites. The dry fabrics (E-glass plain weave) were cut into plies and stacked several layers to form the dry preform. During the stacking, the weft and warp fibers were carefully aligned so that each layer preserves the same directions of reinforcement. The fabric was placed in an open mold with a peeling cloth between the fabric and the mold surface. After laying up the desired number of fabric layers, another peeling cloth and distribution mesh was 25

placed over the preform.

One resin-suction line and one resin-injection line were

attached to the mold surface. The resin-suction line was connected to a resin trap that was connected to a vacuum pump. A vacuum bag was placed over the open mold to enclose all the contents on the mold surface. A typical setup of the VARIM process is illustrated in Fig. 6 below.

Figure 6. setup of VARIM process The epoxy resin and curing agent were mixed according to the prescribed 100:20 ratio. Prior to resin injection, the vacuum was applied to the mold for thirty minutes to debulk the preform. Debulking is done to remove any air within the VARIM setup. After debulking, the resin was infused via vacuum (28.5 to 30 inch mercury) to completely “wet out” the reinforcements and eliminate all air voids in the laminate structure. The

26

resin flowed and wetted the preform in-plane and through the thickness of the preform. After the preform was completely wetted by the resin, the resin-injection line was clamped and blocked off and the preform was left under vacuum at room temperature for one hour. Uniform wettability was attained. The preform was then removed from the VARIM setup, consolidated with a Carver hydraulic hot press and cured at 150 oC for 2 hours.

Dynamical Mechanical Analysis (DMA)

The glass transition temperature (Tg) of the CNF/PNC samples were obtained using a TA Instruments 2980 Dynamical Mechanical Analyzer. To find the Tg of the materials, the temperature was ramped from ambient to 200 oC at a rate of 5 oC min-1. The Tg was determined by the corresponding peak of the tan delta (tan δ) curve.

Thermomechanical Analysis (TMA)

CTE measurements of the CNF/PNC samples were obtained using a TA Instruments Q400 Thermomechanical Analyzer. To find the CTE of the materials, the temperature was ramped from ambient to 200 oC at a rate of 10 oC min-1 and under a load of 0.5 N. The CTE measurement was taken below the Tg of the material.

27

Thermogravimetric Analysis (TGA)

The thermal stability of the CNFs was studied using a TA Instruments 2950 Thermogravimetric Analyzer. The procedure involved a temperature ramp from ambient to 1000 oC at a rate of 20 oC min-1 in air and nitrogen.

Rheology

Rheological experiments were performed on the nanodispersed prepolymers using a TA Instruments AR 2000 Rheometer. Time sweeps were performed at a strain rate of 0.1% and frequency of 1 Hz.

X-ray Photoelectron Spectroscopy (XPS)

X-ray photoelectron spectroscopy (XPS) was used to characterize the effect of chemical surface modification of the CNFs. The surface chemistry was studied with a Surface Science Instruments (SSI) M-probe XPS instrument operated at a base pressure of 8 x 10-10 torr. Binding energy positions were calibrated using the Au 4f7/2 peak at 83.93 eV, and the Cu 3s and Cu 2p3/2 peaks at 122.39 and 932.47 eV, respectively.

28

Raman Spectroscopy

Raman spectroscopy was used to characterize the effect of chemical surface modification of the CNFs. Raman spectroscopy was carried out using a backscattering geometry. The wavelength of 422 nm line of Kimmon Electric’s HeCd laser was used to excite the sample. The laser beam with a nominal power of 80 mW was focused onto a spot 5 µm in diameter.

Mechanical Test Methods

Flexural Testing (3-Point Bending) Flexural properties of neat and CNF/PNCs, and the multiscale FRCs were measured using a SATEC Apex T5000 mechanical testing instrument. A crosshead speed of (1.15 mm min-1) and a span length of (48 mm) was used according to ASTM D790. The ASTM D790 method was used to determine the flexural strength of the CNF/PNCs and multiscale FRCs. The support span to depth ratio (L/d) was 16 to 1. The experimental setup for the three-point bend is illustrated in Fig. 7 below.

29

Figure 7. Experimental setup for 3 point bending test Flexural strength, σ f is determined for three-point bending from the relationship:

σf =

3PL 2bd 2

(1)

where:

σ f = flexural strength, MPa (psi) P= L= b= d=

failure load, N (lbf) support span of beam, mm (in) width of beam tested, mm (in) depth of beam tested, mm (in)

It should be noted that the above equation applies directly to materials for which the stress is linearly proportional to strain up to the point of rupture and for which the strains are small. Since this is not the case in reality, a slight error will arise when the equation is used. It yields an apparent strength based on homogenous beam theory.

30

Short Beam Test (Interlaminar Shear Strength Test (ILSS)) The ASTM D 2344 test method was used to determine the interlaminar shear strength of the neat and CNF/PNC multiscale FRCs. A crosshead speed of (1.00 mm (0.05) min-1) was used. The support span to depth ratio (L/d) was 4 to 1. As shown in Figure 7, this test method involves loading a beam under a three-point bending fixture with a span to depth ratio (L/d) selected such that interlaminar shear failure is induced rather than failure due to bending. For short beam shear test, the apparent ILSS, F sbs is determined from the relationship: F sbs =

3P 4bd

(2)

where: F sbs = flexural strength, MPa (psi) P = failure load, N (lbf) b = width of beam tested, mm (in) d = depth of beam tested, mm (in) It should be noted that proper precautions were taken to prepare the samples for testing. A water-lubricated saw was used to obtain the samples final dimensions and the samples were free of undercuts, notches, delaminated surfaces, and rough surfaces.

31

Morphology

Optical Microscopy (OM) Optical microscopy was used to characterize the fracture surfaces of FRC samples to confirm that delamination of the glass fiber laminates occurred at the glass fiber laminate-resin interface during the ILSS test.

Optical micrographs of the M-FRC

samples were taken using a AXIOPLAN 4MP stereomicroscope.

High Resolution Scanning Electron Microscopy (Hi-Res SEM) High resolution scanning electron microscopy was used to analyze the microstructure of the fracture surfaces of the CNF/PNCs and FRCs. This technique provides an excellent view of the topology of fracture surfaces. SEM micrographs of a freeze fracture surface of CNF/PNC and FRC samples were taken using a JEOL-7000F with a 10-15 kV accelerating voltage.

32

CHAPTER 4

RESULTS AND DISCUSSION

Carbon Nanofiber Characterization

Raman Spectroscopy The effect of CNF surface oxidation was investigated using Raman spectroscopy, because Raman spectroscopy offers useful information concerning the structural changes of CNFs, especially changes due to sidewall modification (37, 38). Fig. 8 shows the Raman spectra of the (a) surface modified CNFs, as well as the (b) surface unmodified CNFs. The peak at 1363 cm-1 is related to the disordered density of the sp3-hybridized carbon atoms of the nanofiber and is designated as the D band, whereas the peak at the 1580 cm-1 is related to the ordered graphitic structure of sp2-hybridized carbon atoms of the nanofiber and is designated as the G band. The D/G band ratio would then represent the amount of sp3-hybridized carbon defects in the nanofiber due to the presence of functional groups.

33

5500

Intensity (Counts)

5000

4500

4000

3500

1100

1200

1300

1400

1500

1600

R a m a n S h if t ( c m

a)

-1

1700

1800

1900

1700

1800

1900

)

6500

Intensity (Counts)

6000

5500

5000

4500

4000

3500 1100

b)

1200

1300

1400

1500

1600

R a m a n S h ift ( c m

-1

)

Figure 8. Raman spectra of a) unmodified CNFs b) modified CNFs, showing peaks at 1363 and 1580, referred to as the D and G bands respectively.

34

These defects (or functional groups) on the CNF now become potential reactive sites for bonding with the polymer matrix. Thus, an increase in the D band intensity is indicative of adding functional groups on the surfaces of the CNFs. Table 5 shows the D/G values for both surface unmodified and modified CNFs. From these values it can be presumed that the surface modification is due to the carboxylic acid oxidation that occurred.

Table 5. D-peak/G-peak values from Raman spectrometry for surface unmodified and modified CNFs Sample

D Peak intensity

G Peak intensity

D/G value

Unmodified

239

1377

0.1735

Modified

566

2545

0.2223

Thermogravimetric Analysis Thermogravimetric analysis (TGA) measures the amount and rate of change in the weight of a material as a function of temperature or time in a controlled atmosphere. Measurements are used primarily to determine the composition of materials and to predict their thermal stability at temperatures up to 1000°C.

The technique can

characterize materials that exhibit weight loss or gain due to decomposition, oxidation, or dehydration. The total weight loss from the TGA experiments showed the effect that carboxylic acid surface modification had on the CNFs results in a nitrogen atmosphere was shown in Fig. 9 and results for tests in air are shown in Fig.10. The surface unmodified CNFs showed no weight loss up to 777oC in nitrogen and suffered only a 0.07% weight loss up to that temperature. The surface modified CNFs showed a 1.35%

35

weight loss up to 820oC in air. The surface modified CNFs in nitrogen exhibited a 5% weight loss on heating to 632oC.

Interestingly, the surface modified CNFs in air

exhibited a 7% weight loss up to 769oC. The surface modified CNFs in both atmospheres exhibited lower thermal stability than the unmodified CNFs suggestive of the introduction of the carboxylic acid surface modification. The lower thermal stability of the modified CNFs can be attributed to the functionalization moieties from the CNFs surface. TGA is useful for quantitatively determining the degree of functionalization of CNFs.

Sayes et al. used TGA to quantitatively determine the degree of surface

functionalization of fullerenes and SWNTs in an argon atmosphere (22, 23). determine the degree of functionalization, this method uses the ratio:

To

Wl Wi -Wl : , Wmol WC

where Wl is the weight loss of the grafted molecule, Wmol is the molecular weight of the grafted functional group, Wi is the initial weight of the modified CNF and WC is the molecular weight of carbon. The weight loss of the grafted molecule and the initial weight of the modified CNF values are taken from the TGA thermogram. The molecular weight of the carboxyl functional group is 45 g/mol, and the molecular weight of carbon is 12 g/mol. The degree of functionalization was determined using TGA from 80oC to 600oC under a nitrogen atmosphere. The weight loss for surface modified CNFs was 5.082%.

The weight loss of surface unmodified CNFs was 0% under the same

conditions. The degree of functionalization of (carbon/carboxyl functional group ratio) for the surface modified CNF was determined to be 70. The degree of functionalization is suggestive of one carboxyl functional group per every 70 carbon atoms on the CNF surfaces.

36

105 495.33°C 99.93%

100

‡ z

z ‡

z ‡

z

z

z

z

z

z

777.79°C

z z

‡ ‡ ‡

397.14°C 97.07%

95

‡ ‡

z

632.06°C

887.33°C 92.92% z

Weight (%)

‡

‡

90

z ‡

z

744.39°C 87.75%

85

‡

‡

CNF unmodified ––––––– z CNF modified – ‡– – –

80

‡

75

‡

0

200

400 600 Temperature (°C)

800

Figure 9. TGA thermograms of surface unmodified and modified CNFs in nitrogen

37

1000

120 565.91°C 98.65%

100

‡z

z ‡

‡z

z ‡

z ‡

z ‡

z ‡

z ‡

566.22°C 92.97%

80 Weight (%)

z ‡

z ‡

820.38°C z

769.03°C z

‡

862.24°C 58.10%

z

‡

60 z

‡

816.54°C 52.36%

40

‡

z

z

‡

20

‡

CNF unmodified ––––––– z CNF modified ––––––– ‡

0

-20

0

200

400 600 Temperature (°C)

‡

800

Figure 10. TGA thermograms of surface unmodified and modified CNFs in air

38

z z

1000

X-ray Photoelectron Spectroscopy X-ray photoelectron spectroscopy is extremely useful in studying the effects of chemically surface modified nanoparticles. This technique is based on the observation of electrons being emitted by atoms undergoing x-ray radiation. The energy of the emitted electrons yields the binding energy of the electron from a particular atom. The energies measured are characteristic of the elements and are sensitive to the electronic environment of the atom. This technique is very useful in quantifying the chemical species on the surface of nanoparticles. XPS was conducted to investigate the effect of surface modification on the CNFs. The surface unmodified and modified CNF atomic concentrations obtained from XPS are listed in Table 6.

Table 6. Atomic concentration of C1s and O1s obtained from XPS spectra Sample

C1s spectra

O1s spectra

Unmodified

98.223

1.777

Modified

86.696

13.304

An XPS scan of the surface modified CNFs is shown in Fig. 11. The oxygen 1s peak is broadened due to the presence of two types of oxygen bonded on to the CNF surface (12). The 1s oxygen peak typically is due to doubly bonded oxygen (C=O), and single bonded oxygen (C-O). The magnitude of the peak is proportional to the number of surface atoms of each type. The percentage of oxygen on the CNF surfaces increased by 7.5 times due to surface modifying the CNFs with 3:1 H2SO4/HNO3. The atomic oxygen on the surfaces of the CNFs is in the form of carboxylic acid (COOH) groups. Similar

39

studies by Glasgow (12) and Lakshminarayanan (39) use XPS to quantify the atomic amounts of oxygen and other atoms on the surface of CNFs. Their findings show that chemically modifying the CNFs with an oxidizing acid led to increased oxygen levels on the surfaces of the CNFs.

3000 2500

Counts

2000 1500 1000 500 0 541

536

531 526 Binding Energy (eV)

Figure 11. XPS scan of the oxygen 1s peak of surface modified CNF

40

521

Epoxy-Carbon Nanofiber/Polymer Characterization

Dynamic Mechanical Analysis Dynamic mechanical analysis (DMA) is used to determine the mechanical properties (i.e., modulus and damping of viscoelastic materials) over a spectrum of time (frequency) and temperature.

DMA can detect molecular motions enabling the

development property-structure-morphology relationships.

DMA was conducted to

investigate the effect of CNFs on polymer nanocomposite relaxation behavior. Plots of the storage modulus (E’) versus temperature for the 1% CNF compositions and dispersion methods are shown in Fig. 12. Table 7 shows the compositions of polymer nanocomposites and their corresponding codes. The glass transition temperature can be determined from the peak of the tan delta curves shown in Fig. 13 and the values are listed in Table 8. Plots of E’ versus temperature for the 0.1% CNF compositions and dispersion methods are shown in Fig. 14.

Table 7. List of samples and their corresponding codes Sample

Code

Neat

Neat

Mechanical Mixed Surface Unmodified

MMU

Mechanical Mixed Surface Modified

MMM

High Shear Mixed Surface Unmodified

HSU

High Shear Mixed Surface Modified

HSM

41

Starting with the lowest CNF concentration of 0.1 wt%, the curves show essentially no difference in the glassy state. However, as the temperature increases to the transition region, the curves for both the HSM and MMM extend to higher temperatures than the neat resin and unmodified sample, suggestive of higher glass transition temperatures (Tgs). The values of the plateau modulus are also increased. The Tg for the MMM 0.1% sample is higher than the HSM 0.1% sample as illustrated in the tan delta curves in Fig. 13 (see Table 8). When the CNF content is increased to 1%, more dramatic relaxation behavior is observed, as shown in Fig. 14. While no significant difference is observed in the glassy state, a marked difference in the slopes of the E’ decrease is observed in the transition region. In addition, the plateau modulus increases by a few tens of MPa. Both of these effects can be attributed to a reduction in molecular mobility, presumably caused by enhanced epoxy-CNF interactions. The tan delta curves shown in Fig. 15 provide additional insight into the relaxation behavior. Of particular note are the curves for the 1% modified samples. Both the MMM and HSM exhibit a first peak which is shifted to a higher temperature relative to the neat and unmodified CNF samples, followed by a second, higher temperature peak. The presence of the higher temperature peak in the modified samples relative to the unmodified samples clearly suggest that there is an interaction between the surface modified CNFs and the epoxy resin. It is also noted that the magnitude of the tan delta peaks for the MMM and HSM samples is lower than that for the neat epoxy and the MMU and HSU. The lower tan delta value is due to the more elastic nature of the resin caused by the presence of the CNFs and the enhanced epoxyCNF interactions.

42

10000

Storage Modulus (MPa)

z‡ 

z‡ 

z‡ 

z‡ 

1000

z‡ 

Neat ––––––– z HSM 0.1% – ‡– – – ––––– · HSU 0.1%  ‡ z 

‡ z 

‡ z 

‡

100

z 

‡ z



‡ 

z

10

1 20

40

60

80

100

120

z

140

z

z

z

10000

Storage Modulus (MPa)

z

z ‡

z ‡

z ‡



z ‡

1000

‡ 

‡ 

160

180

200

Universal V4.2E TA Instruments

Temperature (°C)

a)

Neat ––––––– z MMM 0.1% – ‡– – – ––––– · MMU 0.1%  ‡ z ‡ z  ‡ z  ‡

100

z  ‡ z  ‡ z 

10

1 20

b)

‡ 

‡ 

40

60

80

100

120

‡z

140

Temperature (°C)



‡z



160

‡z



‡z

180



‡

200

Universal V4.3A TA Instruments

Figure 12. Storage Modulus vs. Temperature curves of 0.1 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites

43

1 z ‡ ‡

 z

z

‡

 ‡

 z z

‡ z

0.1

Tan Delta

‡

z

‡ z ‡

z z

z ‡ z ‡

z‡ 

‡         ‡  ‡ ‡ ‡‡‡ ‡   ‡  ‡ ‡ ‡ ‡ ‡ ‡

z z z z z

 z‡

z zz z

z zz z

0.01

z z

zz

z

Neat ––––––– z HSM 0.1% – ‡– – – ––––– · HSU 0.1% 

0.001 20

40

60

80

100

120

140

160

a)

180

200

Universal V4.2E TA Instruments

Temperature (°C) 1 z z z z 

Tan Delta

0.1 z



‡

z  ‡

‡

 z ‡          ‡     z ‡ ‡ ‡‡‡ z z z z ‡‡ ‡ ‡ ‡z ‡ ‡ z ‡ z ‡‡ ‡‡ ‡ ‡‡ zz z z ‡ ‡ ‡ zz z ‡ ‡ z zz z

‡



z ‡

‡ z

z 

‡

‡

z ‡ z

‡

‡

 

0.01

z

0.001 20

b)

Neat ––––––– z MMM 0.1% – ‡– – – ––––– · MMU 0.1%  40

60

80

100

120

140

Temperature (°C)

160

180

200

Universal V4.3A TA Instruments

Figure 13. Tan delta vs. Temperature curves of the 0.1 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites

44

Table 8. Glass transition temperature from the tan delta peaks a) HSM, b) HSU, c) MMM, and d) MMU. (a) HSM CNF loading (%)

Peak 1 (oC)

Peak 2 (oC)

0

117

N/A

0.1

119

N/A

1.0

128

155

(b) HSU CNF loading (%)

Peak 1 (oC)

Peak 2 (oC)

0

117

N/A

0.1

117

N/A

1.0

126

N/A

CNF loading (%)

Peak 1 (oC)

Peak 2 (oC)

0

117

N/A

0.1

123

N/A

1.0

126

154

CNF loading (%)

Peak 1 (oC)

Peak 2 (oC)

0

117

N/A

0.1

115

N/A

1.0

126

N/A

(c) MMM

(d) MMU

45

10000

Storage Modulus (MPa)

z ‡ 

z ‡ 

z ‡ 

z‡ 

Neat ––––––– z HSM 1% – ‡– – – ––––– · HSU 1%  z ‡

1000

‡ z

 ‡ z

 ‡ z



100 z

‡ 

z

‡ ‡



‡ 

z

10

1 20

40

60

80

100

120

z

140

z

10000

Storage Modulus (MPa)

z ‡

z ‡

z ‡

z ‡

‡ z

1000

‡

‡ 

 z

z

160

180

200

Universal V4.2E TA Instruments

Temperature (°C)

a)

––––––– Neat z MMM 1% ‡– – – –  ––––– · MMU 1% ‡ z

‡

 z

‡

 z

100

‡

 z

‡

 ‡

z  z 

10

1 20

b)



40

60

80

100

120

 z

140

Temperature (°C)

‡

‡

‡

 z

 z

 z

160

180

200

Universal V4.3A TA Instruments

Figure 14. Storage modulus vs. Temperature curves of the 1.0 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites

46

1 z z

 ‡ z

 ‡

z

z



Tan Delta

‡

z z ‡ z 

z

z

‡

‡ ‡

  ‡     

z



‡

‡

‡

z

0.1

‡ 

z



z





‡ ‡

z z z

‡ z z





z zz z



‡ ‡‡ ‡

z zz z

0.01

z z

zz

z

0.001 20

Neat ––––––– z HSM 1% – ‡– – – ––––– · HSU 1%  40

60

80

100

120

140

160

180

a)

200

Universal V4.2E TA Instruments

Temperature (°C) 1 z

 

z

‡ z

‡

 

z

‡

‡

z

 z

Tan Delta

0.1

 ‡ z

 ‡ z

 z‡

 ‡ z

‡

‡

z

 z

‡

  

‡



‡

z z

0.01

‡ ‡ ‡          z z z  z z    z z z z  z zz z

  

z z

zz

z

0.001 20

b)

Neat ––––––– z MMM 1% – ‡– – – ––––– · MMU 1%  40

60

80

100

120

140

Temperature (°C)

160

180

200

Universal V4.3A TA Instruments

Figure 15. Tan delta vs. Temperature curves of the 1.0 wt% a) high shear mixed b) mechanically mixed polymer nanocomposites

47

Thermomechanical Analysis Thermomechanical analysis (TMA) is used to measure the dimensional change (i.e., CTE) in materials as a function of time, temperature and force in a controlled atmosphere. TMA was conducted to investigate the effect that the CNFs had on CTE of the polymer nanocomposites. The CTEs for the high shear samples were measured in two directions, parallel and perpendicular to the shear direction to investigate if the effect of alignment on CTE. The CNFs in the mechanically mixed samples are assumed to have no orientation preference and no specific emphasis was made to investigate any alignment effects. The CTE values of all of the PNC samples are listed in Table 9. The neat epoxy resin sample had a CTE value of 85 ppmC-1.

The 1 and 0.1% HSM

CNF/PNCs (parallel to the shear direction) samples showed the best property enhancements with CTEs of 67 and 59 ppmC-1 respectively. The 1 and 0.1% HSM CNF/PNCs (perpendicular to the shear direction) exhibited CTEs of 72 and 77 ppmC-1 respectively. The anisotropic CTE behavior observed suggests that the surface modified CNFs play a significant role in lowering the CTE of the epoxy resin matrix. The CTE values of the mechanically mixed samples fell in between the values of the high shear samples and the neat epoxy resin sample. The nominal property enhancement for the mechanically mixed samples is presumably due to poor dispersion and CNF aggregation (9). Due to the entangled and curved nature of CNFs, processing thermoset polymer nanocomposites using mechanical mixing may not be the best dispersal method in disrupting CNF aggregation. The effect of surface modified CNFs on the CTE of the CNF/PNCs is shown in Fig. 16. The samples made with surface modified CNFs have

48

better CTE properties relative to the samples made with surface unmodified CNFs, for both mixing methods.

Table 9. CTEs (ppm/oC) of mechanically mixed specimens (a) (1.0 & 0.1% wt/wt unmodified and modified CNF/PNCs) and (b) high shear (1.0 & 0.1% unmodified and modified CNF/PNCs) and neat epoxy resin. (a) Mechanical Mixed specimens CNF loading (%)

Unmodified

Modified

0

85

85

0.1

82 (-3)

71 (-14)

1.0

81 (-4)

71 (-14)

CNF loading (%)

Unmodified

Modified

0

85

85

0.1 (perpendicular)

66 (-19)

77 (-8)

1.0 (perpendicular)

75 (-10)

72 (-13)

0.1 (parallel)

65 (-20)

59 (-24)

1.0 (parallel)

69 (-16)

67 (-18)

(b) High Shear Mixed specimens

Chen et al. observed that adding large amounts (2-52% vol/vol) of CNFs reduced the CTE in their CNF/epoxy composites (15). The CNF/epoxy composites were fabricated by stacking plies of CNF mats followed by infusion of epoxy resin into the plies and compression molding. Chen and co-workers reported as CNF concentration increased from (2-52% vol/vol) that the CTE of epoxy CNF/PNCs decreased from 31.29 ppmK-1 to -0.11 ppmK-1 respectively (15). However, our 1 and 0.1 wt% HSM PNC samples

49

showed a 21 and 31% decrease respectively when compared to the Neat. The 1 and 0.1 wt% MMM PNC samples showed a 16.5% decrease when compared to the Neat sample. The CTE decrease in our surface modified CNF/PNC samples is very significant considering that we have a much lower volume fraction (0.54% vol/vol = 1 wt%) and (0.054% vol/vol = 0.1 wt%), which is 4 times less than the Chen et al. reported 2% vol/vol sample. Interestingly, the 1 and 0.1% surface modified CNF/Epoxy PNCs made by both methods show property enhancements relative to the neat epoxy resin sample indicating that only small amounts of modified CNFs are needed to alter the dimensional stability of the resultant polymer nanocomposites. Creating PNCs for high performance applications will need maximum synergistic efforts between the processing approach and surface modified nano-reinforcements.

50

1% 95 90 85

Neat MMU

CTE (ppm/oC)

80 75

HSU pd HSM pd

MMM

70

HSU pa

65

HSM pa

60 55 50 45 40 35 30

a)

0.1 % 100 95 90

CTE(ppm/oC)

85 80

Neat MMU HSM pd

75 70 65

MMM HSU pd

60

HSU pa HSM pa

55 50 45 40 35 30

b) Figure 16. CTE graphs of a) 1% samples b) 0.1% samples (pa = parallel), (pd = perpendicular)

51

Epoxy-CNF/PNC Processing Analysis Rheology is extremely useful in investigating the viscous (fluid), elastic (solid), and viscoelastic (behaves both as a fluid and solid) states of matter. This technique is used to determine the mechanical properties i.e., modulus and damping of fluid, solid and viscoelastic materials over a spectrum of time (frequency) and temperature, detect molecular motions and develop property-structure or morphology relationships. A schematic of the device used to disperse the CNFs was shown in Fig. 4 in the experimental section. Shearing through the syringe results in flow that can be described by Pouiselle’s equation: ΔP = μ L

8Q , where ΔP is the pressure drop, µ is the viscosity, π R4

Q is the flow rate, R is the radius, and L is the length. The shear rate for Newtonian .

fluids can be described by: γ =

4Q , where Q is the flow rate. For non-Newtonian π R3

fluids, the shear rate is an apparent shear rate (40). The shear rate can be determined by using the experimental volumetric flow rate. Based on the inner orifice diameter (id) 0.202 cm the flow rate of the neat epoxy resin is 11.73 mLs-1 and the calculated shear rate is 14495 s-1. The flow rates and apparent shear rates for the CNF/PNCs were calculated and are listed in Table 10. As the CNF loading increased from 0 to 1% the flow rates showed a decrease from 11.73 to 10.44 mLs-1. This change in flow rate was due to the CNFs disrupting flow through the orifice. The apparent shear rates for the CNF/PNC samples were lower than the shear rate of the neat epoxy resin, and can be attributed to the CNFs obstructing flow through the orifice. A flow study was conducted to correlate the rheological behavior of the fluid to the conditions experienced during flow. This was measured by applying a steady shear flow to the samples.

52

Table 10. Flow Rates and Shear Rates Sample

Flow Rate (mLs-1)

Shear Rate (s-1)

Neat

11.73

14495

0.1% unmodified CNF/PNC

10.68

13206

0.1% modified CNF/PNC

11.18

13813

1% unmodified CNF/PNC

10.50

12975

1% modified CNF/PNC

10.44

12907

Fig. 17 and Fig. 18 show the viscosity as a function of shear rate for the neat sample and CNF/PNC samples based on unmodified and modified CNFs. The viscosity of the neat epoxy resin is independent of frequency, consistent with behavior exhibited by a Newtonian fluid. As the nanoparticle content increases, the behavior changes to that of a complex fluid, with shear thinning behavior being exhibited at higher shear rates. The low frequency range is useful for discerning differences in the structure of the systems. As can be seen, the 1% unmodified and modified CNF/PNC samples exhibit high viscosities at low shear rates. This behavior has been observed for other carbon based nanoparticles dispersed in epoxy resins, including spherical particles and CNTs (41, 42). The effect of surface modification on the flow behavior is examined by overlaying the flow curves for the CNF/PNCs containing unmodified and modified CNFs, as shown in Fig. 19. Both systems exhibit higher viscosities than the neat epoxy resin.

53

viscosity (poise)

100.0

Unmodified 1% CNF/PNC Umodified 0.1% CNF/PNC Neat

10.00

1.000 0.1000

1.000

10.00

100.0

shear rate (1/sec)

Figure 17. Unmodified CNF/PNCs

viscosity (poise)

100.0

Modified 1% CNF/PNC Modified 0.1% CNF/PNC Neat

10.00

1.000 0.1000

1.000

10.00 shear rate (1/sec)

Figure 18. Modified CNF/PNC samples

54

100.0

viscosity (poise)

100.0

Modified 1% CNF/PNC Unmodified 1%CNF/PNC Modified 0.1% CNF/PNC Unmodified 0.1% CNF/PNC Neat

10.00

1.000 0.1000

1.000

10.00 shear rate (1/sec)

Figure 19. Overlay of All CNF/PNC samples

55

100.0

Moreover, the 1% modified and unmodified CNF/PNC samples exhibit approximately the same viscosity with respect to the other CNF/PNC samples. However, at high shear rates the 1% modified CNF/PNC sample has a higher viscosity, suggestive of a highly dispersed system and stronger epoxy resin-CNF interactions. On the other hand, the 1% unmodified CNF/PNC sample exhibited more distinct shear thinning behavior, suggestive that the CNFs are restricting resin flow at low shear rates, but at high shear rates the CNFs do not restrict resin flow leading to the resin exhibiting shear thinning behavior, indicative of CNF having poor dispersion. The behavior observed in this study has been observed in studies using MWNTs (41, 43). Abdalla et al. studied the effect of interfacial chemistry on molecular mobility and morphology of multi-walled carbon nanotubes epoxy nanocomposite (43).

The surfaces of the MWNT were functionalized with

fluorine and carboxylic acid groups. They report data similar to what is observed in this study. It is interesting to note that their 0.1 and 1% COOH-MWNT samples exhibited approximately the same viscosities at low shear rates. They show as the MWNT content increases, that the viscosity increases. In both modified and unmodified CNF/PNC samples in our study, similar behavior occurred. The studies by Kim et al. and Abdalla et al. attributed the increase in viscosity is due to the presence of strong interfacial interactions (41, 43). The high viscosity at low shear rates can be attributed to the formation of a percolated structure by the CNFs.

As the shear rate increases, the

percolated structure breaks down and results in all of the CNF/PNCs showing similar viscosities at high shear rates. In light of this, shearing followed by curing allows the well dispersed morphology to be locked in before aggregation can occur.

56

Morphology High-Resolution SEM was used to characterize the dispersion of the CNFs in the epoxy resin matrix.

Representative samples of the 1% unmodified and modified

CNF/PNCs are shown in Fig. 19. The Hi-Res SEM images in Fig. 20a and Fig. 20b shows fracture surfaces for the 1% unmodified CNF/PNC and 1% modified CNF/PNC samples, respectively. Circles are used to show some regions where CNFs are present. Fig. 20a shows that agglomerates are visible, indicative of a poor dispersion of the CNFs within the epoxy resin matrix. Fig. 20b shows that the some CNF agglomerates are present but the overall CNF dispersion is good in the epoxy resin matrix.

57

a)

b) Figure 20. Representative samples of a) 1% unmodified CNF/PNC b) 1% modified CNF/PNC

58

Flexural Testing (3-Point Bending) 3-point bending tests were conducted on the CNF/PNCs to investigate the effect of the CNFs on the flexural properties of the epoxy polymer matrix. A plot of flexural stress vs. strain for the neat, 0.1 and 1% CNF/PNCs are shown in Fig. 21. The flexural strength of the 0.1 and 1% CNF/PNC samples decreased by 33.75 and 14.43% respectively, relative to the neat sample. The flexural moduli of the 0.1% CNF/PNC sample decreased by 8.23% while the 1% CNF/PNC sample increased by 11.92%, relative to the neat sample. The decreased flexural strength is suggestive that possible flaws or defects are present in the CNF/PNC samples. The decreased flexural modulus in the 0.1% CNF/PNC sample is suggestive that the CNFs were not homogeneously dispersed, leading to inefficient load transfer within the CNF/PNC. The flexural modulus increase in the 1% CNF/PNC sample is suggestive that the CNF concentration is high enough to provide more efficient load transfer. Additional 3-point bending tests were conducted on 0.5, 2, and 4% CNF/PNC samples to observe if varying CNF concentrations would show mechanical property enhancements. A plot of stress vs. strain for neat, 0.1, 0.5, 1, 2, and 4 % CNF/PNC samples are shown in Fig. 22. The results of flexural testing of all CNF/PNC samples are listed in Table 11. All CNF/PNC samples show decreased flexural strength properties. There is a 12.01% increase in flexural modulus in the 0.5% CNF/PNC sample relative to the neat sample, followed by 11.92% and 5.81% increases in the 1% and 4% respectively. The 20.92% decrease in flexural modulus for the 2% CNF/PNC sample is suggestive of inhomogeneous CNF dispersion throughout the sample. The decreased flexural strength is suggestive that possible flaws or defects are present in the CNF/PNC samples.

59

Figure 21. Flexural stress vs. strain of neat, 0.1 and 1% CNF/PNC samples

60

Figure 22. Flexural stress vs. strain of neat, 0.1, 0.5, 1, 2, and 4 % CNF/PNC samples

61

A similar effect has been shown by others, and generally has been attributed to poor dispersion of the nanoparticles. Strength is very sensitive to flaws or defects within a sample, and the inhomogeneous nanoparticle dispersion can be considered as flaws. In addition, variations in crosslink topology can lead to molecular scale defects, such as dangling chains which would result in strength decreases (16, 30, 44). Flaws and defects can lead to inefficient load transfer within the CNF/PNC and subsequent premature PNC fracture.

Table 11. Flexural strength and flexural modulus mechanical properties Sample

Flexural Strength (MPa)

% Gain/Loss in Strength

Modulus (GPa)

% Gain/Loss in Modulus

Neat

118.5 (+/-10.7)

0

2.84

0

0.1%

78.5 (+/-9.2)

-33.75

2.60

-8.23

0.5%

87.9 (+/-10.8)

-25.82

3.18

12.01

1.0%

101.4 (+/-6.7)

-14.43

3.18

11.92

2.0%

62.7 (+/-16.2)

-47.08

2.23

-20.92

4.0%

37.8 (+/-3.1)

-68.01

3.00

5.81

62

Multiscale Fiber-Reinforced Composite Characterization

Flexural Testing (3 Point Bending) 3-Point bending tests were conducted on M-FRC to investigate the effect of the nanophased polymer matrix on the FRC. A plot of stress versus strain for the neat, 0.1 and 1% M-FRC samples are shown in Fig. 23. The results of flexural testing results of the neat, 0.1, 1% M-FRC samples are listed in Table 12.

Table 12. Flexural strength and flexural modulus mechanical properties Sample

Flexural Strength (MPa)

% Gain/Loss in Strength

Modulus (GPa)

% Gain/Loss in Modulus

Neat

336.8 (+/-6.5)

0

17.4 (+/-0.3)

0

0.1%

392.7 (+/-5.1)

16.61

21.05 (+/-1.1)

23.43

1.0%

404.2 (+/-18.6)

20.02

22.02 (+/-0.5)

26.44

The flexural strength of the 0.1 and 1% M-FRC samples increased by 16.61 and 20.02% respectively, relative to the neat FRC sample. The flexural modulus of the 0.1 and 1% M-FRC samples increased by 21.05 and 22.02% respectively, relative to the neat FRC sample. The increased mechanical properties are suggestive of enhanced nanophased polymer matrix-glass-fiber interfacial interactions.

Qui et al. (33) studied

epoxy/MWNT/glass-fiber M-FRC using the VARIM process.

They attributed the

observed mechanical property enhancements to enhanced polymer-nanoparticle-fiber interfacial interactions.

63

Figure 23. Flexural stress vs. strain of neat, 0.1 and 1% M-FRC samples

64

Fig. 24a and Fig. 24b show Hi-Res SEM micrographs of the fracture surfaces of 0.1 and 1% M-FRC samples. Both M-FRC samples show that the fibers are embedded within the matrix. It is obvious that the interfacial bonding between the nanophased polymer matrix and glass fiber was so strong that the nanophased polymer matrix deformed significantly during fracture.

The embedded glass-fiber remains in the

nanophased polymer matrix until the fiber broke creating a fracture surface. Since the breaking of the fiber consumed a significant amount of energy, the failure method achieved an efficient load transfer between the reinforcing constituent (glass-fiber) and the nanophased polymer matrix, which improved the mechanical properties of the MFRC samples. At higher magnification, the 1% M-FRC sample shown in Fig. 25a and Fig. 25b, show that CNFs are pulled out of the matrix during fracture. In Fig. 25a, many CNFs are visible in the nanophased polymer matrix.

Some CNFs experience fiber

pullout due to the M-FRC undergoing fracture and other CNFs remain embedded within the nanophased polymer matrix. In Fig. 25b, some CNFs experience fiber pullout and lay adjacent to the micron-sized glass-fibers. At high magnification, similar behavior was not observed for the 0.1% M-FRC sample. This was due to a combination of two limitations; the low concentration of CNFs and the brittle fracture behavior of the nanophased polymer matrix.

65

a)

b) Figure 24. Hi-Res SEM micrographs of a) 0.1% and b) 1% M-FRC samples

66

a)

b) Figure 25. Hi-Res SEM micrographs of 1% M-FRC samples a) 4000X b) 4500X

67

Interlaminar Shear Strength Testing Interlaminar shear strength (ILSS) testing was conducted on the M-FRC to understand the failure mechanism within the M-FRC samples. The high performance capability of glass-fiber FRCs is certainly related to the wettability of its surface by the matrix resin. A plot of shear strength vs. strain of the neat, 0.1 and 1% M-FRC samples is shown in Fig. 26. The results of interlaminar shear strength testing of all M-FRC samples are listed in Table 13.

Table 13. Shear strength and flexural mechanical properties Sample

Interlaminar Shear Strength (MPa)

% Gain/Loss in Strength

Neat

35.49 (+/-1.8)

0

0.1%

43.85 (+/-1.0)

23.55

1.0%

38.48 (+/-2.5)

8.42

The 0.1% M-FRC sample shear strength showed a 23.55% increase relative to the neat FRC sample. Interestingly, the 1% M-FRC sample shear strength showed a nominal increase 8.42% relative to the neat FRC sample. The small increase in the 1% M-FRC is suggestive of an inhomogeneous dispersion of the CNFs within the nanophased polymer matrix and/or inhomogeneous wetting of the glass-fibers with the nanophased polymer matrix. This is caused by the CNFs being filtered by the glass-fiber preforms during MFRC fabrication using the VARIM process. The cross-sectional fracture surfaces of the neat, 0.1 and 1% M-FRC samples representative optical micrographs are shown in Fig. 27a, Fig. 27b, and Fig. 27c, respectively.

68

Figure 26. Shear strength vs. strain of neat, 0.1 and 1% M-FRC samples

69

a)

b)

c) Figure 27. Cross-sectional fracture surface optical micrographs of a) neat, b) 0.1% c) and 1% M-FRC samples 70

All of the M-FRC samples fail in shear. In Fig. 27a, the neat FRC shows that the crack begins at the matrix-glass-fiber interface and propagates through the matrix and is blunted by the glass-fiber. In Fig. 27b, the 0.1% M-FRC sample shows that the crack propagates along the nanophased polymer matrix-glass-fiber interface. In Fig. 27c, the 1% M-FRC sample shows that the crack propagates into the cross-section of glass-fiber tow.

The increase in ILSS is suggestive that the surface modified CNFs do not

compromise the epoxy resin ability to effectively wet the glass-fibers. Dean et al. (30) Haque et al. (31) and Kornmann et al. (32) used LS as the nanoconstituent in their nanophased matrix and glass-fibers as the reinforcement component to fabricate MFRCs. They observed substantial improvements in ILSS, and concluded that the ILSS improvements were due to efficient load transfer between the nanophased polymer matrix and the reinforcing component.

Dynamic Mechanical Analysis Dynamic mechanical analysis was conducted to investigate the effect of the nanophased polymer matrix on the M-FRC relaxation behavior.

A plot of storage

modulus (E’) vs. temperature and tan delta vs. temperature is shown in Fig. 28a and Fig. 28b respectively. The results from dynamic mechanical analysis of all M-FRC samples are listed in Table 14.

71

G' (Pa)

1.000E10

0.1% FRC sample 1% FRC sample Neat

1.000E9

1.000E8 25.0

50.0

75.0

a)

100.0 125.0 temperature (°C)

150.0

tan(delta)

1.000

175.0

200.0

0.1% FRC sample 1% FRC sample Neat

0.1000

0.01000 25.0

b)

50.0

75.0

100.0 125.0 temperature (°C)

150.0

175.0

200.0

Figure 28. a) Storage modulus vs. temperature curves and b) tan delta vs. temperature curves of neat, 0.1, and 1% M-FRC samples

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Table 14. Storage Modulus and Glass Transition Temperature Sample

Storage Modulus in glassy region (GPa)

Glass Transition Temperature (oC)

% Tg Improvement

Neat

4.24

78

0

0.1%

4.97

103

32.05

1.0%

4.43

103

32.05

The 0.1 and 1% M-FRC sample showed an increase in E’ in the glassy region of 17 and 4.5% respectively, relative to the neat FRC sample. The increase in E’ in the M-FRC samples is suggestive that M-FRC samples have a well dispersed nanophased polymer matrix and enhanced interfacial bonding. However, as the temperature increases to the transition region, the curve for both 0.1 and 1% extend to higher temperatures than the neat sample, suggestive of higher Tgs. The Tgs were obtained from the maximum of the tan delta peak in Fig 28b. Interestingly, the Tgs of 0.1 and 1% M-FRC samples increased by 25 oC, relative to the neat FRC sample. However, when the CNF content increased from 0.1 to 1%, no significant increase in Tg was observed. The Tg increase observed in the M-FRC samples relative to the neat FRC sample is suggestive of an increase in crosslink density, mechanical reinforcement arising from a percolated CNF morphology, and/or restriction of segmental relaxation of chain segments near the CNFs. Enhancing the nanophased polymer matrix relaxation behavior allows for enhancing the M-FRC relaxation behavior.

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Thermomechanical Analysis Thermomechanical analysis was conducted to investigate the effect of the nanophased polymer matrix on the CTE of M-FRC. The CTE values of all FRC samples are listed in Table 15.

Table 15. CTEs (ppmC-1) of the neat, 0.1 & 1% samples Sample

CTE value (ppm/oC)

% Reduction

Neat

71

0

0.1%

44

37

1.0%

33

53

The neat FRC sample had a CTE value of 71 ppmC-1. The 0.1 and 1% M-FRC samples show CTE values of 44 and 33 ppmC-1 respectively. In light of this, the 0.1 and 1% MFRC samples show a 37 and 53% decrease respectively, relative to the neat FRC sample. Interestingly, both M-FRC samples show that substantial property enhancements relative to the neat sample are obtainable with small amounts of surface modified CNFs. This suggests that small amounts of CNFs significantly alter the dimensional stability of the nanophased polymer matrix, and as a result will significantly affect the dimensional stability of M-FRC.

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CHAPTER 5

CONCLUSIONS

We have studied the processing-property relationships of M-FRCs with a hierarchal microstructure ranging from nanoscale fibers to micron size fibers have been studied. Surface modification of the CNFs promoted dispersion of the CNFs in the epoxy resin and led to enhanced relaxation behavior (Tg) and thermal dimensional stability (CTE) in the CNF/PNC nanophased matrix.

Enhancing the CNF/PNC nanophased

matrix properties allows for the development of M-FRCs with improved properties. Addition of the CNF in the epoxy resin results in a nominal increase in resin viscosity, and is within the desired range for VARIM processing. Multiscale fiber-reinforced composites were prepared using a combination of VARIM and compression molding. The Tgs of the M-FRC samples increased (25oC) for both 0.1 and 1% M-FRC samples respective to the neat FRC sample. The CTE properties improved (37-53%) for both M-FRC samples respective to the neat FRC sample. The resulting 0.1 and 1% MFRCs exhibited a maximum flexural strength and modulus enhancement (16-20%) and (23-26%) respectively. The 0.1% M-FRC sample showed a 23% increase in ILSS, while the 1% M-FRC showed a nominal increase in ILSS (8%). The small increase in ILSS for the 1% M-FRC sample is due to an inhomogeneous dispersion of the CNF/PNC nanophased matrix within the M-FRC caused by filtering of the CNFs by the glass-fiber

75

preforms during VARIM processing of M-FRC. The enhanced relaxation behavior, thermal dimensional stability and mechanical properties in the 0.1 and 1% M-FRCs may be due to synergistic interfacial interactions between the CNF/PNC nanophased matrix and glass-fibers. The M-FRCs thermomechanical and mechanical property enhancements attained with very low filler loaded-nanophased matrices coupled with the ease of production makes them a very promising new class of materials. The unique combination of their key properties and potential low production costs paves the way for their usage in an array of potential applications. The very low filler amount needed to result in property enhancements make them competitive with other materials.

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