EFFECTS OF NANOPARTICLE AND MATRIX INTERFACE ON
NANOCOMPOSITE PROPERTIES
A Dissertation
Presented to
The Graduate Faculty of The University of Akron
In Partial Fulfillment
of the Requirements for the Degree
Doctor of Philosophy
Sandi G. Miller
August, 2008
EFFECTS OF NANOPARTICLE AND MATRIX INTERFACE ON
NANOCOMPOSITE PROPERTIES
Sandi G. Miller
Dissertation
Approved: Accepted:
Advisor Department Chair Dr. Darrell H. Reneker Dr. Mark D. Foster
Committee Member Dean of the College Dr. Ali Dhinojwala Dr. Stephen Z.D. Cheng
Committee Member Dean of the Graduate School Dr. Gary R. Hamed Dr. George R. Newkome
Committee Member Date Dr. Michael A. Meador
Committee Member Dr. Sadhan C. Jana
ii
ABSTRACT
The objectives of this work were to functionalize two nanoparticles, layered silicate
clay and expanded graphite, and evaluate the effects of surface modification on polymer
nanocomposite properties. Two thermosetting resin systems were evaluated, a polyimide
for high temperature applications, and a general use epoxy. The chemistry of the
modifier or the particle surface was tailored in each case to optimize nanocomposite
properties such as: particle dispersion, thermal oxidative stability (TOS), electrical
conductivity, strength, and toughness.
Dispersion of layered silicate clay into the two separate matrices demonstrated an apparent affinity between the silicate surface and aromatic compounds. Steps were taken in each case to disrupt that attraction; resulting in improved material properties. The dispersion of layered silicate clays into a thermosetting polyimide demonstrated that improved thermal oxidative stability was achieved only when the clay was modified with a combination of an aromatic diamine and an alkyl ammonium ion. When such a system was employed, the nanocomposite TOS improved by 25% over that of the base polyimide. Attention to the interactions between clay and aromatic containing compounds was also necessary for silicate modification and dispersion in an epoxy blend.
Here, preferential contact between the clay and the aromatic containing sections of the blend was observed; resulting in nanocomposites exhibiting little enhancement to epoxy properties. By forcing the clay into the non-aromatic component, the material yield stress
iii increased by up to 65%, Young’s modulus increased by up to 80%, and increases in Tg of up to 11oC were observed relative to the base resin.
Within nano-graphite containing materials, trade-offs in functionalization, dispersion, and properties were evaluated. Functionalization of graphite proved beneficial in terms of dispersion. For example, an epoxy functionalized graphite nanoparticle resulted in acceptable dispersion throughout the matrix, with a minimal level disruption of the sp2 hybridization within the graphene sheet. As a result, the nanocomposite structure increased yield stress by 30% at a filler loading of 0.67 vol%. Electrical conductivity increased by 5 orders of magnitude at this same loading. Graphite materials that did not disperse well, or were more heavily oxidized exhibited conductivity as loading increased to 1.5 vol%. In the poorly dispersed expanded graphite material, a 35% increase in yield stress was observed, but with significant reduction in ductility. With the heavily oxidized graphene sheets, 50% increase in yield stress was observed, following adjustments to the resin stoichometry.
iv
ACKNOWLEDGEMENTS
First and foremost, I wish to thank Dr. Michael Meador for the guidance, encouragement, kindness, and support that he provided as a mentor. I would also to sincerely thank Professor Darrell Reneker for taking the time to work with me through the later stages of my graduate career at the University of Akron. I would also like to extend that gratitude to all my committee members, including Professor Ali Dhinojwala,
Professor Gary Hamed, and Professor Sadhan Jana. I sincerely appreciate their time, effort, and critical evaluation of my work.
This research for this project was conducted at the NASA Glenn Research Center.
As such, there are a number of people who have helped me over the years. The entire
Polymers Branch, both past and present, has been an incredible source of knowledge and guidance. In particular, I would like to thank Daniel Scheiman, Linda Inghram, Linda
McCorkle, Gary Roberts, Justin Littell, Lee Kohlman, and Paula Heimann. Additionally,
I would like to acknowledge the many researchers outside of the Polymers Branch who have been invaluable resources for assistance with material characterization.
Specifically, Dave Hull, Joy Buelher, Joe Lavelle, Dorothy Lukco, Rick Rogers, Bob
Mattingly, and Anna Palczer.
I would also like to thank the friends who have encouraged me over the many, many years it has taken to complete this degree, and my parents and brother for their un-ending
v support. Finally, I would not have been able to complete this work without the patience, love, and support I received from my husband Kyle and daughters, Kennedy and Ashley.
vi
TABLE OF CONTENTS
Page
LIST OF TABLES…………………………………………………………………….…xii
LIST OF FIGURES……………………………………………………………………..xiv
CHAPTER
I. INTRODUCTION…………………………………………………………………….1
1.1 Polymer Nanocomposites…………………………………………………6
1.2 Layered Silicate Clay Nanocomposites…………………………………...8
1.2.1 High Temperature Polyimide Matrix……………………………...8
1.2.2 Epoxy Blend Matrix……………………………………………….9
1.3 Graphite Nanocomposites………………………………………………..10
1.3.1 Epoxy Blend Matrix……………………………………………...10
II. OBJECTIVES………………………………………………………………………..12
2.1 PMR-15 Matrix Nanocomposites………………………………………..12
2.2 Epoxy Resin Matrix Nanocomposites…………………………………...14
2.3 Summary…………………………………………………………………15
III. BACKGROUND……………………………………………………………………17
3.1 Conventional Fillers……………………………………………………...17
3.2 Comparison of Conventional and Nano Fillers………………………….18
vii 3.3 Layered Silicate Clays…………………………………………………...23
3.4 Organic Modification of Layered Silicate Clays………………………...25
3.5 Dispersion of Layered Silicate Clays…………………………………….30
3.6 Dispersion Mechanism…………………………………………………...35
3.7 Polymer-Clay Interface…………………………………………………..38
3.8 Nanocomposite Properties……………………………………………….39
3.9 Summary of Layered Silicate Clay Nanocomposites……………………40
3.10 Expanded Graphite……………………………………………………….41
3.11 Dispersion of Expanded Graphite………………………………………..45
3.12 Dispersion Techniques…………………………………………………...46
3.13 Polymer-Graphite Interface……………………………………………...48
3.14 Nanocomposite Properties……………………………………………….49
3.15 Conclusion of Expanded Graphite Nanocomposites…………………….52
3.16 Summary…………………………………………………………………52
IV. EFFECTS OF SILICATE-MODIFIER ORIENTATION ON NANOCOMPOSITE INTERFACE AND PROPERTIES IN A HIGH TEMPERATURE POLYIMIDE MATRIX…………………………………………………………………………….53
4.1 Introduction………………………………………………………………53
4.2 Experimental Details……………………………………………………..56
4.2.1 Materials…………………………………………………………56
4.2.2 Synthetic Procedures……………………………………………..57
4.2.3 Nanocomposite Synthesis………………………………………..57
4.2.4 Characterization………………………………………………….59
4.3 Results and Discussion…………………………………………………..61
viii 4.3.1 Ion Exchanged Clays…………………………………………….61
4.3.2 Orientation Characterization by Infrared Spectroscopy…………66
4.3.3 Characterization of Silicate Dispersion in PMR-15……………...68
4.3.4 Processing PMR-15 Nanocomposites……………………………71
4.3.5 Thermal Stability of PMR-15 Nanocomposites………………….76
4.3.6 Polymer Matrix Composites……………………………………..82
4.4 Conclusions………………………………………………………………86
V. EFFECT OF INTERFACE ON SILICATE DISPERSION IN A CO-EPOXY SYSTEM………………………………………………………………………….…89
5.1 Background………………………………………………………………89
5.2 Introduction………………………………………………………………90
5.3 Experimental……………………………………………………………..92
5.3.1 Materials…………………………………………………………93
5.3.2 Nanocomposite Preparation……………………………………...93
5.3.3 Characterization………………………………………………….95
5.4 Results……………………………………………………………………96
5.4.1 Characterization of Silicate Dispersion………………………….96
5.4.2 Interaction Between Cloisite 30B and Aromatic or Aliphatic Compounds………………………………………………………99
5.4.3 Glass Transition Temperature…………………………..………103
5.4.4 Tensile Tests…………………………………………………....105
5.4.4.1 General Trends Observed From Tensile Testing…….…111
5.4.4.2 Relation of Matrix Mobility…………………………….111
5.4.5 Permeability…………………………………………………….115
ix 5.4.6 Coefficient of Thermal Expansion……………………………...115
5.5 Discussion………………………………………………………………117
5.5.1 Microstructure…………………………………………………..117
5.5.2 Preferential Interaction Between Silicate Clay and Aromatic Compounds……………………………………………………..121
5.5.3 Pre-Swelling Clay with the Aliphatic Component……………...121
5.5.4 Interpretation of Results………………………………………...122
5.5.4.1 Glass Transition Temperature…………………………..122
5.5.4.2 Mechanical Properties…………………………………..123
5.6 Conclusions……………………………………………………………..123
VI. FUNCTIONALIZED GRAPHITE NANOCOMPOSITES……………………….125
6.1 Background……………………………………………………………..125
6.2 Introduction……………………………………………………………..125
6.3 Experimental……………………………………………………………127
6.3.1 Materials………………………………………………………..127
6.3.2 ATI Preparation………………………………………………...127
6.3.3 Nanocomposite Preparation………………………………….…128
6.3.4 Organic Modification of Graphite……………………………...128
6.3.5 Characterization………………………………………………...129
6.4 Results…………………………………………………………………..130
6.4.1 Characterization of Graphite Materials…………………………130
6.4.1.1 Expanded Graphite Characterization…………………...132
6.4.1.2 Characterization of Adherent Technologies Graphite.…135
x 6.4.1.3 Characterization of Exfoliated Graphite (FGS)………...140
6.4.2 TGA Determination of Graphite Functionalization…………….143
6.4.3 Dispersion in an Epoxy Matrix…………………………………146
6.4.4 Physical Property Analysis………………………………….….146
6.4.4.1 Glass Transition Temperature…………………………..150
6.4.5 Electrical Conductivity…………………………………………150
6.4.6 Mechanical Properties…………………………………………..157
6.5 Discussion……………………………………………………………....166
6.6 Conclusions……………………………………………………………..168
VII. CONCLUSIONS AND FUTURE WORK………………………………………...174
7.1 Layered Silicate Clay Nanocomposites………………………………...174
7.2 Graphite Nanocomposites………………………………………………176
7.3 Future Work………………………………………………………….…177
REFERENCES………………………………………………………………………....179
xi
LIST OF TABLES
Table Page
3.1 Literature data comparing physical properties of Nylon 6, and Nylon 6/ clay composites; where clay is dispersed by varying degrees…………21
4.1. TGA data indicating the quantity of diamine present in clays…………………...63
4.2 Increase in silicate gallery height after ion exchange of aromatic diamines in SCP and PGV…………………………………………………………………65
4.3 FTIR characterization of ion exchanged clays…………………………………...70
4.4 Squeeze flow index and crosslinking enthalpy of neat PMR-15 and PMR-15 nanocomposites…………………………………………………………………..75
4.5 Squeeze flow index and crosslinking enthalpy of neat PMR-15 and PMR-15 nanocomposites…………………………………………………………………..83
4.6 Tg of post cured PMR-15 and PMR15 nanocomposites…………………………85
5.1 Tg measurements by DSC………………………………………………………104
5.2 Yield stress as determined from tensile tests……………...……………………106
5.3 Strain to failure as determined from tensile tests……………………………….107
5.4 Young’s modulus as determined from tensile tests…………………………….108
5.5 Toughness as determined from tensile test………………………………….….109
5.6 Change in saturation stress, Young’s modulus, and strain to failure on silicate dispersion……………………………………………………………….112
5.7 Oxygen permeability of epoxy blend based resins and nanocomposites……….114
5.8 Values for the coefficient of thermal expansion…………………...…………...116
xii 6.1 Summary of graphite nanoparticle characterization……………………………131
6.2 Calculated volume percent of graphite used for each corresponding weight percent…………………………………………………………………………..148
6.3 Measured viscosity on increased mixing time…………...……………………..149
6.4 Glass transition temperature as measured by DSC…………………...………...151
6.5 Electrical resistivity of graphite-epoxy nanocomposites……………………….152
6.6 Results of EG nanocomposite tensile tests……………………………………..155
6.7 Results of ATI tensile tests, comparing all nanocomposites to the stoichiometric neat resin………………………………………………………..159
6.8 Results of ATI tensile tests, comparing all nanocomposites to the corresponding neat resin………………………………………………………..163
6.9 Results of FGS tensile tests, comparing all nanocomposites to the stoichiometric neat resin………………………………………………………..164
6.10 Change in nanocomposite mechanical properties relative to stoichiometric base resin………………………………………………………………………..165
6.11 Tensile test data of nanocomposites and neat resins prepared with excess amine curing agent……………………………………………………………...167
xiii
LIST OF FIGURES
Figure Page
1.1 Schematic of carbon fiber reinforced composite, illustrating a 0-90o weave……..3
1.2 Schematic of a high bypass turbofan engine temperature profile…………………4
1.3 Components of Ares V cargo launch vehicle…...………………………………...5
1.4 Chemical structure or image of commonly investigated nanoparticles…………...7
3.1 Depiction of ‘tortuous path’ incurred by a permeant traveling through polymer filler with nanoclay at different orientations and degrees of dispersion………………………………………………………………………...20
3.2 Chemical structure of montmorillonite clay………..……………………………24
3.3 Arrangement of modifiers between layers as determined by X-ray diffraction and Fourier Transform Infrared Spectroscopy……………………….27
3.4 Potential polymer-clay morphologies……………………………………………31
3.5 Structure of layered natural flake graphite……………………………………….42
3.6 Expansion and dispersion of natural flake graphite……………………………...44
4.1 Synthesis of PMR-15………………………………………..…………………...55
4.2 Chemical structure of ion exchange diamines ..…………………………………58
4.3 4.3 (a) XRD patterns of unmodified and organically modified SCP, and (b) XRD patterns of unmodified and organically modified PGV…………...64
4.4 Figure 4.4 (a) orientation of protonated aromatic diamine in SCP; low cation exchange capacity clay, and (b) orientation of protonated ABDM in PGV; high cation exchange capacity clay…………………………………….67
4.5 FTIR spectra of SCP-BAX and PGV-BAX……………………………………...69
xiv 4.6 TEM images of PGV-MDA dispersed in PMR-15………………………………72
4.7 Dispersion of PGV-BAX in PMR-15 resin, and dispersion of SCP-BAX in PMR-15 resin……………………………………………………...73
4.8 Weight loss of neat PMR-15 and PMR-15 nanocomposites on heating at 288oC for 1000 hours………………………………………………………….78
4.9 Molecular length of methylene dianiline, dodecylamine and octylamine (from optimized structures)………………………………………………………79
4.10 XRD patterns of organically modified PGV clay………………………………..80
4.11 TEM images of PGV(MDA-C12) dispersed in PMR-15………………………..81
4.12 Weight loss of neat PMR-15 and PMR-15 nanocomposites on heating at 288oC for 1000 hours, where silicate is modified with a diamine and an aliphatic amine…………………………………………………………………...84
4.13 Weight loss of PMR-15 matrix and PMR-15 nanocomposite matrix carbon fiber composites on aging at 288oC for 1000 hours, where the PGV modifications are ABDM and 50:50 MDA/C12…………………………………88
5.1 Chemical structures of Epon 826, Araldite DY3601, Jeffamine D230, and the organic modifier on Cloisite 30B………………………………………..94
5.2 XRD patterns of 90:10, 70:30, and 50:50 nanocomposites, respectively………..97
5.3 TEM images of 2 wt% and 5 wt% 30B, respectively, in 70:30 systems, and illustrates dispersion of 2 wt% 30B in 90:10 resin………………………………98
5.4 Cloisite 30B aggregates in propylene glycol (left) and dispersed in chlorobenzene (right)………………..………………………………………….100
5.5 XRD pattern of Cloisite 30B and 2% 30B pre-swollen in DY3601…………....101
5.6 TEM images of 2% 30B in 70:30 epoxy matrix prepared by (a) simple mixing, and (b) pre-swelling the clay layers……………………………………..……...102
5.7 Stress – strain curves of 90:10 and 70:30 series, respectively…………….……110
5.8 SEM image of the fracture surface of the 70:30 resin containing 5 wt% 30B…119
5.9 SEM images of fracture surface of 50:50 resin containing 5 wt% 30B………...120
xv 6.1 XRD patterns of (a) acid intercalated natural flake graphite, and (b) expanded graphite…………………………………………………………………………133
6.2 SEM image of expanded graphite illustrating regions of closely spaced graphite layers, and well separated graphite layers…………………………….134
6.3 SEM images illustrate decrease in EG particle size on increased sonication time……………………………………………………………………………..136
6.4 XRD pattern showing decreased peak intensity as EG aggregate size is reduced………………………………………………………………………….137
6.5 XRD pattern of ATI graphite showing intense peak corresponding to graphite layer stacking….……………………………………………………....138
6.6 SEM images of untreated graphite and graphite covalently bonded with epoxy monomer……………………………………………………………...... 139
6.7 XRD pattern of FGS showing complete separation of the graphene layers……141
6.8 SEM image of FGS……………………………………………………………..142
6.9 TGA plot of as received ATI and organically modified ATI…………………..144
6.10 TGA plot of as received FGS and organically modified FGS………………….145
6.11 TEM images of EG, ATI, and FGS, respectively, in an epoxy matrix…………147
6.12 SEM images of 3 wt% EG in epoxy and 1.0 wt% FGS of epoxy………………154
6.13 Stress-strain curves of EG nanocomposites…………………………………….156
6.14 Stress-strain curves of ATI nanocomposites in comparison to the stoichiometric neat resin………………………………………………………..160
6.15 Fracture surface of EG, 0.5% EG and 3.0% EG nanocomposites……………...169
6.16 Fracture surface of ATI nanocomposites containing 0.5 wt% graphite and either 10% excess amine (top), or 0% excess amine (bottom)…...... …….170
6.17 Fracture surface of FGS nanocomposites containing either 0% excess amine (top), or 10% excess amine (bottom)……………………………………171
xvi CHAPTER I
INTRODUCTION
The wide spectrum of properties available to polymeric materials has afforded numerous practical applications ranging from common household goods to biomedical materials and aerospace components. Often, however, the inherent properties of the polymer alone are insufficient to meet the structural demands of an application. As a result, blending with a stronger or stiffer material is necessary to improve upon the mechanical performance of the pristine polymer.
A composite structure is a combination of two or more different components resulting
in a material having better performance than each individual constituent.1 Such
constructions impart mechanical strength and stiffness to a polymeric matrix. Composite
structures have been widely applied to areas including, but not limited to, the automotive,
aerospace and defense industries.
In the aerospace community, specifically, there has been a steady growth of applications to which polymer matrix composites (PMCs) are well suited.2 PMC’s are lightweight compared to the traditionally used metallic materials, thereby enabling airplanes, missiles, and spacecraft to operate with less fuel or increased payload.3 The
majority of PMC’s used in aerospace and aeronautics applications utilize a high
performance resin matrix with carbon fiber reinforcement.4 Figure 1.1 depicts the
1 components of a carbon fiber reinforced polymer. Typically, the carbon fiber is present
in 50 to 60 wt% of the overall composite, with the resin comprising the remaining 40 to
50 wt%.
The primary material requirements of high performance, aerospace resins are thermal
stability and mechanical performance.5 Increasing the thermal oxidative stability of a
PMC allows incorporation of these materials into applications requiring higher operating temperatures, such as those experienced in compressor components of aircraft engines.
The temperature profile of a high bypass turbofan engine is illustrated in Figure 1.2.6
Primarily, PMC’s are suited for engine components such as ducts, feedlines, and fan casings,4-5,7 PMC’s are also a candidate material for space applications; for example
NASA has evaluated many designs for composite cryogen propellant tanks.8-11
Additionally, it has been proposed that components of the new Ares V Launch Vehicle, including the interstage connector, Figure 1.3, could be made of a PMC.12
As new technologies continue to place increasingly stringent demands on the
performance of polymeric materials, it is becoming clear that traditional polymer matrix
composites can not meet these requirements. As a result, new composite materials are
being developed, utilizing the multifunctional nature of several nano-scale materials.13-15
Multifunctional composites are currently attracting considerable interest in academia, government, and industry. Such materials provide structural integrity, as well as serve additional functions, for example conduct electricity or impart enhanced gas barrier performance.
2
Matrix (grey)
Carbon Fiber (black)
Figure 1.1. Schematic of carbon fiber reinforced composite, illustrating a 0-90o weave
3
Figure 1.2. Schematic of a high bypass turbofan engine temperature profile
4
Figure 1.3. Components of Ares V Cargo Launch Vehicle
5 1.1 Polymer Nanocomposites
The term “polymer nanocomposite” broadly describes any number of multicomponent
systems, where the primary component is the polymer and the filler material has at least
one dimension below 100nm.16 Polymer nanocomposites are generally lightweight, require low filler loading, are often easy to process, and provide property enhancements extending orders of magnitude beyond those realized with traditional composites.
Rapid advancements in nanocomposite technologies have been realized as new classes of nanoscale fillers continue to emerge. Nanoparticles commonly dispersed in polymer matrices include: Polyoctahedral silsesquioxane (POSS),17-19 layered silicate clays,20-22 carbon nanofibers,23-25 carbon nanotubes,26-28 and graphite nanoflakes.29-31 The structure of these nanoparticles are shown in Figure 1.4. The nano- particles listed above differ in chemistry, morphology, aspect ratio and aggregate size.
The nanoparticle chosen for dispersion in a resin is dependent on the intended application. However, realization of significant enhancement in properties with any of
the nano-materials requires that the nano-particle is well dispersed throughout the matrix.
Furthermore, it is desirable that the particle bond with the matrix.
Generally, composite components should interact either chemically or physically, with
the matrix to create a strong interface. In all cases, the interface between the composite
components plays a defining role in the overall material properties.32 This becomes
critical as the available interfacial area is increased by the dispersion of nano-sized
particles throughout the polymer matrix.
6
Figure 1.4. Chemical structure or image of commonly investigated nanoparticles
7 Creating a strong bond with the matrix requires an understanding of the interactions
taking place at the nano-particle/matrix interface. In this work we examine two matrices;
a high temperature polyimide resin, PMR-15, which is commercially used in aircraft
engine components, and a toughened epoxy system as a model resin for low temperature
space applications. Regardless of the resin, the common goal is optimization of the
interface to realize the full potential of the nanoparticles within a composite structure.
Both layered silicate clays and expanded graphite nanoparticles are chosen for the
purpose of this work. They are similar in that both are composed of stacked nano-sized
platelets, however differ significantly in processing requirements and surface chemistry.
Working with these materials for high performance applications introduces issues that
have not received significant attention in the literature, such as the thermal oxidative
stability of the particle surface modifier, or extensive functionalization of graphite
particles.
1.2 Layered Silicate Clay Nanocomposites
The layered silicate clay work has been separated into two sections with include
dispersion in a high temperature polyimide matrix, and dispersion in an epoxy matrix.
1.2.1 High Temperature Polyimide Matrix
The majority of past work in layered silicate clay nanocomposites has focused on
thermoplastic or epoxy matrices, which are processed at relatively low temperatures,
allowing the clay to be modified with a long chain alkyl ammonium ion.33-37 However, on dispersion into a high temperature resin, such a modifier degrades at the high curing or post curing temperatures required for polymer processing. There has been research into
8 the development of thermally stable modifiers, however little attention has been focused
on how that modifier affects the viscosity or cure of a thermosetting resin. However,
several groups have evaluated successful alternatives to the traditional clay modification.
Delozier, et al.38 prepared reduced charge organoclays. This process reduces the
number of intragallery cations available for exchange with an organic ammonium salt.
As a result, the organic content of the modified clay is reduced and the effects of surfactant degradation during polymer processing are diminished.
Liang et al.39 chose thermally stable, rigid rod aromatic modifiers for use in polyimide
nanocomposites. The results were positive in that, below 3 wt% clay content, the nano-
structure strengthened the neat resin. The thermally stable silicate also led to a 25oC increase in the temperature of 5% weight loss on heating, compared to the pristine polyimide.
This work also focuses on the development of a thermally stable nanoclay for dispersion in high temperature resins. This is accomplished by:
(1) Evaluating a series of modifiers for use in a high temperature resin matrix, PMR-15.
(2) Understanding the orientation of the modifier within the clay galleries.
(3) Understanding the effect of the amine on the resin melt viscosity.
(4) Evaluating the thermal stability of the resin nanocomposite and carbon fiber reinforced nanocomposites.
1.2.2 Epoxy Blend Matrix
Epoxy resins have been widely utilized as a matrix resin for silicate clay nanocomposites.40-45 These resins offer affordability, commercial availability, and ease of
processing. With little exception the epoxies used for nanocomposite studies have been
9 single component epoxy systems. However, the high performance resins required for
aerospace applications often utilize a combination of epoxy resin and rubber toughener.
Synergistic effects between toughened epoxy and a nanocomposite structure have been
reported,46 and are likely due to the generally improved mobility of the clay. For this
work, we attempt to answer the following questions with regard to a blend of aromatic
and aliphatic based epoxy resins, and a clay nanoparticle.
(1) Will the clay prefer one region of the blend, e.g. aromatic over aliphatic?
(2) How will properties such as strength, ductility, and permeability be affected by a
nanocomposite blend?
(3) Can the preference of the clay be overcome by appropriate treatment of the silicate
surface?
1.3 Graphite Nanocomposites
The graphite nanocomposite work focuses only on an epoxy matrix, with the oxidation
of the graphite differentiating the nanocomposite materials.
1.3.1 Epoxy Blend Matrix
Over recent years, numerous publications have reported advances in expanded
graphite nanocomposites. They have attracted considerable interest as the graphite
morphology is very similar to that of layered silicate clays. Graphite also has a strength and conductivity resembling that of carbon nanotubes, but at a fraction of the cost. The issue associated with dispersion of graphite is the low surface energy of the planar sheet, resulting in poor wetting by the matrix. Dispersion and wetting are improved by
10 oxidation of the basal plane, but at the expense of transport properties. Therefore the purpose of this work is to:
(1) Disperse expanded graphite, epoxy functionalized graphite, and oxidized graphite in an epoxy matrix.
(2) Find a balance between graphite surface treatment and dispersion.
(3) Optimize the mechanical properties and electrical conductivity of the nanocomposites.
11
CHAPTER II
OBJECTIVES
Two separate resins are used as nanocomposite matrices in this work; a high temperature polyimide (PMR-15) and a commercial epoxy resin. The matrices are chosen for distinct purposes. PMR-15 is prepared by Polymerization of Monomeric
Reactants, where the average molecular weight between crosslinks is 1500 g/mol. PMR-
15 is the current state of the art high temperature polyimide, and is used as a matrix for
the composite duct in the F404 jet engine. Improving upon the existing thermal and
mechanical properties of this material is desirable to extend the lifetime and durability of
the resin. This translates to increased safety over existing parts. The epoxy is chosen to
represent a material for low temperature aerospace applications, such as cryogenic
propellant tanks. Additionally, the epoxy matrix serves as a proof of concept vehicle for multifunctional nanocomposites. Both are thermosetting resins, however PMR-15 is prepared in solution and cured at 315oC, whereas the epoxy resin is prepared from a melt
and cured at 125oC.
2.1 PMR-15 Matrix Nanocomposites
PMR - type polyimides are thermosetting polymers which combine excellent processability, mechanical properties, and thermal oxidative stability (TOS).47 Many
12 approaches have been taken to further improve the thermal oxidative stability and
mechanical performance of PMR-15, commonly by modifying the chemical structure of
the resin itself.48-51
As an alternative, incorporation of layered silicate clays into PMR-15, and other
high performance polyimides, has been investigated as a means of enhancing the
properties of this class of materials.52-56 Dispersion of the silicate layers has proven
challenging as the solvent based synthesis offers little shear forces to promote mechanical
separation of the layers.57 For such systems, alternative dispersion methods have been investigated. Jana and co-workers58 forced the layer separation by intercalation of a low molecular weight monomer or oligomer within the clay galleries. This increases the elastic forces inside the galleries, resulting in exfoliation. Researchers at NASA Langley
Research Center (LARC) have used a combination of high shear homogenization and sonication to introduce shear forces into the solution. This technique led to exfoliation of the clay in cured specimens.38 Many researchers varied the organic treatment on the clay surface to promote favorable interactions with the polymer, often finding that thermally stable and chemically reactive clay modifiers promote bonding with the polymer matrix and improve the mechanical and thermal properties of PMR-15.39, 58-60
The goal of this work is to improve the thermal oxidative stability of PMR-15 by
dispersion of organically modified layered silicate clays. The high processing
temperatures required to cure PMR-15, as well as the absence of shear forces to aid in
dispersion, present a challenge for incorporating an organically modified silicate into the
PMR-15 matrix. The processing temperature of up to 315oC degrades the alkyl
ammonium surfactants that are commonly used to modify clays. The focus of this work
13 is to develop a modification system for the silicate surface that (1) aids in silicate layer separation, (2) covalently bonds with the matrix, and (3) tolerates the polyimide processing temperatures. To realize these goals, it is necessary to develop a fundamental understanding of the orientation of the organic modifier and the effect the modified clay has on the oligomer melt viscosity and crosslinking reaction.
2.2 Epoxy Resin Matrix Nanocomposites
The use of epoxy resins in high performance applications places numerous demands
on these materials. Currently, several epoxides are being evaluated for application in
lightweight, carbon fiber reinforced composite cryogen storage tanks. This function
requires that the material resist microcracking on thermal cycling, have low permeability
to liquid and gaseous hydrogen and/or oxygen, and perform over a wide temperature
range. The rigorous criteria require optimized strength, toughness, barrier performance,
and dimensional stability of the resin. Therefore, layered silicate clay nanocomposites
are investigated as a means of meeting the multiple material performance requirements.
Furthermore, epoxy is a workhorse resin and therefore may serve as a proof of
concept matrix for new nanocomposites. Nanocomposites that impart multifunctionality
to a polymer, for example, are being widely investigated in industry.61-64 Exfoliated graphene nanocomposites are attracting interest as a multifunctional material capable of imparting both strength and stiffness to a resin, as well as enhanced thermal and electrical conductivity. Such capabilities expand the range of applicability for lightweight polymer components. For example, a thermally conductive polymer may be used in a higher temperature environment, if it is able to conduct the heat away from itself. An
14 electrically conductive polymer may be able to direct current to a specific location, providing protection against lightning strike. The challenge presented by graphite nanoparticles is in determining the balance between wetting and dispersion of the pristine graphite to yield optimized transport properties, or oxidation of the graphite to improve wetting and mechanical strength, at the expense of transport properties.
2.3 Summary
The addition of nano-particles to a polymer matrix has proven to be an effective
mechanism to enhance polymer matrix properties. Dispersion of nano-sized fillers can
improve matrix strength, modulus, thermal stability, impact resistance, barrier
performance, and conductivity. Materials such as POSS, silica, layered silicates,
exfoliated graphite, and nanotubes have been successfully dispersed in numerous
polymer systems.17-31 The effects of nano-fillers on the properties of thermosets, thermoplastics, rubbers, liquid crystalline polymers, copolymers, block copolymers, and polymer matrix composites have been reported.66-70 The opportunities to produce and optimize polymer nanocomposites are endless, due to the ability to functionalize the nano-particles; thereby strengthening the interface and maximizing interaction with the polymer matrix.
The purpose of this research is to demonstrate the importance of the interface in
polymer-nanocomposites systems. We investigate layered silicate clay, expanded
graphite, exfoliated graphite, and epoxy functionalized graphite. Two polymer matrices
were chosen, a thermosetting polyimide developed for high performance, high
15 temperature aerospace applications, and a general use thermosetting epoxy resin. In each case, the variables influencing the interfacial interaction are identified and optimized.
Due to the number of nanoparticles and resins investigated, this dissertation is divided into three sections. Chapters 4 and 5 contain the research and results from the layered silicate clay nanocomposite work. The scope of the layered silicate clay work is broad, and includes modification of the interface between clay and a PMR-15 matrix for applications in the 290oC temperature range, and interface considerations with a commercial epoxy resin matrix for cryo-tank applications requiring material performance as low as -250oC. Chapter 6 describes the research and results of the graphite
nanocomposite investigations. The matrix resin in this section consists solely of the
commercial epoxy resin. The goal is to optimize dispersion, interfacial chemistry, and
prove the multifunctional capabilities of the graphite nanomaterial.
16 CHAPTER III
BACKGROUND
3.1 Conventional Fillers
Particulate fillers have played a vital role in the development of commercial uses for
polymers. The primary filler types utilized may be classified as either natural or
synthetic fillers. Naturally occurring fillers are generally a pure, crystalline mineral. A
common natural filler is calcium carbonate, CaCO3, which has been used in polyvinyl chloride, polypropylene, elastomers, and unsaturated polyesters. Other filler materials include micron sized clay particulates such as kaolinite and metakaolin. These materials have been used to impart electrical resistivity to polyvinyl chloride for cable insulation applications. Crystalline silicas and calcium sulphate are additional commonly employed, natural fillers. Synthetic fillers are prepared by chemical processes, with carbon black being among the most well known. Carbon black has been widely used for rubber reinforcement. Synthetic silicas are another example of a synthetically prepared
filler. Silica particulates have found wide application in silicone elastomers.71
Originally, fillers were primarily considered cheap diluents. However, their ability to
beneficially modify polymer properties was quickly realized. Some of the main reasons
for using particulate fillers include: cost reduction, improved processing, thermal
conductivity, controlled thermal expansion, flame retardancy, and improved mechanical
17 properties. Of course no single filler has provided all of these benefits. Ideally a filler improves some properties without negatively affecting others. The magnitude of the property change observed is not only a function of the filler composition but is strongly influenced by particle size, shape, and surface chemistry.71 Particle size, shape, and ability to bond with the polymer matrix are all important factors in determining filler performance.
3.2 Comparison of Conventional and Nano Fillers
The literature is limited in direct comparison of conventional and nanofillers. Those publications that have addressed the difference in material properties that resulted from reduced filler dimensions all report that nanofillers outperform the conventional fillers.72-
73 Other researchers have compared aggregated to fully dispersed nanofillers.
Nazarenko, et. al.74 examined the effect of clay layer aggregation on the gas barrier performance of several polystyrene/ clay nanocomposites. The extent of aggregation was varied by selecting (1) an ion exchange compound which favored polystyrene intercalation, (2) a compound that favored layer exfoliation, and (3) they mixed unmodified clay with the polymer to yield an aggregated polymer/ clay composite. The exfoliated system however was not completely exfoliated, but rather consisted of randomly oriented stacks of disordered layers. This is representative of the morphology
observed in most ‘exfoliated’ nanocomposite systems. The researchers were able to
demonstrate that ordered stacking, or ‘intercalation’ of layers provided lower
permeability than the disordered stacking. Furthermore, the data showed that the
18 composites composed of clay aggregates displayed the highest permeability of the three
clay systems.
The barrier performance of polymer clay nanocomposites has been modeled by
several groups including the work of Bharadwaj,206 who explained that the reduced permeability may be attributed to the increased path length over which a permeant must travel to diffuse through the material. This ‘tortuous path’ model was also dependent on layer aggregation and orientation, as illustrated in Figure 3.1.
Dasari, et. al.75 prepared Nylon-6 clay nanocomposites with varying degrees of exfoliation. As with the work of Nazarenko,74 they also used unmodified clay to prepare a composite containing aggregated clay layers. They found that blending organically modified clay into Nylon-6 resulted in reasonably good layer separation; however this separation was enhanced by pre-swelling the clay in water before processing. The composite tensile properties are listed in Table 3.1. The researchers found that greater layer separation led to the most significant improvements in Young’s modulus and yield strength. However, the increase in strength and stiffness of the nanocomposite materials resulted in a significant decrease in the elongation at break, from 96% in the neat Nylon-6 material, to 16% in the highly exfoliated nanocomposite.
Hussain, et. al.29 described the transition from micro-particles to nano-particles as having dramatic effect on the physical properties of polymers. This was attributed to the large surface area that nanoscale additives have for a given volume.
As many chemical and physical interactions are governed by surfaces and surface properties, a nanostructured material will impart substantially different properties compared to a larger-dimensional material of the same composition. For layered
19
Permeant Permeant Permeant
Permeant Permeant Permeant
Figure 3.1. Depiction of ‘tortuous path’ incurred by a permeant traveling through
polymer filler with nanoclay at different orientations and degrees of dispersion
20
Table 3.1. Literature data comparing physical properties of Nylon-6, and Nylon-6/ clay
composites; where clay is dispersed by varying degrees75
Tensile Properties Specific Wear Rate,(10-6mm3N-1m-1) Sliding velocity Sample Young’s Yield Elongation 0.08 m/s 0.12 m/s 0.20 m/s Modulus Strength at Break (GPa) (MPa) (%) Neat Nylon 6 2.16 ± 0.11 71.1 ± 0.7 96 ± 18 7.2 ± 2.7 33.4 ± 3.2 107 ± 5.3 Nylon 6/ 2.55 ± 0.05 74.1 ± 0.7 25 ± 4 26.7 ± 4 86.3 ± 5.7 195.8 ± 14 unmodified clay Nylon 6/ 2.77 ± 0.07 73.8 ± 1.11 24 ± 4 17.4 ± 40.6 ± 4.4 76.3 ± 6.5 organically 1.6 modified clay Nylon 6/ 3.32 ± 0.03 79.7 ± 0.9 16 ± 5 21.4 ± 57.7 ± 3.9 119 ± 7.3 organically 0.6 modified clay, dispersed with aid of swelling agent.
21 materials, such as silicates or graphenes, the surface area per unit volume is inversely
proportional to the material’s diameter, therefore, the smaller the diameter, the greater the
surface area per unit volume. Reducing the dimensions of a platelet material from the
micrometer to nanometer size range affects the surface area to volume ratio by three
orders of magnitude. Therefore, effective dispersion of the particles is essential to realize
the full potential of the nanoparticles increased surface area.
The current interest in polymer matrix nanocomposites has been fueled in large part
by the 1991 findings of Toyota, where dispersion of layered silicate clays in a Nylon 6
matrix greatly increased the heat distortion temperature and the mechanical properties of
the pristine polymer.76 Coincidently, in the early 1990’s, Ijima77 reported on the synthesis of single wall carbon nanotubes. This discovery in particular sparked considerable interest in the potential of nanotechnology, driving the consistent increase in funding and research over the past 15 years.
Polymer matrix nanocomposites offer several advantages over conventionally filled composites. (1) They are lighter in weight, as less filler is required to attain high levels of strength and stiffness, (2) the platelet morphology of clays or graphenes offers exceptional diffusional barrier properties, (3) carbon based nanoparticles offer exceptional thermal and electrical conductivity at low filler loading, and (4) the low filler loading allows polymer processing without the need for costly tooling.78-79
While a vast number of nanoparticles are available, this work will focus on the
characterization, modification, and dispersion of layered silicate clays and graphite
nanoparticles.
22 3.3 Layered Silicate Clays
The term layered silicate clay nanocomposite generally refers to the dispersion of a 2
or 3 layer silicate clay into individual, nanometer thick sheets throughout a polymeric matrix. The layered silicates utilized for such constructions have an expandable gallery between sheets, unlike the more familiar silicates such as talc or mica.
The principal elements of the layered silicates are two-dimensional arrays of silicon- oxygen tetrahedral sheets and two dimensional arrays of aluminum- or magnesium- oxygen-hydroxyl octahedral sheets, as illustrated in Figure 3.2. The combination of an octahedral sheet and two tetrahedral sheets is referred to as a unit layer. Most clay minerals consist of unit layers stacked parallel to each other. The distance between a plane within a unit layer and the corresponding plane in the next unit layer is called the
001 spacing, the basal spacing, or the d-spacing. In montmorillonite clay, the clay primarily of interest for polymer nanocomposites, the 001 spacing between consecutive layers equals 1 nm. The lateral dimension of the unit layer may range from 100 nm to 1
µm.80
Within the tetrahedral sheet of montmorillonite, tetravalent Si is partly replaced,
naturally, by trivalent Al. In the octahedral sheet, there may be replacement of trivalent
Al by divalent Mg, without complete filling of the third vacant octahedral position. Al
atoms may also be replaced by Fe, Cr, Zn, Li, or other atoms. The small size of these
atoms permits them to take the place of the Si and Al atoms, therefore, the replacement is
often referred to as isomorphous substitution. Commonly, an atom of lower positive
valence replaces one of higher valence, resulting in a net negative charge within the clay
unit layer. The excess negative lattice charge is naturally compensated by the adsorption
23
Figure 3.2. Chemical structure of montmorillonite clay
24 of cations on the layer surfaces. These cations, often Na+ or Ca2+, are too large to be
accommodated in the interior of the lattice and therefore may be easily exchanged by
other cations when available in solution. The total amount of these “exchangeable
cations” may be determined analytically. This quantity is expressed in milliequivalents
per 100 grams of dry clay, and is termed the cation exchange capacity, (CEC).81 The CEC represents an upper limit to the quantity of organic modifier which may be bound to a clay and therefore has a significant effect on the interfacial performance of a polymer- silicate nanocomposite.
3.4 Organic Modification of Layered Silicate Clays
Replacing the naturally occurring metal cations between clay layers with a protonated organic compound serves two purposes. First, it increases the spacing between silicate sheets, facilitating intercalation of a monomeric or polymeric species. Secondly, the ion exchange renders the clay ‘organophilic’ rather than hydrophilic. The versatility of the ion exchange process allows one to tailor the chemistry of the clay surface to match that of the polymer matrix. As a result, the choice of organic modifier may have a profound influence on the dispersion and properties of the resulting bulk nanocomposite. A commonly utilized silicate modifier is an alkyl ammonium ion, of varying carbon chain length. Therefore, significant research has focused on the effect the length of the carbon chain has on nanocomposite formation, as well as understanding the orientation and reactivity of the modifier. All aspects are essential to the design of an organically modified silicate clay nanocomposite.
25 The mobility and reactivity of the organic modifier tethered to the silicate surface is of
interest as such characteristics influence nanocomposite dispersion and properties. It has been reported that the orientation of an alkyl ammonium ion is dependent on the packing density, temperature, and chain length.82 Such conditions dictate whether the long carbon chain surfactants lie parallel to the silicate surface as mono- or bilayers, or extend away
from the surface in mono- or bimolecular arrangements. Figure 3.3 illustrates these
arrangements. The structural characterization of organically modified clay is commonly
elucidated from d-spacing values calculated from X-ray Diffraction (XRD) analysis. As
illustrated in the figure, a paraffin type arrangement of the modifier offers the greatest
increase in layer separation, and may therefore be the easiest to disperse in a matrix resin.
Additional characterization has provided insight into the chain length and packing density requirements that direct a desired modifier arrangement.82
Vaia et. al.82 utilized Fourier Transform Infrared Spectroscopy (FTIR) to characterize the molecular conformation of alkyl ammonium surfactant chains within montmorillonite clay galleries. Their study yielded three primary conclusions. First, the disordered
(gauche) conformation was favored over the ordered (trans) conformation. Second, increased packing density or chain length lead to a more ordered arrangement of the chains, and third, increased temperature favored the disordered conformation.
Wang and co-workers83 employed solid state NMR to examine the conformation and mobility of alkyl ammonium modifiers on montmorillonite clay. Resonance peaks for both ordered (trans) and disordered (gauche) conformations were resolved. The conditions which led to the trans and gauche conformers confirmed the earlier study by
Vaia.82 With respect to the mobility, Wang et. al. found that the molecules in the ordered
26
N+ N+ N+ N+
N+ N+
N+ N+ N+ N+ N+
N+ N+
N+ N+ N+
Figure 3.3. Arrangement of modifiers between layers as determined by XRD and FTIR
27 conformation were as rigid as those in most solid crystalline materials, and that the
molecules in the disordered conformation were similar in mobility to liquid crystalline
materials. On heating, the molecules in the disordered conformation remained
unchanged while those in the ordered conformation attained a more liquid-like quality.
Additionally, researchers have varied the modifier chain length to optimize
modification for dispersion in a specific resin system. Pinnavaia, et al.,40 found that when
using an alkyl ammonium modifier, a chain length of at least 8 carbons was necessary to
disperse the silicate in an epoxy matrix. The group also noted that complete exchange of
the naturally occurring sodium cations by the organic ammonium ion was equally
important for silicate layer separation.
Usuki, et al.,59 modified a montmorillonite clay with a series of ω-amino acids
+ [H3N (CH2)n-1COOH] of increasing chain length. The organically modified clay was intercalated with ε-caprolactam and polymerized, to prepare Nylon-6 nanocomposites.
The greatest dispersion was observed with clays where the modifier was an amino acid having chain lengths greater than 11 carbons.
Similarly, Fuskushima, et al.,84 report unlimited swelling of a montmorillonite clay ion exchanged with protonated 12-aminododecanoic acid, in a Nylon-6 matrix. The researchers concluded that the driving force for the silicate layer separation was two fold and included the polymerization energy of ε-caprolactam and the attractive force between the ε-caprolactam and the interlayer cation.
A study by Delozier, et al.85 dispersed a series of alkyl ammonium modified clays into a polyimide matrix. They noted a darkening of the nanocomposite films following imidization which was not observed in the control material. This darkening was
28 attributed to degradation of the alkyl ammonium modifier on the clay surface. XRD
studies verified the degradation by a marked shift of the d001 diffraction peak, corresponding to a decrease in silicate gallery height. A similar result was reported by
Gintert et al.,86 where PMR-15 nanocomposites were prepared with alkyl ammonium modified clay. In this case, XRD data also showed a decrease in d001 spacing on high temperature curing of the imide oligomers. As a result of studies like these, researchers have looked to rigid, aromatic diamines to modify clay for dispersion in high temperature polymers.38-39,56,58,86 The aromatic structure offers thermal stability, while the additional amine functionality provides a mechanism for chemical reaction with polymer precursors.
Wei, et al.87 investigated di- and tri- functional amine modifiers. They reported a significant increase in the onset temperature of polymer degradation upon addition of the
modified clay into a polyimide matrix. Other observations included a reduction in the
coefficient of thermal expansion and an increase in the percent elongation.
Liang et al.88 compared nanocomposite properties obtained following dispersion of
hexadecyl ammonium modified clay with those of protonated aromatic amine modifiers.
They reported exfoliation in all cases when clay loading was below 3 wt%. On heating,
XRD demonstrated a collapse in the clay gallery height of the hexadecyl ammonium
modified clay, whereas the aromatic treated clay retained its morphology. Finally, the
nanocomposites prepared with aromatic modified clay exhibited better mechanical and
thermal properties.
The amount of modifier present between clay layers, as well as the chemical structure of the modifier, affects the interlayer chemistry of the layered silicate clays. This, in turn,
29 contributes to the morphology of the polymer nanocomposite. As a result, the choice of
silicate modification will be sensitive to the matrix resin of interest. Furthermore, a
modifier must be chosen that is amenable to the processing requirements of the resin.
Consequently, there has been significant research into the role of the organic modifier on
nanocomposite processing and properties.89-92
3.5 Dispersion of layered silicate clays
Three generalized morphologies exist upon introduction of stacked layered platelet particulates, such as clay or graphite, to a polymeric system. Those morphologies include aggregation, intercalation, and exfoliation. These structures are depicted in Figure 3.4.
Both intercalated and exfoliated morphologies are considered nanocomposites because in these cases, each 1 nm thick silicate layer is in contact with the polymer. This provides the large interface that is necessary to achieve optimal nanocomposite properties.82
However, within the field of polymer/ clay nanocomposites, it is generally agreed that the greatest extent of layer separation will provide the best properties, such as improved barrier and mechanical performance compared to the base resin. Considerable effort is therefore focused on obtaining the optimum modification and dispersion of the clay in a matrix.
However, nanoparticles are notoriously difficult to disperse, due to their large surface area which gives rise to strong interparticle electrostatic interactions. Heinz et al.93 showed that van der Waals forces are the primary force keeping organically modified clay from readily dispersing. Their work demonstrated that these interactions are very
30
Figure 3.4 Potential polymer-clay morphologies
31 strong, but only at distances less than 0.5 nm. These forces must be overcome during
nanocomposite processing in order to exfoliate clay into individual platelets.
There have been a number of papers reported in the literature describing the driving force for layered silicate clay nanocomposite formation. Vaia and Giannelis 94 have used a simple lattice-based thermodynamic model that examined the entropic and enthalpic contributions during the formation of a polymer layered-silicate nanocomposite to understand the driving forces for intercalation and exfoliation of organically modified layered silicates by long-chain polymers. In their estimation, despite the expected loss of conformational entropy of the confined polymers, the gain in conformational entropy of the surfactant tails compensated. This lead to the conclusion that the enthalpy of mixing dominated the free energy considerations.
Meneghetti and Qutubuddin95 extended this model beyond a polystyrene matrix to include polymers of varying polarity. They showed that the resulting nanocomposite morphology was dependent on the free energy change of the system, ∆f, as described by the equation.
∆f = f(h) – f(h0) = ∆E - T∆S
The change in free energy was influenced by both the change in internal energy (∆E), and the change in entropy (∆S). The internal energy change was associated with layer
expansion, polymer intercalation, and the formation of new intermolecular interactions,
whereas the change in entropy was associated with conformational change of the polymer
and the organic modifiers on the silicate surface. They verified the model proposed by
Giannelis and Vaia94 by demonstrating that polymers of increasing polarity and
interaction with the silicate, increased silicate layer separation.
32 Models by Balazs, Singh, and Zhulina96 predicted that the length of the surfactant should influence the clay dispersion by affecting the degree to which the entropic penalty to an intercalating homopolymer was reduced. They concluded that lengthening the surfactant should facilitate intercalation by providing a greater reduction in the entropic penalty to the intercalating homopolymer.
A recent model by Mackay et. al.97 suggested that the dispersion of nanoparticles in a polymer was a result of a favorable enthalpy of mixing due to increased molecular contacts between the polymer and the dispersed nanoparticles. This was attributed to the increased accessible area on the nanoparticle caused by dispersion. Their results indicated that nanoparticles were capable of being dispersed in polymers if the size of the nanoparticle was smaller than the radius of gyration of the polymer matrix.
The dispersion challenge has been met by numerous dispersion techniques. The most general methods include in-situ polymerization, solution mixing, and melt blending.
Other common techniques, such as high shear mixing, can be considered a subset of the previously mentioned methods. In-situ polymerization is appropriate for materials where polymerization may readily take place in the presence of clay, and for thermosetting resin which are not melt processable. Nylon-6/ clay nanocomposites are often prepared by the in-situ polymerization of ε-caprolactam.59 This method allows pre-swelling of the silicate
layers by the ε-caprolactam monomer, thereby facilitating separation of the layers on
polymerization. Clay nanocomposites prepared by this method, dispersed in Nylon 6,
have produced some of the largest clay layer separations observed, relative to other
matrices. Several groups have demonstrated that the method of pre-swelling clay with a
33 monomeric reactant, followed by in-situ polymerization, may also produce clay
dispersion in epoxy matrices as well.59,98-100
Melt processing requires that the formed polymer is capable of sufficient melt flow, so
that it may intercalate and disperse silicate layers. For example, dispersion of an
organically modified montmorillonite in polystyrene was achieved by melting PS in the
presence of clay, and allowing the polymer to penetrate the clay layers.101
Often, more rigorous mixing is required. Dennis et al.102 melt processed Nylon-6 nanocomposites using different types of extruders and varying screw design. By this method, they were able to study the effect that shear speed had on silicate dispersion.
They found that both residence time in the extruder and shear speed had affected dispersion. Specifically they noted that high shear intensity was necessary to initiate the
dispersion process, and that this high shear intensity reduced the residence time
requirement. At lower shear, longer residence times were required to disperse the layers.
Aside from high shear, melt processing may also be assisted by the introduction of
chemical functionalities onto the clay surface, to provide interaction sites for the
intercalating polymer. Kawasumi et al.103, described a technique to melt blend organically modified clay and polypropylene (PP). In this case, PP was mixed with distearyldimethylammonium ion, and a polyolefin oligomer containing polar telechelic
OH groups. This compound served as a compatibilizer. By this process, the compatibilizer first intercalated the clay layers through strong hydrogen bonding between the OH groups of the compatibilizer and the oxygen groups on the silicate. This swelled the clay interlayer spacing and therefore allowed layer separation on introduction of the melted PP.
34 3.6 Dispersion Mechanism
Increasing the degree of clay layer dispersion increases the interfacial area available to
the matrix. Therefore, silicate dispersion is a primary requirement to obtaining desirable
nanocomposite properties. As a result, a considerable amount of research has focused on
understanding the mechanism for dispersion.
Giannelis et al.42 used time-resolved, high temperature XRD to understand the exfoliation mechanism of a high-Tg epoxy nanocomposite. They found that the clay interlayer expansion mechanism could be resolved into three stages; initial interlayer expansion, steady state interlayer expansion, and the cessation of interlayer expansion. It was found that the activation energies of interlayer expansion and curing influence the final nanocomposite morphology. In the first stage, interlayer expansion was initiated by intragallery polymerization. Throughout the second stage, expansion was characterized by a linear increase in interlayer spacing. Expansion during this stage was dependent on diffusion of unreacted resin into the intragallery spacing. It was observed that samples exhibiting large interlayer expansion also had an activation energy of expansion that was lower than the activation energy of curing. The third stage was described by a slowing of
interlayer expansion. Layer separation then stopped as the modulus of the extragallery polymer became equal to the modulus of the intragallery polymer.
Benson-Tolle and Anderson104 studied the relationship between processing and
morphology development in a thermosetting epoxy matrix nanocomposite. The study
used small angle x-ray scattering (SAXS) to follow changes in the d001 clay layer stacking
during isothermal experiments. They found that the silicate layers expanded as
polymerization and network formation proceeded. The separation was monitored by
35 XRD and provided evidence of layer separation by a shift in the d001 diffraction peak, as well as a broadening of this peak. This suggested that the network formation between
layers may not be uniform and therefore created a loss of stacking registry. As the
experiment continued, the peak intensity continued to decrease until the layers were
considered exfoliated. This study also showed that the cure temperature played a role in
exfoliation. As the cure temperature was increased, the onset of exfoliation began earlier
and was completed more quickly than at lower temperatures.
Other researchers have focused on the chemistry of exfoliation. For example,
Pinnavaia et al.105 found that the alkyl ammonium ion used to modify the silicate may initiate epoxy polymerization. The studies concluded that the rate of intragallery polymerization should not be much higher than that of the extragallery polymerization,
which could lead to clay layer aggregation. Here again, the researchers found the choice
of alkyl ammonium ion to be a critical variable. Specifically, they stated that intragallery
catalysis by the modifier was evidenced by a decrease in layer exfoliation with
decreasing Bronsted acidity of the exchange ion. In terms of silicate exfoliation, the
research concluded that with appropriate curing conditions, greater concentrations of
epoxy and curing agent could penetrate the gallery space and allow intragallery
polymerization to occur at a rate comparable to extragallery polymerization.
Consequently, the galleries continued to expand as the degree of polymerization
increased and a monolithic exfoliated nanocomposite was formed. However, if the
curing temperature was too low and the rates of epoxy and curing agent intercalation
were too slow, than extragallery polymerization would be faster and intercalated
nanocomposites will form.
36 Park and Jana106 took this a step further and studied the influence that elastic forces
developed in the clay galleries had on exfoliation. They found that the elastic force
exerted by the cross-linking epoxy molecules inside the clay galleries pushed out the
outermost clay layers from the tactoids, against the opposing forces arising from
electrostatic and van der Waals attraction, and shear. Exfoliation continued until the
extragallery epoxy turned into a gel. It was found that complete exfoliation of the clay
was produced when the ratio of shear modulus to complex viscosity was maintained
above 2 - 4 l/s, such that the elastic forces inside the galleries outweighed the viscous
forces of the extragallery epoxy.
Clay dispersion in a polymer melt, under conditions of relatively high shear,
introduces alternative mechanisms to dispersion. Intuitively it would seem that high
shear would yield well dispersed clay. However, this is not always the case. Dennis, et
al.102 reported separate conditions for dispersion of clays in a Nylon-6 matrix. They found that the degree of shear introduced to the system during processing affected the
mechanism by which the clay layers separated. The firs method involved high shear. In
this case, the number of platelets in the clay tactoids was reduced by the individual
platelets sliding apart. An alternative, low shear, mechanism showed the polymer chains
entering the clay galleries and pushing the end of the platelets apart. This pathway
involved diffusion of the polymer into the clay galleries, and was driven by either
physical or chemical affinity of the polymer for the organoclay surface. Such a
mechanism would be facilitated by increased residence time in an extruder.
37 In short, while there are numerous techniques available to compatiblize and disperse
silicate clays into just about any polymer, the level of dispersion and the organic modifier
remain key in ensuring optimal interfacial interaction.
3.7 Polymer-Silicate Interface
Using nanocomposites in design-critical applications requires an understanding of
their structure-property-function relationships. Furthermore, the ability to tailor the
filler/matrix interaction and an understanding of the impact that the interface has on
macroscopic properties are key to obtaining nanocomposites with the desired properties.
Understanding the nature and chemistry of the interface introduces tailorability of the
nanocomposite, providing compatibility with a number of different polymer matrices.
Tailoring is often achieved by grafting short molecules or polymer chains with precise
chemical structure from the nanoparticle surface.
Korley et al.107 reported the tailorability of an elastomeric polyurethane nanocomposite through selective interactions between a layered silicate clay and the matrix. Their research found that in one system, the hydrophilic, polar soft block polyurethane segments dominated the clay/polyurethane interactions. In this case, strain induced alignment of the soft segment chains was suppressed within the nanocomposite, resulting in a reduction of polymer toughness and extensibility, as compared to the neat
resin. Comparably, the silicate layers in a polyurethane containing a hydrophobic soft
segment resulted in the clay favoring interaction with the hard segment. This
morphology offered enhanced polymer toughness and modulus.
38 3.8 Nanocomposite Properties
Enhancements in polymer properties as a result of nano-filler addition vary from
polymer to polymer and can be dependent on the level of silicate dispersion, the organic
modifier, silicate loading, and the test temperature.40,108-109 Increased strength, modulus,
barrier properties, and dimensional stability (reduced coefficient of thermal expansion,
CTE) have been observed in most systems as a result of dispersing the rigid nanoparticle
in the softer polymer matrix. However, many nanocomposites exhibit a toughness lower
than the corresponding neat resin.110
Drzal, et al.111 dispersed layered silicate clays into an amine cured diglycidyl ether of bisphenol A (epoxy resin). They reported a 50% increase in the room temperature storage modulus at a 10 wt% clay loading. Enhancements in the tensile strength were only observed with 2.5 wt% loading, and the Izod impact strength was radically decreased with increasing clay loading.
Bao et al.112 also reported a reduction in impact toughness following dispersion of
layered silicate clay in high density polyethylene, where the reduction in toughness
corresponded with an increase in Young’s modulus. A reduction in impact toughness
was also observed on clay dispersion in HDPE as well as in rubber toughened HDPE.
Other researchers have noted that the negative influence of the nanofiller on toughness is not observed when testing above the glass transition temperature. Misra et al.113 reported that the addition of clay to polypropylene increased the impact strength in the temperature range of 0oC to +70oC, whereas the clay had little to no effect at
temperatures ranging from -40oC to 0oC.
39 114-115 Pinnavaia et al. noted that flexible resin systems with low Tg showed a much larger increase in modulus and tensile strength with the addition of an organoclay, than did rigid systems.
Giannelis et al.116-117 provided experimental evidence that tensile loading of a
nanocomposite at a temperature much higher than Tg, allowed nanoparticle alignment in the direction of the stress and thereby provided a mechanism for energy dissipation.
There is less discrepancy in the results observed in barrier properties, with several researchers reporting order of magnitude reductions in the gas permeability of the layered silicate nanocomposites, compared to the pristine resins.
The dimensional stability of the nanocomposite is often of interest in polyimide nanocomposites, where elevated temperature applications are a common goal, for example in microelectronics, semiconductors and composites. Researchers often report a decrease in the coefficient of thermal expansion, generally in the range of 25% to 50%, depending on the silicate loading. The reduction in CTE is attributed to the dispersion of rigid particles of high surface area which restrict the motion of the polymer chains, thereby improving dimensional stability.118-119
3.9 Summary of layered silicate clay nanocomposites
The past fifteen years have seen a considerable amount of work performed in the field
of polymer and layered silicate clay nanocomposites. Most often, the clay is modified
with an organic compound prior to introduction to the matrix material. The choice of
modifier and the chemical nature of the matrix result in an interface that is unique to each
40 polymer-modified clay system. Therefore, careful thought should be applied to each
composite system to achieve the desired nanocomposite bulk properties.
3.10 Expanded Graphite
Expanded graphite (EG) is analogous to layered silicate clay with respect to its
stacked platelet composition. The individual graphene platelets are held together by relatively weak intermolecular van der Waal’s forces. Therefore, when dispersed throughout a polymer matrix, these sheet-like materials may provide many of the property enhancements that have been reported in the layered silicate clay nanocomposites.
There are features inherent to this material; however that separate it from layered
silicate clays. The uniqueness of the graphite nanoparticles is in the strength and conductivity of the graphene sheet. The individual graphene layers are strong, with
single crystal graphite having an elastic modulus of over 1 TPa.120 They are also highly conductive, with single sheet electrical conductivities reported as 104 S/cm at room temperature.121 Therefore, expanded graphite provides an opportunity to develop truly multifunctional materials that can bear mechanical loads as well as provide electrical conductivity, for example.
The precursor to expanded graphite is natural flake graphite (NFG). Like EG, NFG is composed of numerous graphene layers held together by van der Waals forces, as illustrated in Figure 3.5. The dimensions of the stacked platelets comprising NFG range from 5 to 20 µm in length with 0.34 nm spacing between layers. The weak interplanar
41
Figure 3.5. Structure of layered natural flake graphite
42 force between graphene layers allows acid intercalation. Typical EG preparation
proceeds by intercalation of natural graphite flakes with a combination of concentrated
sulfuric and nitric acids. The graphite intercalation compound (GIC) is rapidly heated to
approximately 1000oC, expanding the particles in the c direction to approximately 300 times that of the original dimension.122,123 The resulting expanded graphite structure
consists of stacked aggregates of graphene sheets. This structure can be separated into
smaller aggregates by sonication in alcohol or upon addition to a polymer matrix, as
illustrated in Figure 3.6.124-125
Attaining useful composite properties with expanded graphite is dependent on the
extent of its dispersion throughout a resin matrix. The literature offers numerous
examples of nanocomposites prepared by sonication or simple mixing of EG into a
polymer matrix, or a monomer solution.126-130 Enhanced polymer properties have been reported, however, the graphite is often present in regularly stacked aggregates composed of several graphitic layers. Such aggregation is evident by XRD, where an intense
o diffraction peak at 2θ = 27.6 is observed, corresponding to the d002 diffraction plane of
NFG.131-133
EG aggregate dimensions commonly reported range from 2-5 nm in thickness to as
high as 20 to 50 nm in thickness.122 As the thickness of the aggregate increases, the
surface area available to the matrix decreases, and the benefits of the nanostructured
material are diminished. Yasmin and Daniel134 reported improved properties of an epoxy resin by the addition of 2.5 wt% and 5 wt% expanded graphite. They acknowledged that the graphite was not well dispersed in this study, and the average aggregate thickness was in the range of 250 nm. Regardless, they noted a 50% reduction in the epoxy CTE
43
Figure 3.6. Expansion and dispersion of natural flake graphite
44 as well as increased glass transition temperature, onset of decomposition, and storage
modulus. Therefore, it was anticipated that dispersion of the graphite aggregates to a
finer level, would significantly benefit material performance.
3.11 Dispersion in Expanded Graphite
Controlling the dispersion of nanoparticles in polymeric matrices is the most
significant impediment in the development of high performance polymer nanocomposite
materials and results primarily from the strong interparticle interactions between the
nanoparticles. The literature studies to date, on the dispersion mechanism of exfoliated graphite are considerably limited relative to those of layered silicate clays.
While dispersion mechanisms have not been extensively investigated within graphite
nanocomposites, the dispersion of carbon nanotubes has. Carbon nanotubes are similar to
graphite in chemistry and by the nature of their inert surface. Carbon nanotubes can be
difficult to disperse into a polymer for a number of reasons, including the large surface
area which encourages aggregation due to intermolecular interactions, such as van der
Waals forces.135 Such interactions are present in expanded graphite as well. Furthermore,
the planar surface of EG is inert, and functionalization occurs primarily along the plane
edge. Functional groups often observed include semiquinone, carboxyl, and hydroxyl
moieties.
In general, graphite dispersion to the level of individual graphene sheets has proven difficult. As previously described, this is attributed to the low surface energy of the graphene sheet, and subsequent poor wetting by the matrix polymer. There are several
45 examples in the literature of approaches that have been taken to improve the interaction
between graphite and polymer, thereby improving dispersion.136-140
3.12 Dispersion Techniques
It is generally accepted that improving graphite separation as well as increasing
interfacial strength should benefit graphite nanocomposite performance. The dispersion
methods used for graphite-polymer nanocomposites are similar to those employed with
silicate clays, as both materials have a layered structure and require monomer or polymer
intercalation as a preliminary step to dispersion. However, because the materials differ chemically, the surface treatment of graphite can not follow the same chemistry as that of the layered silicate clays.
Unlike layered silicate clays, graphene sheets are not readily ion exchanged with long
chain, alkyl ammonium ions to swell the layers. Furthermore, the highly non-polar basal
plane of the graphene sheets offers little driving force for intercalation of polar resins.
Several processes have been developed to disperse graphite and the applied technique is
dependent on the processing and properties of the matrix resin. Direct mixing of
expanded graphite and a polymer is commonly used with low viscosity thermosetting
matrices, e.g. epoxy resins.141 Melt compounding is commonly used with polyolefins, e.g. polyethylene or polystyrene. Solution intercalation requires that the polymer is soluble.
With this technique, the expanded graphite is mixed with the dissolved polymer, followed by solvent evaporation to leave the nanocomposite. PMMA/graphite142 and
PP/graphite143 nanocomposites have been prepared by this method. Additional
techniques include in-situ polymerization, where a monomer is polymerized in the
46 presence of expanded graphite. Nylon 6/graphite,144 polystyrene/graphite,145 and polyacrylonitrile/graphite146 nanocomposites have all been prepared by this method. A
polymerization filling technique147 has been utilized to prepare epoxy matrix or
polystyrene matrix nanocomposites. By this method, in-situ polymerization occurs in the
presence of initiator intercalated graphite. The intercalated graphite is mixed with a
monomer, and graphite separation occurs during exfoliation.
The standard technique for acid intercalation and expansion of natural flake graphite
results in minimal functionalization of the graphite sheet edge. However, it is often
desirable to also chemically functionalize the graphene basal plane. The method most
commonly used to functionalize the basal plane begins with preparation of graphite
oxide. Preparation methods were developed in 1859 by Brodie,148 and later (1958) by
Hummers,149 both resulted in graphite layers decorated with epoxy, hydroxyl, carboxyl,
and quinone functionalities within the basal plane, however the layers retained a stacked
registry detectable by XRD. Such moieties allowed for polar interactions with a resin, or
covalent functionalization with a surfactant.150 Furthermore, graphite oxide is hydrophilic
and readily adsorbs water or other polar liquids.151
Taking the graphite oxide functionalization a step further, groups at Princeton152 and
Northwestern153 Universities separated the stacked graphite oxide layers by rapid thermal treatment. The rapid heating procedure degraded many of the functionalities present on graphite oxide, and generated CO2 in the process. The CO2 expansion on rapid heating
split the stacked graphite oxide sheets into individual graphene layers.154 The procedure reduces the graphite oxide to an extent, however several functionalities remained, as characterized by NMR studies. Addition of this material to both epoxy and
47 polymethyl(methacrylate) matrices has resulted in significant improvement in the
mechanical properties of these resins.155-156
Oxidation of the layered graphene sheets serves to increase the surface energy of the
graphite, thus facilitating both dispersion and adhesion to the matrix. The presence of functional groups also allows chemical modification of the basal plane, further increasing the strength of the interface.157-158
3.13 Polymer-Graphite Interface
Characterization and modification of the polymer – graphite interface has been rather limited, with most literature focusing on dispersion and properties.159-162 Generalized statements regarding the strength of the polymer-graphite interface may be made by examining the nanocomposite fracture surface at break. Drzal, et al.141 have worked to
improve the graphite interface and dispersion in a polypropylene matrix. They
functionalized graphite by application of a PP coating onto the graphite surface, followed
by sonication in alcohol to reduce aggregate size. The graphite was introduced into the
matrix by processing techniques such as extrusion, injection molding, and compression
molding. At 10 vol%, the nanocomposites showed an 8% higher strength and a 60%
higher modulus over composites made by simple melt mixing. These results
demonstrated the benefit of coating the graphite to improve dispersion and interfacial
strength. More importantly, the SEM images of the material fracture surface showed
large aggregates in the melt mixed material and a much finer dispersion of the coated
material.
48 3.14 Nanocomposite Properties
Improved electrical conductivity is the most commonly reported property enhancement
with expanded graphite nanoparticles. This is because graphite layer separation is
difficult, therefore while mechanical properties may not improve on incorporation of the
graphite, conductivity generally does, as separation of the expanded graphite is separated
into individual sheets is not necessary for conduction. Conductivity simply requires that
the percolation threshold for conduction is reached. Dispersion of the sheets, however,
does allow the conductive network to form at a lower volume fraction of filler.
Dispersion is always desirable and is therefore the starting point of most research. There
are numerous publications in the literature that describe the electrical properties of
expanded graphite nanocomposites.163-168
Zheng and Wong166 dispersed both graphite and expanded graphite into a
polymethylmethacrylate matrix and compared the electrical conductivity of the resulting
nanocomposites. The percolation threshold was reached at 1 wt% loading of the
expanded graphite, whereas addition of 3.5 wt% graphite was necessary for conductivity.
The conductivity after percolation was an order of magnitude greater for the expanded
graphite composites as compared to the graphite composites. This difference was
attributed to the greater dispersion of the expanded graphite.
Kim et al.167 added 10 wt% of expanded graphite to a polystyrene matrix. The graphite was initially intercalated with potassium metal of varying molar ratios to expand the distance between graphene layers. They provided XRD and TEM analysis. Both suggested that the graphite remained aggregated in groups of several graphene layers in
49 thickness. Regardless of the poor dispersion, they reported over 100-fold improvement in conductivity.
She et al.168 employed an unsaturated polyester resin to modify expanded graphite.
The modified graphite was then powdered into small particles to facilitate dispersion in a high density polyethylene matrix. In this case, percolation was reached at 5.7 wt% of the modified expanded graphite, whereas 22 wt% of the conventional expanded graphite was required to reach percolation.
Zheng et al.169 dispersed expanded graphite at 3 wt% and 5 wt% loading. The lower loading was “treated” graphite, whereas the higher, 5 wt% loading, was “untreated” graphite. The materials were mixed with high density polyethylene (HDPE) in a Haake mixer or twin-screw extruder to disperse the graphite. The percolation threshold was reached above 2 wt% for both materials, however the conductivity of the composite containing the 3 wt% “treated” graphite was four orders of magnitude higher than that of the “untreated” graphite at a higher loading. This difference was attributed to variation in the levels of dispersion for each material. The graphite aggregate size was analyzed by
SEM which showed that the “untreated” graphite had an average sheet size (lateral dimensions) of 0.4 mm, and aggregate thicknesses approximating 5µm. The treated graphite was much better dispersed. The lateral dimensions averaged 20µm and the average thickness was 100 nm. It should be noted that both nanocomposites had lower strain to failure than neat HDPE, however the “untreated” graphite nanocomposite displayed the lower strain to failure.
Treating graphite by chemical or physical means is a benefit in that such techniques may improve dispersion of the nanoparticles throughout the matrix. However, it should
50 be recognized that such treatments may affect the transport properties of the
nanocomposite. For example, chemical functionalization or heavy oxidation disrupts the
sp2 hybridization of the graphene sheets, thereby negatively affecting conductivity.
Physical treatments such as pulverization or ultra sonication reduce graphite sheet dimensions, thereby affecting the ability to form a continuous conductive network. Such treatments however are often necessary to achieve the desired level of graphene dispersion; therefore a trade-off may be necessary in order to optimize dispersion and properties. As with the layered silicates, many properties, especially mechanical, vary from polymer to polymer and are highly dependent on graphite loading and dispersion.
Mechanical properties are also commonly reported, however these benefits from graphite fillers are less reliable, as compared to the electrical conductivity. While most of the available literature reports at least a modest increase in conductivity with the dispersion of graphite, the mechanical performance can, in some cases, be compromised.
The mechanical properties in general tend to depend on particle size and dispersion.
Drzal, et al.170 dispersed expanded graphite into a high temperature, thermosetting, polyimide (PETI-5). The graphite particle size was reduced by sonication, then further reduced by vibratory ball-milling. They reported a significant increase in storage modulus as graphite concentration increased, and particle size decreased.
Yasmin and Daniel,118 however, observed negative effects on the mechanical
performance of graphite nanocomposites from continued sonication. They reported
reinforcing epoxy with 1-2 wt% of EG by direct mixing, sonication, shear mixing, and a
combination of shear and sonication. They found that combining shear mixing with
51 sonication produced enhanced resin elastic modulus and tensile strength. However, as
sonication times were increased, the mechanical properties degraded.
3.15 Conclusion of Expanded Graphite Nanocomposites
Research in the area of graphite nanoparticles is limited relative to other nanomaterial
work to date. However, in recent years research in the area has exploded has the many
benefits of graphite have become recognized. The material offers the strength, stiffness,
and morphology of layered silicate clays, and the conductivity of carbon nanotubes. The primary barrier to rapid advance in the field had been the difficulty in dispersing the particles. The graphene surface is inherently inert and far more difficult to functionalize than that of layered silicate clays. However, as expertise in oxidation and surface modification continues to develop, the field will grow and undoubtedly offer new materials with interesting properties.
3.16 Summary
The background broadly describes the considerations made when organically modifying and dispersing layered nanomaterials into a polymeric matrix. The next chapters focus on optimizing the chemical modification of either layered silicate clay or nano-graphite for dispersion in thermosetting resin systems.
52 CHAPTER IV
EFFECTS OF SILICATE – MODIFIER ORIENTATION ON NANOCOMPOSITE
INTERFACE AND PROPERTIES IN A HIGH TEMPERATURE POLYIMIDE
MATRIX
4.1 Introduction
The durability and reliability of materials used in aerospace components is a critical
concern. Among the material requirements are a high glass transition temperature (Tg), high temperature stability in a variety of environments, and good mechanical properties over a wide range of temperatures.48 PMR-type polyimides have met the performance requirements listed above, with the added benefit of simplified processing. PMR-15
synthesis proceeds in a low boiling solvent, such as methanol, rather than traditional
polyimide solvents such as N-methyl pyrrolidinone or dimethylacetamide. Processing
from a solvent such as methanol eliminates the need to apply vacuum during processing
for solvent removal. It also reduces voids in the cured resin that commonly result due to
evaporation of the higher boiling solvents during processing.
PMR-15 is commercially available and prepared in two stages from three monomer reactants: 2-carbomethoxy-3-carboxy-5-norbornene (nadic ester, NE), 4,4’-
methylenedianiline (MDA), and the dimethyl ester of 3,3’,4,4’-
benzophenonetetracarboxylic acid (BTDE). Curing under heat and pressure results in a
53 highly crosslinked network structure.49 The PMR-15 synthesis is outlined in Figure 4.1.
There has been a significant amount of research aimed at increasing the thermal oxidative
stability (TOS) of PMR-15 by altering the structure of the dianhydride,50 the diamine,49,171-172 or the end-cap51. Modification of the chemical structure has afforded improved material properties, but often at the expense of processability. Dispersion of a layered silicate into the polymer matrix provides an alternative means of increasing TOS.
Layered silicates have quickly become recognized as useful fillers in polymer matrix composites. Their platelet morphology and high aspect ratio results in greatly improved thermal,173-174 mechanical,103,175-176 and barrier properties.177-178 Because the interface is a
dominant factor in the nanocomposite performance, care should be taken to address the
many variables that exist in modification of the silicate surface. The strength of the
interface may be tailored by choosing an organic modifier that forms a strong interaction
or bond with the matrix. The concentration of the modifier may be controlled by
choosing a silicate with a greater or lesser cation exchange capacity. The amount of the
modifier tethered to the silicate surface has a profound effect on the orientation of the organic molecule, and the subsequent material properties. For example, Vaia et al.82 report that a high surfactant concentration between layers will force alkyl ammonium chains to extend perpendicular to the silicate layers, whereas a low concentration will allow the chains to extend diagonally from the basal surface.
In order for the silicate to perform in a high temperature polymer matrix, the organic modification chosen for the clay surface must not degrade during processing or at the intended use temperature. Degradation of the modifier can result in residual organics in the matrix negatively impact performance. Furthermore, degradation of the organic
54
O O O O H CO OCH 3 3 H NNCH H+ OCH3 HO OH + 2 2 2 OH O O MDA O BTDE (n+1 moles) NE (n moles) (2 moles) MeOH
o 120-232 C
O O O O O N[ CH2 N NC] H2 N O O O o 2.087 O 315 C
O O N N O CH O 2 CH2 O O N N O O
Figure 4.1. Synthesis of PMR-15
55 modifier may result in re-aggregation of the layered silicate.85 To avoid such degradation in high temperature polyimides, researchers have modified clay with aromatic diamines that are not only thermally stable, but allow for covalent bonding between the modified clay and the matrix.86,87 Such modifiers were used in this study, but proved ineffective.
Therefore, a novel modification system was identified where both aromatic and aliphatic compounds were tethered to the silicate surface.
The focus of this chapter is to examine the effects of the silicate modification on the overall nanocomposite properties. Specifically the relationship between the protonated diamine chain length and orientation, including its effect on oligomer melt viscosity and polymer crosslink density is investigated.
4.2 Experimental details
Experimental details of the synthesis and characterization of PMR-15/ silicate nanocomposites are described.
4.2.1 Materials
Bentonite (Bentolite-H, cation exchange capacity (CEC) = 90 meq/100g) was received from Southern Clay Products and will be referred to as SCP. Montmorillonite (PGV- grade, CEC = 145 meq/100g) was received from Nanocor and will be referred to as PGV.
Ion exchange amines: p-phenylene diamine (pPDA), methylene dianiline (MDA), octylamine (C8), and dodecylamine (C12) were purchased from Aldrich. 4,4’(1,4- phenylene-bismethylene) bisaniline (BAX), was purchased from the Maverick
Corporation and 4,4’-bis(4-aminobenzyl)diphenylmethane (ABDM) was synthesized by
Prof. David Klopotek at St. Norbert College.49 The chemical structures of the ion
56 exchange diamines are shown in Figure 4.2. NE and 3,3’,4,4’-
benzophenonetetracarboxylic acid dianhydride (BTDA) were purchased from Christev.
All materials were used as received.
4.2.2 Synthetic Procedures
PMR-15 Synthesis: BTDE was prepared by refluxing BTDA in an amount of
methanol calculated to yield a solution containing 50 wt% solids. The reflux was carried
out for two hours after the solid BTDA had dissolved. Dissolution indicated the
formation of the acid ester (BTDE) species. The solution was used immediately for resin
synthesis, or refrigerated for later use.
PMR-15 resins were fabricated in several steps. The three monomers (NE, MDA, and
BTDE) were dissolved in methanol followed by solvent evaporation, on a hot plate, at
60o to 70oC. B-staging the mixture at 204o to 232oC in an air circulating oven produced a low molecular weight imide oligomer. The oligomer was then cured in a mold at 315oC under 2355 psi to produce the crosslinked polymer. The polymer was post cured in an air circulating oven for 16 hours at 315oC to further crosslinking. The average number of
imide rings was kept constant by using a stoichometry of 2NE/ (n+1)MDA/ nBTDE
(n=2.087) corresponding to an average molecular weight of 1500.179
4.2.3 Nanocomposite Synthesis
Ion exchange of the interlayer cations of SCP and PGV clays with the protonated forms of the listed amines was performed by dissolving the amine (5mmol) in 450 mL of a 0.005M aqueous HCl solution at 60oC. In the case where the diamine did not dissolve in aqueous solution, ethanol (25 – 50 mL) was added. The silicate (5g) was dispersed in the solution and the resultant mixture was stirred at 60oC for three hours. The solution
57
pPDA H2N NH2
H2NNCH2 H2 MDA
BAX H2NCH2 CH2 NH2
ABDM H2NCH2 CH2 CH2 NH2
Figure 4.2. Chemical structure of ion exchange diamines
58 was filtered and the clay was washed thoroughly with distilled water, which was heated
to 60oC. Based on literature recommendations for maximizing ion exchange, the
procedure was repeated for a total of three exchange reactions.40 The silicate was then dried overnight in a vacuum oven at 100oC. The clays modified with various amines will
be identified by the silicate and the ion exchange amine, for example, SCP-amine or
PGV-amine for the organically modified forms of SCP and PGV respectively.
The procedure used to prepare PMR-15 nanocomposites was identical to that of the
neat resin, except that 2 wt% of organically modified silicate was added to the methanol
solution of the monomers. All nanocomposites were prepared with 2 wt% organically
modified clay, unless otherwise stated.
4.2.4 Characterization
X-ray diffraction (XRD) patterns were obtained using a Philips XRG 3100 X-ray
diffractometer with Ni-filtered CuKα radiation. Ion exchanged clays, B-staged PMR-15
nanocomposites, and post cured PMR-15 nanocomposites were ground into powder and
the XRD data was recorded in the range of 2θ = 2o to 32o. An increase in the basal layer spacing, which was determined from a shift in the (001) peak position, indicated ion exchange or polymer intercalation between the silicate layers. Disappearance of the
(001) peak suggested an exfoliated morphology.
Transmission electron microscopy (TEM) specimens were prepared by microtoming sections of post cured PMR-15 nanocomposites, 20 to 70 nm thick, and floating the sections onto Cu grids. Micrographs were obtained with a Philips CM 200, using an acceleration voltage of 200 kV. The TEM images shown throughout this work are
59 representative of the dispersion observed throughout several sections, taken from various
regions, of each nanocomposite sample.
A Perkin Elmer High Pressure Differential Scanning Calorimeter (HP- DSC) was used
to evaluate the crosslink reaction enthalpy of the imide oligomers after B-staging. The
oligomer (8-12 mg) was weighed into a sealed aluminum DSC pan. The tests were
performed at 200 psi under nitrogen and the temperature was ramped from 25oC to 450oC at a rate of 10oC/min.
The glass transition temperature of post-cured PMR-15 nanocomposites was measured in air on a TA Instruments Dynamic Mechanical Analyzer (DMA). The average sample size was 17.4 mm x 5.0 mm x 1.7 mm. The analysis was performed with a ramp rate of
2.5oC/min and 10 µm amplitude.
Isothermal aging of the PMR-15 nanocomposites was performed to determine the
TOS. Post-cured samples were cut into 1.02 cm by 0.64 cm coupons and placed in an air circulating oven at 288oC for 1000 hours. The weight loss was measured at regular intervals by removing the coupons from the oven, allowing them to cool to room
temperature, and weighing the sample.
The melt viscosities of B-staged nanocomposites were compared by measurement of the materials’ squeeze flow index (SFI).180 Comparison of the SFI provided a relative indication of the melt flow. The value of the squeeze flow index was determined by initially cold pressing the powders (0.5g) into a pellet using a KBr pellet mold, as used in infrared spectroscopy. The pellet was placed between Kapton sheets in a 25 ton press and held at 315oC and 200 psi for two minutes. This process resulted in a circular “splat”
of resin, due to the resin flow. The diameter of the “splat” was measured in several
60 places and the average was used to calculate the area of resin flow. This area was
recorded as the squeeze flow index. Each sample was run in triplicate.
4.3 Results and Discussion
The aromatic diamines chosen as the ion exchange material were similar in chemical
structure to, or the same as, the diamine found in PMR-15. Tethering one end of the
diamine to the silicate leaves the second free for reaction with BTDE during PMR-15
synthesis. Wei et al.173 demonstrated this as a viable method of achieving irreversible
swelling of the silicate in a thermoplastic polyimide matrix.
4.3.1 Ion Exchanged Clays
The clays used in this study differ in CEC (PGV =145 meq/100g, SCP = 90
meq/100g) therefore fewer diamines are exchanged with SCP. A comparison of the
amount of exchanged diamine was made by thermogravimetric analysis (TGA), and the
data is shown in Table 4.1.181 Measurement of the weight loss on heating the untreated clays to 800oC provides the content of adsorbed and structural water within each clay.
Heating the organically modified clays to 800oC, and subtracting out the weight loss
corresponding to water content, gives an indication of the amount of diamine exchanged
into each clay. TGA measurements suggest that an average of 7% to 14% more diamine,
by weight, was exchanged into PGV. This is consistent with the difference in the CEC of
each clay. The exception is ion exchange of pPDA, where approximately 2% by weight
more diamine is exchanged into SCP.
The d-spacing of the organically modified silicates is calculated from XRD data.
The XRD patterns of SCP and PGV clays exchanged with each diamine are shown in
61 o Figures 4.3a and 4.3b. The d001 peak of untreated SCP appears at 2θ = 7.07 (d001 = 1.25
o nm). Ion exchange with pPDA shifts the peak to 2θ = 5.97 (d001 = 1.48 nm). After this initial increase in gallery height, there is little change in d-spacing as the number of phenylene linkages in the diamine increases.
The values of the clay d-spacing are listed in Table 4.2, including those of PGV exchanged with the same diamines. Unlike SCP, ion exchange in PGV produces a monotonic increase of the d-spacing with increasing diamine length. The d001 spacing is
calculated from the peak at half width, representing the d-spacing of the greatest number
of stacked silicate layers. There is some breadth to the diffraction peaks, indicating that
the spacing between layers of organically modified sheets varies within the larger
aggregates. It should be noted that the observed increase in d-spacing is small, the
smallest being that of PGV-pPDA. This is explained by the fact that the d-spacing listed
for untreated PGV, 1.22 nm, includes 0.94 nm of the silicate sheet thickness. Therefore,
the actual empty space between silicate layers in untreated PGV is 0.28 nm (1.22 nm –
0.94 nm). Following ion exchange with pPDA, the d-spacing obtained from XRD is
calculated as 1.28 nm, meaning that the space between layers has increased to 0.34 nm.
Such spacing is sufficient to accommodated an aromatic ring lying flat between layers.207
The observed variation in d-spacing between organically modified SCP and PGV
suggests that the exchange diamines may adopt a separate configuration in each silicate.
Organically modified SCP shows little change in gallery height on increasing the length
of the diamine, suggesting that the longer diamines may adopt a folded conformation, as
illustrated in Figure 4.4a, representing one potential orientation which would account for
62
Table 4.1. TGA data indicating the quantity of diamine present in clays
Ion Exchange SCP clay % Wt% PGV clay % Wt.% Diamine on Clay Residue at modifier in Residue at modifier in 800oC clay 800oC clay Un-modified clay 94% ------95% ------pPDA 81% 13% 84% 11% MDA 88% 6% 82% 13% BAX 75% 19% 72% 23% ABDM 68% 26% 55% 40%
63
(a)
ABDM ity
s BAX n e t
In MDA pPDA
SCP 246810 2 Theta
(b)
ABDM ity
s BAX n
te MDA In pPDA PGV
246810 2 Theta
Figures 4.3. (a) XRD patterns of unmodified and organically modified SCP, and (b) XRD
patterns of unmodified and organically modified PGV
64
Table 4.2. Increase in silicate gallery height after ion exchange of aromatic diamines in
SCP and PGV
Silicate d001 interlayer spacing (nm) d001 interlayer spacing (nm) Modification (SCP) (PGV) Untreated 1.25 (+/- 0.14 nm) 1.22 (+/- 0.15 nm) pPDA 1.48 (+/- 0.23 nm) 1.28 (+/- 0.18 nm) MDA 1.52 (+/- 0.23 nm) 1.39 (+/- 0.16 nm) BAX 1.47 (+/- 0.18 nm) 1.46 (+/- 0.25 nm) ABDM 1.52 (+/- 0.18 nm) 1.50 (+/- 0.22 nm)
65 the small increase observed in silicate gallery spacing. The monotonic increase observed
for ion exchange of PGV suggests the aromatic diamines are not folding, but may extend
away from the silicate surface. However, the increase in gallery spacing is small and
therefore the chains are likely not perpendicular to the silicate, but may lie at an inclined
angle with respect to the silicate surface, as illustrated in Figure 4.4b. Such an
arrangement allows for π−π interactions between the aromatic rings, which will favor this
extended arrangement.182
Such variation in packing and orientation has been demonstrated previously.
Giannelis et al. have shown that the conformation of alkyl ammonium cations within a silicate gallery is dependent on the chain length and layer charge density of the clay. 82,183
Such arrangements were elucidated based on XRD and FTIR data. XRD data alone has been used to identify the arrangement of single ring aromatic compounds as lying flat between clay layers. The observed increase in d-spacing approximated 0.34 nm.207
Due to the low cation exchange capacity of SCP, there is a large surface area available
to the chains so that they may fold back toward the silicate sheet. As the CEC increases,
the chain packing density of the aromatic exchange diamines increases and the available
surface area decreases. This prevents folding and forces the diamine away from the
silicate surface, resulting in the observed dependence of PGV d-spacing on the aromatic
chain length.
4.3.2 Orientation Characterization by Infrared Spectroscopy
The orientation differences among various alkylammonium ions, within silicate
galleries, has been verified by FTIR.82,83 This technique is useful in these cases as the
conformational freedom of the methylene chain is restricted within the clay galleries.
66
(a) + + H3N H3N NH2
NH2 pPDA BAX
+ + H3N H3N NH2
NH2 MDA MMM
(b) ++ H3N H3N
NH2 NH2
Figure 4.4. (a) Proposed orientation of protonated aromatic diamine in SCP; low cation
exchange capacity clay, and (b) orientation of protonated ABDM in PGV; high cation
exchange capacity clay
67 Unfortunately, the amine series used in this work lacks the conformational freedom
available to the aliphatic chains. However, FTIR spectra of each organically modified
clay were collected and compared.
The images in Figure 4.5 illustrate several differences in the FTIR spectra of the clay
BAX samples. The spectra of these samples are convenient for comparison as the
molecular chain length is long enough to affect molecular packing and conformation. A
significant difference is noted in the shift of the vibrational stretching of the aromatic
C=C bands.184 The aromatic C=C stretching band of SCP-BAX, where it is predicted
− that the chains fold and remain in close contact with the silicate clay, appears at υ =
− 1521.1 cm-1. This band shifts to a lower frequency, υ = 1516.2 cm-1, within the PGV-
BAX material. This is consistent with the proposed orientation. The closer contact
between the clay surface of SCP and the aromatic moieties of BAX, results in an
increased force constant relative to the PGV-BAX sample. Therefore the C=C stretching
frequency is increased within SCP-BAX. A similar shift was observed throughout the
series of modified clays, as shown in Table 4.3. However, based on this data, it is
difficult to draw conclusions of the aromatic conformation, as this particular vibrational
mode is complex, and the band behavior is not as well understood as that of the aliphatic
series.208
4.3.3 Characterization of Silicate Dispersion in PMR-15
The XRD patterns of B-staged PMR-15/ SCP and PGV nanocomposites show a shift
in the d001 peak position, indicating oligomer intercalation. Curing the intercalated
68 3404.7 cm-1
2916.7 cm-1 1521.1 cm-1 3294 cm-1 3336 cm-1
3624.3 cm-1
2921.6 cm-1 3394 cm-1 3170 cm-1
1516.2 cm-1 3624.3 cm-1
Figure 4.5. FTIR spectra of SCP-BAX (upper) and PGV-BAX (lower)
69
Table 4.3 FTIR bands of the ion exchanged clays
Modifier SCP clay (Low CEC) PGV clay (High CEC) Frequency (cm-1) Frequency (cm-1) pPDA 3620 3612 1506 1501 MDA 3620 3612 1510 1505 BAX 3624 3624 3404 3394 3336 3170 3294 2917 2921 1521 1516 ABDM 3616 3614 3402 3334 3208 3009 2912 2905 1513 1509
70 oligomer results in further separation of the silicate layers, and disruption of the stacking registry; as suggested by the absence of a diffraction peak. The exception is PMR-15/
PGV-ABDM where the presence of the diffraction peak indicates the intercalated morphology is maintained on curing.
The level of silicate dispersion within the polymer was confirmed by TEM. Figures
4.6a-e are representative photomicrographs showing the dispersion of PGV-MDA within
several different regions of the PMR-15 matrix. Figures 4.7a-b illustrate the dispersion
of PGV-BAX, and Figure 4.7c shows a very similar level of dispersion for SCP-BAX in
PMR-15. It is generally difficult to fully exfoliate an organically modified silicate clay
into a PMR-15 matrix.86 As noted by Gintert et. al.,86 Such nanocomposites are prepared in solution; therefore exfoliation must occur in the absence of external shear forces. As such, silicate layer dispersion is dependent on intragallery crosslinking to force the layer separation. The addition of aromatic diamine modified clays, both SCP and PGV, to
PMR-15 results in a mixture of intercalated and exfoliated nanocomposite morphologies.
However, it was observed in this work that differences in the proposed orientation of the ion exchange diamine translated into processing and property characteristics unique to each silicate-amine system.
4.3.4 Processing PMR-15 Nanocomposites
The squeeze flow index (SFI) provides a comparison of the melt viscosity between
samples. Pressing the B-staged powder between Kapton sheets results in a circular
“splat” of the melted oligomer. A “splat” with a smaller area indicates a high melt
viscosity, relative to a low viscosity matrix which yields a larger area of melted resin.
The data shown in Table 4.4 reveals an increase in the melt viscosity of SCP
71 (a) (b)
(c) (d)
(e)
Figures 4.6.(a-e) Representative TEM images of PGV-MDA dispersed in PMR-15
72 (a) (b)
(c)
Figure 4.7.(a-b) Representative TEM images showing the level of dispersion of PGV-
BAX in PMR-15 resin, and (c) dispersion of SCP-BAX in PMR-15 resin
73 nanocomposites compared to the neat resin. Nanocomposites prepared using PGV also show an increase in melt viscosity for clay that has been ion exchanged with pPDA and
MDA. However, nanocomposites prepared with PGV-BAX or PGV-ABDM show an identical or an increased SFI, suggesting a similar or even lower oligomer melt viscosity, as compared to PMR-15.
The energy released during endcap crosslinking was measured by HP-DSC for B- staged nanocomposites and is listed in Table 4.4. The data in Tables 4.4a and 4.4b suggest a correlation between the oligomer melt viscosity and the crosslinking enthalpy.
HP-DSC thermograms of B-staged SCP nanocomposites show a decrease in the amount of energy released during the crosslinking reaction, compared to the neat resin, with little dependence on the length of the ion exchange diamine. The melt viscosity of the SCP nanocomposites is increased compared to the neat resin, again with no dependence on the modifier length. In the B-staged PGV nanocomposites series, the nanocomposite melt viscosity decreases on increasing modifier chain length. Corresponding to the reduction in viscosity is an observed increase in the crosslinking exotherm, compared to neat PMR-
15, with the exception of PGV-pPDA.
It has been demonstrated in polymeric samples that the relaxation of polymer chains either tethered to or in close proximity to a host surface such as clay are dramatically altered relative to chains unaffected by a rigid surface.79 A similar effect may occur in these systems, in which the motion of the PMR-15 oligomers may be restricted within the silicate gallery. Inhibition of chain motion can slow oligomer rearrangement to conformations that are energetically favorable for crosslinking to occur through the NE endcap, thus decreasing the crosslink density. With PGV nanocomposites, the diamine
74
Tables 4.4.(a-b) Squeeze flow index and crosslinking enthalpy of neat PMR-15 and
PMR-15 nanocomposites
(a) PGV (CEC = 145 meq/100g) Sample (clay SFI (cm2) DSC (mcal/mg) modification) PMR-15 93 +/- 0.4 14* PPDA 68 +/- 0.1 12* MDA 82 +/- 0.3 14* BAX 93 +/- 0.1 15* ABDM 110 +/- 0.3 15*
(b) SCP (CEC = 90 meq/100g) Sample (clay SFI (cm2) DSC modification) (mcal/mg) PMR-15 93 +/- 3 14 +/- 1 PPDA 77 +/- 7 12 +/- 1 MDA 78 +/- 8 9 +/- 2 BAX 82 +/- 10 9 +/- 1 ABDM 74 +/- 6 10 +/- 2
* Data obtained from one sample
75 packing density is increased and, based on the monotonic increase in d-spacing, the
tethered diamines appear to radiate away from the silicate surface. In these samples, with the exception of PGV-pPDA and PGV-MDA nanocomposites, the oligomer melt viscosity of the nanocomposite is equal to, or lower than that of neat PMR-15. Therefore we observe a crosslink density similar to that of the neat resin. It should be noted that the reported error in squeeze flow index is based on batch to batch related error in the SCP samples, whereas the PGV error is a measure of flow differences of separately evaluated samples of the same batch. The values of DSC exotherm for the PGV material were taken from 1 sample, therefore error based on sample to sample variation could not be calculated. However, batch to batch variation similar to that observed in the DSC exotherm observed of the SCP nanocomposites would be expected in the PGV nanocomposites as well.
4.3.5 Thermal Oxidative Stability of PMR-15 Nanocomposites
An anticipated advantage of polymer-clay nanocomposites is an improvement in the
polymer thermal stability over that of the neat resin. However, PMR-15 is a fairly stable
thermosetting polymer. Disruption of its highly crosslinked structure could lead to a
reduction in the TOS. As illustrated in Figure 4.8, there is a decrease in TOS, i.e. an
increase in oxidative weight loss, in all SCP and PGV nanocomposites, with the
exception of PGV-ABDM nanocomposites. This material has a weight loss
approximately equal to PMR-15 and is the only nanocomposite with an oligomer melt
viscosity lower than that of the neat resin. The observed reduction in TOS may therefore
be due to incomplete oligomer crosslinking in the nanocomposite samples due to
increased melt viscosity.
76 To support the influence of oligomer melt viscosity on the nanocomposite TOS, the melt viscosity of the B-staged PGV-MDA nanocomposite was decreased and the TOS measured. To decrease the melt viscosity of B-staged PMR-15/ PGV-MDA, PGV was ion exchanged with 50 mol% MDA and 50 mol% protonated dodecylamine, C12. The longer chain length of C12 (15.0 nm) compared to MDA (10.7 nm) apparently inhibits the oligomer from coming within close contact with the silicate, but still allows reaction between the tethered diamine and BTDE. For comparison, PGV was also modified with
50 mol% MDA and 50 mol% protonated octylamine, C8, where the alkyl amine chain length (10.0 nm) was similar to that of MDA. The conformations and dimensions of the
MDA, C12, and C8 structures are shown in Figures 4.9a-c, as calculated by HyperChem
Software.
From XRD data, Figure 4.10, the d-spacing of PGV(MDA-C12) increased to 1.47 nm,
compared to 1.39 nm for PGV-MDA. Co-exchange of MDA and C8, however, offered
no increase in d-spacing. PGV(MDA-C8) had a d-spacing of 1.40 nm, little change from
that of the PGV-MDA d-spacing. The increase in layer separation provided by C12
facilitates dispersion of the PGV(MDA-C12) clay in the PMR-15 matrix.
Modification of PGV with both MDA and C12 offers greatly improved silicate
dispersion within the matrix. Representative TEM images of this system are shown in
Figure 4.11. Co-exchange of a long alkyl chain in the presence of an aromatic diamine
may influence the orientation of the aromatic rings by disrupting the π−π interactions,
thus increasing diamine mobility.79 It is observed that the flexible chain also aids in reducing melt viscosity. The reduction in melt viscosity may be due to the alkyl ammonium chain allowing increased molecular motion near the oligomer-silicate
77
7
s 6.6 PMR-15 Los 6.2 pPDA ght i
e MDA W t 5.8 BAX n e
c ABDM r
Pe 5.4
5 SCP PGV
Figure 4.8. Weight loss of neat PMR-15 and PMR-15 nanocomposites on heating at
288oC for 1000 hours
78
(a) Molecular length of methylene dianiline = 10.7 nm
(b) Molecular length of dodecylamine = 15.0 nm
(c) Molecular length of octylamine = 10.0 nm
Figure 4.9. Molecular lengths of ion exchange amines
79 (a) y
it PGV-C12 s n
e PGV-(MDA-C12) t In PGV-MDA PGV
2 5 8 11141720 2 Theta
(b)
PGV-C8 ity
s PGV-(MDA-C8) n
te PGV-MDA In PGV
2 5 8 11141720 2 Theta
Figure 4.10.(a-b) XRD patterns of organically modified PGV clay
80
Figure 4.11. Representative TEM images of PGV(MDA-C12) dispersed in PMR-15
81 interface. Table 4.5 lists the melt viscosity and crosslinking enthalpy for B-staged
nanocomposites prepared with the mixed amines. The data demonstrates that the
oligomer melt viscosity and crosslink density can be manipulated by introduction of a
long chain alkylamine onto the silicate. Figure 4.12a shows the TOS of the PGV-(MDA-
C12) nanocomposites of increasing clay concentration compared to neat PMR-15. Up to
a 25% decrease in weight loss was observed on aging the samples in air for 1000 hours at
288oC. This is a significant improvement compared to the thermal oxidative stability of the PGV-MDA nanocomposites. For comparison, the TOS data of nanocomposites containing 2 wt% PGV-MDA, PGV-(MDA-C8) and PGV-C8 is included in 4.12b. It is clear that the benefits observed from the co-exchange of MDA and C12 are not observed
when the shorter alkyl ammonium ion, C8, serves as the co-exchange ion.
Unfortunately, the low molecular weight alkylamine can act as a plasticizer in the
o polyimide matrix, reducing Tg. The Tg of neat PMR-15 as measured by DMA is 347 C.
Addition of 2 wt% PGV-MDA results in no change of Tg, where addition of 2 wt% PGV-
o o C12 lowers the Tg to 330 C. Co-exchange of MDA and C12 results in a Tg of 345 C, which is not a significant change from that of the neat PMR-15. A comparison of the
PMR-15 nanocomposite Tg values is shown in Table 4.6.
4.3.6 Polymer Matrix Composites
Carbon fabric reinforced composites with a PMR-15/ silicate nanocomposite matrix were prepared. A side by side evaluation of the TOS of composites prepared with the nanocomposite matrix as well as the neat resin matrix was conducted. Prepreg was
prepared by brush application of the PMR-15 monomer solution onto T650-35 carbon fabric, to give a final fiber content of 60 wt%. The silicate used in the nanocomposite
82
Table 4.5. Squeeze flow index and crosslinking enthalpy of neat PMR-15 and PMR-15
nanocomposites
Ion exchange amine in Squeeze Flow DSC (mcal/mg) PGV Index (cm2) + + PMR-15 (neat resin) 93 /- 0.4 14 /- 1 + MDA 82 /- 0.3 14* + C8 87 /- 5 13* + C8/MDA 90 /- 2 14* + C12 115 /- 3 15* + C12/MDA (50:50) 125 /- 5 17*
* Data obtained from one sample
83
8 s 7.6 Los 0% silicate 7.2 ght
i 1% silicate e 6.8 3% silicate W t n 6.4 5% silicate e c r 6 7% silicate e P 5.6
1000 hours (288oC)
8 s 7 Los PMR-15 ght i MDA e 6
W C8 t n
e MDA-C8 c
r 5 e P 4
1000 hours (288oC)
Figure 4.12.(a-b) Weight loss of neat PMR-15 and PMR-15 nanocomposites on heating at
288oC for 1000 hours, where silicate is modified with a diamine and an aliphatic amine
84
Table 4.6. Tg of post cured PMR-15 and PMR15 nanocomposites.
o o Sample (clay Tg ( C) Tg ( C) modification) SCP PGV PMR-15 347 +/- 2 347 +/- 2 PPDA 343 +/- 4 348 +/- 4 MDA 342 +/- 4 350 +/- 3 BAX 342 +/- 4 343 +/- 1 ABDM 342 +/- 4 340 +/- 3 C12 n/a 330 +/- 2 MDA-C12 n/a 345 +/- 2
85 matrix composites was either PGV-ABDM or PGV-(MDA-C12). The prepreg sheets
were cut into eight, 10.2 cm by 10.2 cm, plies and placed in a metal mold. The mold was
initially heated to 232oC to imidize the monomers. The mold temperature was then
raised to 315oC, and the matrix was cured with application of 500 psi for 2 hours. The composites were post cured in an air circulating oven at 315oC for 16 hours. The TOS was evaluated by aging the specimens (2.54 cm by 1.27 cm) in an air circulating oven heated to 288oC, and monitoring weight loss for 1000 hours.
The amount of resin in each composite was determined by acid digestion. The composite quality and void content were evaluated by C-scan. Both tests indicated that the neat resin matrix and nanocomposite matrix composites had a similar resin content of
40 wt%, and all composites were of comparable quality. The TOS data presented in
Figure 4.13 shows an average decrease of 20% (average of three samples) in the weight loss of the nanocomposite matrix composites compared to the neat resin matrix composite after aging for 1000 hours. The main route to degradation of PMR-15 composites is through oxidation of the surface layer followed by microcracking in the matrix. Microcracking allows permeation of oxygen into the bulk of the sample, furthering oxidative degradation.185 Dispersion of the silicate in the matrix inhibits the permeation of oxygen through the bulk, enhancing the TOS of the composite.86
4.4 Conclusions
The orientation and packing density of the aromatic diamine used to modify SCP and
PGV clays affects the oligomer melt viscosity of a PMR-15/clay nanocomposite. An increase in oligomer melt viscosity is accompanied by a decrease in the oligomer
86 crosslinking enthalpy, as well as a decrease in the TOS of the cured nanocomposite. The
oligomer melt viscosity can be modified by co-exchange of an aromatic diamine and a
long chain alkyl amine on the clay. Incorporation of the novel PGV/( MDA-C12) co-
modifier into the PMR-15 matrix has little affect on Tg. In addition, the TOS of the nanocomposite is improved compared to the neat resin or nanocomposites prepared with clay exchanged only with the aromatic diamine. Dispersion of PGV-ABDM or
PGV/(C12-MDA) in the matrix of a PMR-15/ carbon fiber composite results in a significant increase in the TOS of the composite. Thus, this work demonstrated that the use of nanocomposites as matrix resins is a viable approach to the development of polymer matrix composites with improved high temperature stability. However it is evident that the difference in interaction of the clay with both alkyl and aromatic components becomes relevant when deciding to add the clay to a multicomponent system. This is addressed further in the next chapter.
87
3 PMR-15 PGV(ABDM) 2.5 PGV(MDA-C12) 2 Loss
ght 1.5 i e 1 W
% 0.5 0 1000 hours in air at 288oC
Figure 4.13. Weight loss of PMR-15 matrix and PMR-15 nanocomposite matrix carbon
fiber composites on aging at 288oC for 1000 hours, where the PGV modifications are
ABDM and 50:50 MDA/C12
88
CHAPTER V
EFFECT OF INTERFACE ON SILICATE DISPERSION IN AN EPOXY BLEND
5.1. Background
It is envisioned that several of the next generation space vehicles, including the Ares
V Launch Vehicle, will contain integral liquid hydrogen (LH2) and liquid oxygen (LOX) cryogenic fuel tanks. These tanks will not only function to contain fuel, but must also perform as load-carrying structures for the vehicle.186 As such, they must be able to withstand extreme flight loads and operating temperatures without loss of cryogenic fuel due to microcracking or delamination.
Traditionally, metallic, tanks have been used for housing cryogenic fluids. The advantages of such tanks include high strength and stiffness, and low permeability.
While it appears that the replacement of traditional metallic cryogenic fuel tanks with
PMC tanks may lead to significant weight reductions, there are also critical inherent
disadvantages to composites which reduce their applicability as cryogen storage
materials.
A main concern with these materials is matrix microcracking at cryogenic
temperatures.187 Microcracking allows permeation of the cryogenic fluid into the bulk
polymer matrix which can result in catastrophic failure. Microcracking occurs on
thermal cycling of a composite due to the mismatch in thermal expansion characteristics
89 of the matrix resin and the carbon fiber reinforcement. Such mismatch leads to stress
build up and ultimately results in microcracking. Reducing the coefficient of thermal
expansion (CTE) of the matrix to better match that of the fiber reinforcement is one
means by which to reduce the stress concentrations at the fiber/ matrix interface, and
thereby reduce microcracking.
Another approach to limit microcracking is enhancing the toughness of the epoxy
matrix, as toughened epoxies tend to resist microcracking better than untoughened
systems.188 Toughening the material offers drawbacks, however, as the mechanical strength of the epoxy is typically decreased. The decrease in strength may be compensated for by dispersion of a filler, such as layered silicate clays, into the resin.
For composite tanks to be a viable replacement to the current metallic tanks, an approach must be taken which reduces the permeability of the composite materials but also enhances the strength and toughness of the composite matrix.
5.2 Introduction
Layered silicate clay has been investigated as a potential filler for PMCs intended for cryogenic fluid storage.189-190 Overall, the mechanical performance of polymer-clay nanocomposites can vary significantly and is often dependent on resin properties, silicate loading, the level of silicate dispersion, and the organic modifier on the clay.40, 108-109 It is generally recognized that a trade off exists between material modulus and ductility, where many nanocomposites exhibit an increased modulus but a reduced ductility relative to the base resin.110 Several researchers have noted that this trend does not hold when tensile load is applied at a temperature greater than the nanocomposite Tg.
90 114-115 Pinnavaia et al. have noted that flexible resin systems, with low Tg, show a much
greater increase in modulus and tensile strength with the addition of an organoclay, than
do rigid systems. Giannelis et al.116-117 provided evidence that clay mobility in a matrix allowed for layer orientation, thereby providing a mechanism for energy dissipation.
Similar studies have shown an increase in nanocomposite toughness in high Tg epoxies blended with rubber tougheners, where property enhancements in these resins far exceed those of the untoughened nanocomposites.191-192
Data presented in the previous chapter illustrated a certain affinity between aromatic
amine salts and layered silicate clay, as compared to those derived from aliphatic amines.
Recently, Giannelis et al.193 reported that phenyl functionalized polydimethylsiloxane
(PDMS) dispersed organically modified silicate to a much greater extent than unfunctionalized PDMS. The driving force was not that interactions between the phenyl moieties and silicate were favorable, just that they were more favorable than the aliphatic
- silicate interactions.
Similarly, Korley et al.107 demonstrated preferential contact between ‘hard’ (aromatic) and ‘soft’ (aliphatic) polyurethane segments, depending on the hydrophilicity of the silicate. They reported that, for either unmodified or hydrophilic clays dispersed in a polyurethane containing a polar, hydrophilic soft block, the soft block dominated the polymer-clay interactions. In this case, the soft block was made of (PEO-PPO-PEO) segments. However, employing a hydrophobic soft block shifted the interaction to the aromatic containing hard section. Such data is relevant in that it demonstrates the ability to control clay location within the nanocomposite by tailoring the specific interfacial
91 interactions. Such data is valuable when designing epoxy blends containing both rigid aromatic and flexible aliphatic components.
While a few initial studies are beginning to show the preference of clay for aromatic over aliphatic components of block copolymers, there are no studies which investigate the influence of such placement on the mechanical performance of the epoxy systems.
Furthermore, there have been no reports in which a pre-swelling technique is used to direct placement of the clay. Pre-swelling is a valid approach to force the less desirable interactions, as it yields the added benefit of improved dispersion. Improved dispersion is the primary reason that researchers have used pre-swelling in the past.59,98-100
The purpose of this chapter is to demonstrate that nanocomposite properties may be
optimized by specifically placing clay nanoparticles in contact with the more mobile region of an epoxy blend. For the case of a toughened epoxy system, the more mobile region would be a rubber toughener. In this case, a reactive diluent is used for toughening. Pre-swelling organically modified clay in the flexible, aliphatic (PPO) monomer, forces the clay to reside in the mobile regions of the blend. Such manipulation optimizes the benefits of the clay to the resin, resulting in significant enhancements to epoxy strength, toughness, and dimensional stability.
5.3 Experimental
Experimental details of the preparation and characterization of epoxy/ silicate
nanocomposites are described.
92 5.3.1 Materials
Epoxy resin, Epon 826, was supplied by Resolution Performance Products. Araldite
DY3601, a polypropylene oxide based epoxide, and Jeffamine D230 curing agent were
supplied by Huntsman Chemicals. The organically modified clay, Cloisite 30B, was
supplied by Southern Clay Products. The structures of Epon 826, DY3601, D230, and
the organic modification of Cloisite 30B are shown respectively in Figure 5.1.
5.3.2 Nanocomposite Preparation
Resin plaques of Epon 826 (aromatic) and DY3601 (aliphatic) epoxy blends were prepared in 90:10, 70:30, and 50:50 equivalent epoxy ratios; with the first number corresponding to the EPON 826 content, and the second number referring to the DY3601 content. Resin plaque preparation at a 70:30 ratio required mixing Epon 826 (18.4 g) and
DY3601 (7.875g) in a jar, followed by stirring 40oC.
Either 2 wt% or 5 wt% of Cloisite 30B was added and the mixture was stirred with a stir bar for 3 hours. The epoxy/clay mixture was cooled and the D230 curing agent (7.5 g) was added. The contents of the jar were poured into a 10.2 cm by 10.2 cm mold. The
resin was degassed at 40oC for 3 hours then cured at 75oC for 2 hours and 125oC for 2 hours.
The above procedure was followed for all sample ratios prepared. The monomer quantities varied as follows: Plaques with a 50:50 blend of the epoxy resins were prepared, using Epon 826 (13.52g), DY3601 (13.52g), and D230 (6.75g), and the 90:10 ratio plaques contained: Epon 826 (22.95g), DY3601 (2.55g), and D230 (8.25g).
Resin plaques containing pre-swollen clay were prepared by sonicating 2 wt% or 5 wt% 30B, where clay concentrations were based on final nanocomposite weight, with 5%
93
Epon 826
O OH O CH2CHCH2 [OOCH2CHCH2]O OCH2CH CH2 n Mn <700, n~1
DY3601 Jeffamine D230 O CH3 O NH2CHCH2 [ OCH2CH ] NH2 H2CCHCH2 [ OCH2CH] OCH2 CH CH2 x n CH CH 3 3 x = 2.6 Mn < 385
Cloisite 30B Modifier
CH2CH2OH + CH3 NT
CH2CH2OH
Where T is Tallow (~65% C18, ~30% C16, and ~5% C14
Figure 5.1. Chemical structures of Epon 826, Araldite DY3601, Jeffamine D230, and the
organic modifier on Cloisite 30B
94 or 10%, respectively, DY3601 epoxy resin. The clay and epoxy mixture was sonicated
for 2 hours to allow for intercalation into the clay galleries. Following sonication, the
swollen clay was added to EPON 826 and any additional DY3601 that would be required.
The mixture was stirred with a stir bar for 3 hours. The epoxy/clay mixture was cooled
and the calculated quantity of D230 curing agent was added. The contents of the jar were poured into a 10.2 cm by 10.2 cm mold. The resin was degassed at 40oC for 3 hours then cured at 75oC for 2 hours and 125oC for 2 hr.
5.3.3 Characterization
XRD and TEM experiments proceeded as previously described in Chapter 4. CTE
was characterized using a TMA 2940 Thermomechanical analyzer. The tests were run at
a ramp of 5 oC/min, using a 2g load. A Perkin Elmer High Pressure Differential
Scanning Calorimeter (HP- DSC) was used to determine Tg of the epoxy samples. The resin (8-12 mg) was weighed into a sealed aluminum DSC pan. The tests were performed at 200 psi under nitrogen and the temperature was ramped from -50oC to
250oC at a rate of 10oC/min.
Tensile tests were run according to ASTM D638. The tests were performed on MTS
800 instrument at a displacement rate of 0.55 inch per minute, using a 500 pound load cell. Optical measurement techniques using digital image correlation, as opposed to strain gages, were made using ARAMIS software. In image correlation, a random speckle pattern is painted on to the specimen. Cameras then track the displacements of the speckled dots, and displacement fields and strains are calculated by specialized computer algorithms. Once calibrated, the software can measure specimens under
95 loading and output strain and displacement results through automated methods without
user intervention being required.194
5.4 Results
The experimental results detailing silicate dispersion and nanocomposite properties are detailed in the following sections.
5.4.1 Characterization of silicate dispersion
Nanocomposite morphology analysis included both XRD and TEM. The XRD pattern
o of Cloisite 30B displays an intense diffraction peak at 2θ = 4.9 , d001 = 1.8 nm. Within
the 50:50 nanocomposites, the XRD pattern suggests exfoliation, based on the absence of
a diffraction peak corresponding to Cloisite 30B clay layer stacking. As the content of
the aliphatic component decreases, intercalation becomes prevalent. The 70:30 resin
o containing 5 wt% Cloisite 30B exhibits a low intensity diffraction peak at 2θ = 2.51 , d001
= 3.5 nm, thus implying an intercalated nanocomposite morphology. Both 90:10 nanocomposites exhibit diffraction peaks corresponding to intercalated clay. Peaks in
o both samples appeared at 2θ = 2.51 , d001 = 3.5 nm. Representative XRD patterns are shown in Figures 5.2a-c.
XRD patterns may be deceiving when classifying nanocomposite morphology. The
absence of a diffraction peak is not sufficient to confirm the clay layer separation as
exfoliated. The absence of the diffraction peak may only suggest that the layers have lost
regular spacing between sheets. This would indicate that resin has intercalated into a
portion of the clay galleries. Alternatively, the clay tactoids may be reduced to a size that
96 * 5% 30B
y * 2% 30B t i * s n e t In
Cloisite 30B
246810 2 Theta
* * 5% 30B y
it 2% 30B s n e t In
Cloisite 30B
246810 2 Theta
5% 30B y t i 2% 30B ns e t In
Cloisite 30B
246810 2 Theta
Figures 5.2.(a-c) XRD patterns of 90:10, 70:30, and 50:50 nanocomposites, respectively
97
(a) (b)
(c)
Figure 5.3.(a-b) Representative TEM images of 2 wt% and 5 wt% 30B, respectively, in
70:30 systems. (c) illustrates dispersion of 2 wt% 30B in 90:10 resin
98 is not detected by XRD. Therefore, TEM is useful to provide a clear understanding of the nature of silicate separation within each material. The representative TEM images shown in Figure 5.3 indicate a mixed nanocomposite morphology; with all samples containing regions of both intercalated and exfoliated silicate layers.
5.4.2 Interaction between Cloisite 30B and aromatic or aliphatic compounds
The interaction between the silicate clay and an aromatic or aliphatic compound was empirically evaluated by the addition of 0.5 g Cloisite 30B to 20mL of either chlorobenzene or hexanes. On addition to hexanes, the clay remained in powder form and immediately settled to the bottom of the solution. Mixing Cloisite 30B with chlorobenzene resulted in instantaneous swelling and dispersion of the clay throughout the solvent, as pictured in Figure 5.4. The swelling indicates a much greater attraction between the aromatic compound and the silicate clay.
The observation that Cloisite 30B has a certain affinity for aromatic components suggests that, within the epoxy blend nanocomposite, the clay may reside in closer contact with the more rigid, aromatic component of the epoxy blend. Therefore, samples were prepared by pre-swelling the clay in the rubbery, aliphatic component. This served two purposes. Primarily, the pre-swelling step forces the clay into preferential contact with the rubbery component of the blend, thereby reinforcing the weaker regions of the material. Secondly, pre-swelling facilitates silicate layer dispersion. The effect of pre- swelling on the distance between clay layers is demonstrated in Figure 5.5. The XRD
o pattern of the pre-swollen clay shows two diffraction peaks. The peak at 2θ = 4.8 , d001 =
1.7 nm, corresponds to the unswollen Cloisite 30B. Following pre-swelling in the aliphatic component of the blend, this peak decreases in intensity by approximately 50%
99
Figure 5.4 Cloisite 30B aggregates in hexanes (left) and dispersed in chlorobenzene
(right)
100
y it *
ns * e Pre-swollen 30B Int
Cloisite 30B
0 5 10 15 20 25 30 2 Theta
Figure 5.5. XRD pattern of Cloisite 30B and 2% 30B pre-swollen in DY3601
101
(a) (b)
Figure 5.6(a-b) Representative TEM images of 2% 30B in 70:30 epoxy matrix prepared
by (a) simple mixing, and (b) pre-swelling the clay layers
102 o and a second diffraction peak appears at 2θ = 2.4 , d001 = 3.4 nm, corresponding to the
clay layers which are intercalated with the aliphatic component. Figure 5.6 illustrates the
difference in layer separation that pre-swelling provides. The TEM images compare the
layer dispersion in a 70:30 2 wt% Cloisite 30B sample, with and without having pre-
swollen the clay. A much greater level of dispersion is achieved following pre-swelling.
5.4.3 Glass Transition Temperature
Resin and nanocomposite Tg values are listed in Table 1. The resins containing a
higher aromatic content results in a higher Tg. The clay offers little to no effect on Tg
within the 90:10 series of resins. Within the 70:30 and 50:50 resins, the matrix Tg
decreases upon simple mixing of the clay and epoxy, i.e. the clay is not pre-swollen with
the rubbery material.
Pre-swelling Cloisite 30B in the aliphatic component has a dramatic effect on Tg within the 70:30 series. Following dispersion of 2 wt% Cloisite 30B into this resin, the
o Tg initially drops by 5 C. Pre-swelling the clay in the rubbery component, however,
o increases the Tg above that of the base resin by 7 C. The same trend is observed at 5 wt% loading, but to a greater extent.
The purpose of tuning Tg by modification of aromatic content is to allow comparison
of the nanocomposite tensile properties, based on matrix mobility. The blend approach
allows preparation of matrices that are either rubbery (50:50) or glassy (70:30, 90:10), at
room temperature.
103
Table 5.1. Tg measurements by DSC
o o o Clay Tg (C) Tg (C) Tg (C) Content 90:10 70:30 50:50 0% clay 63 +/- 2 32 +/- 2 0 +/- 1 2% 30B 62 +/- 2 29 +/- 1 -1 +/- 2
2% 30B 67* 41* 0* (pre- swollen) 5% 30B 66 +/- 1 30 +/- 8 -2 +/- 2
5% 30B 65* 45* 0 +/- 3 (pre- swollen) * Data from one sample
104 5.4.4 Tensile Tests
The yield stress (σy), strain to failure (εfailure), and Young’s modulus (E) of the silicate-
epoxy nanocomposites were determined from the stress strain curves. The material
toughness was calculated by the area under the stress strain curves. These values are of
interest in determining the strength, ductility, and stiffness of the nanocomposites. The
yield stress is the maximum stress that the material can sustain. As the yield stress is
approached, damage within the material causes plastic deformation, and the initial state
of the material can not be recovered by relieving the stress. Young’s modulus is a
measure of the stiffness of the material, and is calculated from the slope of the linear
elastic region of the stress-strain curve. The resin stiffness is an important parameter for
resins utilized as the matrix of carbon fiber reinforced composites. Under compressive
load, a composite composed of a resin with increased modulus will experience reduced
fiber buckling. The compressive strength of carbon fiber reinforced composites is
considered to be a limiting property for many polymer composite applications.195
The strain to failure is of interest as it is a measure of the ductility of the sample.
Additionally, strain to failure, along with the yield stress, contributes to the overall area under the stress-strain curve. This area is a measure of material toughness, and is important in applications, such as composite cryo-tanks, where material microcracking is of concern.
The tensile test results are listed in Tables 5.2-5.5. Sample curves are illustrated in
Figures 5.7a and 5.7b.
105
Table 5.2. Yield stress as determined from tensile tests
Clay Content *σy (psi) σy (psi) *σy (psi) 90:10 70:30 50:50 0% clay 9741 +/- 200 2568 +/- 6 189 +/- 21
2% 30B 9549 +/- 170 2281 +/- 71 270 +/- 13
2% 30B (pre- 8770** 3673 +/- 198 328 +/- 5 swollen) 5% 30B 8355 +/- 221 3324 +/- 31 352 +/- 16
5% 30B (pre- 7690 +/- 451 4193 +/- 129 365 +/- 5 swollen) * Stress at failure.
** Data from one sample due to air bubbles within the material.
106
Table 5.3. Strain to failure as determined from tensile tests
Clay Content εfailure (%) εfailure (%) εfailure (%) 90:10 70:30 50:50 0% clay 11 +/- 4 45 +/- 8 27 +/- 1
2% 30B 4 +/- 1 47 +/- 2 30 +/- 2
2% 30B (pre- 2 43 +/- 4 30 +/- 2 swollen) 5% 30B 2 +/- 0 51 +/- 2 40 +/- 1
5% 30B (pre- 2 +/- 0 34 +/- 5 37 +/- 3 swollen) * Data from one sample due to air bubbles within the material.
107
Table 5.4. Young’s modulus as determined from tensile tests
Clay Content E (psi) E (psi) E (psi) 90:10 70:30 50:50 0% clay 4509 +/- 57 1124 +/- 53 9 +/- 1
2% 30B 4504 +/- 1 997 +/- 179 11 +/- 1
2% 30B (pre- 4701* 1640 +/- 162 13 +/- 2 swollen) 5% 30B 4762 +/- 70 1575 +/- 147 12 +/- 2
5% 30B (pre- 4687 +/- 98 2023 +/- 182 15 +/- 1 swollen) * Data from one sample due to air bubbles within the material.
108
Table 5.5. Toughness as determined from tensile tests
Clay Content Toughness Toughness Toughness (psi) (psi) (psi) 90:10 70:30 50:50 0% clay 1110 +/- 200 1270 +/- 90 34 +/- 3
2% 30B 190 +/- 20 1020 +/- 40 31 +/- 3
2% 30B (pre- 100 1430 +/- 110 74 +/- 5 swollen) 5% 30B 120 +/- 20 1450 +/- 40 39 +/- 4
5% 30B (pre- 90 +/- 10 990 +/- 150 79 +/- 10 swollen)
109
10000 Neat Resin 8000 ) i
s 6000 p Neat Resin ess (
r 4000 2% 30B t S 2% 30B preswell 2000 5% 30B 5% 30B preswell 0 03691215 Strain (%)
(a)
5000 5 wt% 30B, pre-swell 4000 2 wt% 30B, pre-swell 5 wt% 30B si)
p 3000 Neat Resin ( ss
e 2 wt% 30B r 2000 t S 1000
0 0 102030405060 Strain (%)
(b)
Figures 5.7(a-b) Stress – Strain curves of 90:10 and 70:30 series, respectively
110 5.4.4.1 General trends observed from tensile testing
The resin modulus tends to increase with increasing aromatic content, and increasing
clay content. An additional increase in modulus is observed when the clay is pre-swollen
with aliphatic component. The yield stress of the resin also increases with increasing
aromatic content. The yield stress is also raised on addition of clay, and generally to a
greater extent with pre-swelling. This is not true within the 90:10 series. In this case, the
nanocomposites do not yield, therefore, the reported σy is actually the stress at failure. As the nanocomposites failed earlier than the neat resin, the reported σy of the nanocomposites is lower than that of the base resin.
Strain to failure data is somewhat less reliable of a data point than the yield stress as it greatly depends upon flaws within the test specimen. Generally, an increase in E and σy leads to reduced strain to failure. This is observed within the glassy specimens (70:30 and 90:10), however in the 50:50 series, which is rubbery at room temperature, the strain to failure increases in the nanocomposite despite the increase of strength and modulus.
The toughness of the material tends to increase, with the addition of clay, within the
70:30 and 50:50 samples. The 90:10 samples, however, exhibit a drastic reduction in toughness.
5.4.4.2 Relation to matrix mobility
Table 5.6 summarizes changes in σy, E, and εfailure for each nanocomposite, as
compared to the base resin. The literature reports that improving the mobility of
nanoclay within the matrix enhances the toughness of the nanocomposite material as
compared to the base resin.46,114-117 This is reflected, in this case, as an increased strain to
failure. The 50:50 samples, rubbery at room temperature, showed improvement in strain
111
Table 5.6. Change in yield stress, Young’s modulus, and strain to failure on silicate
dispersion
Clay 90:10 70:30 50:50 Content
∆σy ∆E ∆εfailure ∆σy ∆E ∆εfailure ∆σy ∆E ∆εfailure 0% clay ------
2% 30B -2% within -64% -11% within within +43% within within error error error error error 2% 30B -10% +4% -82% +43% +46% within +74% +44% within (pre - error error swollen) 5% 30B -14% +6% -82% +29% +40% within +86% within +48% error error 5% 30B -21% +4% -82% +63% +80% within +93% +67% +37% (pre- error swollen)
112 to failure following dispersion of 5 wt% layered silicate. Within that material, pre- swelling the clay in the more mobile component has little effect on strain to failure, as the overall matrix is already rubbery at room temperature. Pre- swelling in the 50:50 resin does, however, reinforce the more flexible regions of the blend, resulting in an overall stronger and stiffer composite, which accounts for the more than 100% increase in the calculated toughness. The enhancement to yield stress and toughness is pronounced in this material, relative to either of the two glassy systems. This is an accomplishment considering that the increases in strength and stiffness are achieved along with improved ductility. It should be noted that the 50:50 resins also do not yield, therefore the yield stress represents the stress at failure. This is increased in all the nanocomposite samples, relative to the neat resin.
Selectively placing the clay in the more mobile component is of significant benefit to the yield stress and modulus of the 70:30 resins. The values of both properties are enhanced on strengthening the flexible portion of the resin blend. On addition of 2 wt% pre-swollen clay, this is achieved with no corresponding decrease in strain to failure, relative to the base resin. As a result, the area under the stress-strain curve increases; meaning the nanocomposite structure imparts a degree of toughness to the resin. The toughness of the 5 wt% pre-swollen clay sample does decrease as the strain to failure decreases.
The least dramatic difference was observed with the 90:10 resins. In this case, pre- swelling the clay had no effect on the overall properties. This may be attributed to the already small concentration of flexible component present in the system. Additionally, as with the 50:50 resins, this system does not yield. Therefore, the σy values of the
113
Table 5.7. Oxygen permeability of epoxy blend base resins and nanocomposites.
Clay Permeability Permeability Permeability Content (cc·mil/m2·day) (cc·mil/m2·day) (cc·mil/m2·day) 90:10 70:30 50:50 0% clay 1149 +/- 42 303 +/- 37 1277 +/- 75
2% 30B 1351 +/- 44 355 +/- 72 2249 +/- 106
2% 30B 992 +/- 62 317 +/- 36 2198 +/- 65 (pre- swollen) 5% 30B ------323 +/- 22 2110 +/- 71
5% 30B 990 +/- 26 263 +/- 18 2366 +/- 53 (pre- swollen)
114 nanocomposites in the 90:10 series are reduced relative to the base resin because these
nanocomposites fail earlier than the base resin.
5.4.5 Permeability
The permeability data listed in Table 5.7 offers some difficulty in explanation. There
are two points of confusion. First, it was initially believed that increasing aliphatic
content would result in increased permeability, due to the more mobile nature of this
component. This was not observed, and the permeability of the 70:30 resin is
considerably lower than that of either the 90:10 or 50:50 resins. Furthermore, it was
expected that dispersion of the layered silicate would reduce the permeability of the resin,
as is most often observed in nanocomposites.86, 177-178, 206 This again was not observed. In fact, the permeability of the 50:50 nanocomposites was up to 85% higher than that of the base resin. Within the 90:10 resins, a small but real reduction in permeability was observed on pre-swelling the aliphatic component with clay.
The 70:30 resin exhibited a permeability that is significantly lower than the other two resins. Here again, the permeability of the nanocomposites tends to be within experimental error of the base resin permeability. There is no explanation available for the high permeabilities observed.
5.4.6 Coefficient of Thermal Expansion
CTE is a critical concern for resins used in composite structures. Residual stresses are generated on thermally cycling composites where matrix expansion/contraction is significantly greater than that of the reinforcing phase. The addition of nanoclay to an epoxy resin is reportedly a viable method to reduce resin CTE.196-197 Table 5.8 compares
115
Table 5.8. Values of the coefficient of thermal expansion
Clay CTE CTE CTE Content (µm/moC) (µm/moC) (µm/moC) 90:10 70:30 50:50 0% clay 66 66 66
2% 30B 65 76 50
2% 30B 60 72 44 (pre- swollen) 5% 30B 60 66 45
5% 30B 65 60 41 (pre- swollen)
116 the CTE values of the neat resins and corresponding nanocomposites at temperatures
below Tg.
A reduction in CTE is observed in most nanocomposites as a result of the rigid clay
nanoparticle inhibiting the motion of polymer chains which are in contact with the
silicate surface.118-119 By restricting chain motion, the dimensional stability of the resin is improved. Within the 90:10 resins, there is little reduction in CTE observed for the nanocomposites. This may be attributed to the fact that the system is initially very rigid and therefore no benefit is derived from dispersion of the rigid filler particles. However, the nanocomposite structure offers a large reduction in the CTE of the low Tg, 50:50
resin. In this material, up to a 40% reduction in CTE is noted. The CTE of the
nanocomposites is reduced as the clay loading increased, and further reductions are
observed on pre-swelling the clay with the aliphatic component. Both variables serve to
restrict chain motion on heating.
5.5. Discussion
The following sections offer an interpretation of the experimental data based on the
epoxy blend microstructure and the dispersion of clay in the blend.
5.5.1 Microstructure
The microstructure of the blend employed throughout this chapter is important for
interpretation of the observed results. The single Tg reported within each system is a
good indication that the components are miscible, and there is no macrophase separation
of the aromatic and aliphatic components of the blend.
117 However, the results indicate that there is a level of phase separation which allows for
manipulation of the material properties. Within the 70:30 system for example, the
majority of the blend is composed of the rigid, aromatic components. While the rubbery
component is likely well integrated into the blend, there will remain large regions of the
blend composed solely of the aromatic component. The relative volumes of the aromatic
to aliphatic components present in the blend will produce regions rich in the rigid
segments. Therefore, an affinity between the clay and the aromatic segment can pull the
clay into these regions; leaving the rubbery component of the blend deficient in clay
content. The SEM image in Figure 5.8 shows distinct regions of the 70:30 blend
containing significant quantities of clay, while adjacent regions of the blend contain little
visible filler.
The 50:50 series contains essentially equal quantities of both aromatic and aliphatic
components. Again, a single Tg is observed indicating no macrophase separation. SEM images of the fracture surface show much greater homogeneity in the silicate separation, relative to the 70:30 blend, where regions rich in silicate clay are clearly identifiable.
The SEM image of Figure 5.9 represents the dispersion of Cloisite 30B in the 50:50 blend. Within the 50:50 series, all mechanical properties, (E, σy, and εfailure) are improved on simply dispersing Cloisite 30B into the epoxy blend. This can be attributed to rubbery nature of the 50:50 resin and the consequent mobility of the silicate layers under tensile loading. Additional improvements in mechanical properties are obtained with pre- swelling of the clay. While there is no phase separation in the blend, a portion of the weaker blend component is strengthened by pre-swelling. Therefore, E and σy are observed to increase even in the more miscible blend.
118
Figure 5.8. Representative SEM images of the fracture surface of the 70:30 resin
containing 5 wt% 30B
119
Figure 5.9. Representative SEM images of fracture surface of 50:50 resin containing 5
wt% 30B.
120 5.5.2 Preferential interaction between silicate clay and aromatic components
The notion that a clay may prefer one component of a blend over another is gaining
attention. With respect to specific affinity of the clay to aromatic over aliphatic
components, a more comprehensive study reported in the literature198 describes the ability
of hexadecyl-functionalized synthetic clays to disperse and form gels in various organic
solvents. It was found that at 5-10 wt%, the organically modified clays form gels in
selected aromatic solvents. The gels of all clays used in the study were stronger in
toluene than in a branched alkyl solvent. The conclusion was that the synthetic clays
were potentially better suited for dispersion into polystyrene, which is structurally more
similar to aromatic solvents, than in polymers structurally similar to the alkyl solvent, such as polypropylene.
5.5.3 Pre-swelling
Pre-swelling a layered silicate with a reactive monomer is not a novel concept.
However, the purpose has generally been to improve clay layer separation. Ishida, et
al.100 used an epoxy monomer or a polydimethylsiloxane chain to increase silicate gallery
spacing prior to introducing the clay into a host of matrices. They found that by
matching, or significantly mismatching, the χ parameter of the swelling agent and that of the matrix, they were able to improve the level of silicate dispersion in a number of matrices.
Park and Jana106 also employed the pre-swelling technique to improve silicate dispersion in an epoxy matrix. Their work reported that filling the clay galleries with a low molecular weight epoxy, relative to the molecular weight of the surrounding matrix, results in high elastic forces within the clay galleries on curing. These elastic forces are
121 sufficient to overcome the viscous forces acting on the exterior of the clay, thereby
facilitating exfoliation. Gintert and Jana extend this theory to include high temperature
polyimides, where the silicate gallery is pre-swollen by a low molecular weight imide
oligomer to enhance dispersion in a crosslinked polyimide matrix.56 Many other researchers have pre-swollen silicate with a reactive monomer to aid in dispersion.59,74,98-
100 The work presented in this chapter is unique in that pre-swelling the silicate in
monomer was specifically designed to selectively place the clay in a particular region of a
blend to enhance specific properties.
5.5.4 Interpretation of Results
It has been noted in the literature that clay layer mobility is relevant to obtaining
nanocomposite mechanical properties that are improved compared to those of the pristine
resin. The results of this work show that it is equally important to reinforce the regions of
a polymer blend that dominate the properties of interest.
5.5.4.1 Glass Transition Temperature
The Tg data shows that the greater the amount of aromatic compound, the higher the
glass transition temperature. Furthermore, Tg generally dropped following simple mixing of Cloisite 30B into the blend. There are two primary factors which affect the glass transition temperature in thermosetting nanocomposites. One is the effect of the organic modifier, while the other is any change in the crosslink density which may occur due to the presence of the clay. Both scenarios are plausible contributors to the observed drop in
o Tg upon dispersion of silicate. Within the 70:30 resins, the Tg decreases up to 6 C.
The effect of pre-swelling the clay in the rubbery component is to recover the initial
Tg drop due to the presence of the clay. Pre-swelling forces the clay into the more mobile
122 component. Reinforcing the regions that contribute to overall lower Tg results in an increase of the blend Tg. In the case of the 70:30 blends, that recovery pushes the Tg higher than that of the base resin, and an overall increase of up to 11oC is observed.
5.5.4.2 Mechanical Properties
The most significant benefit of pre-swelling is observed within the mechanical property data. In this case, positioning the silicate reinforcement in the rubbery regions allows improved mobility of the clay, translating into improved composite strength and modulus.
Generally speaking, the clay benefits the weaker, rubbery materials to a greater extent than the more rigid materials, as is outlined in Table 5.5. Likewise, forcing the clay into the more rubbery components of a blend has a similar result, and the overall properties of blend properties improved. This data will be useful in designing toughened epoxy systems for any number of applications.
5.6 Conclusions
The observed preference of Cloisite 30B for aromatic containing compounds over aliphatic containing compounds enables preferential placement of the clay within the rubbery regions of an epoxy blend. This placement is achieved by pre-swelling the clay in the aliphatic component of the blend prior to addition of the aromatic component. This pre-swelling step also improves silicate dispersion, as evidenced by XRD and TEM.
The increase in Tg and decrease in CTE on pre-swelling the clay provides evidence
that the clay restricts the motion of epoxy chains within the mobile component of the
blend. Tensile tests of epoxy blends reveal that mobility of the silicate layers within the
123 matrix offers improved resin ductility. This is seen by the improved strain to failure and
significant enhancement in strength within the 50:50 resins, relative to other resin in the
series which are glassy at room temperature.
Placing the clay in the mobile component offers additional enhancements to both
strength and modulus. As the aliphatic component is also the mechanically weaker
component of the blend, the pre-swelling step provides a mechanism to reinforce the
weaker component, thereby offering additional enhancements to the strength, modulus, and glass transition temperature of the epoxy nanocomposites.
124
CHAPTER VI
FUNCTIONALIZED GRAPHITE NANOCOMPOSITES
6.1 Background
The previous chapters illustrate the importance of considering the chemistry and
nature of the surface treatment applied to a layered silicate clay. Layered graphites offer
similar opportunities for surface modification. Such modification is typically achieved
through a multi-step process including oxidation followed by functionalization with an
organic compound. This chapter will demonstrate the trade-off in material properties
observed as the level of graphite functionalization is varied. An understanding of such
trade-offs is necessary to tailor the functionalization of the nanostructure and provide
nanocomposites with the desired properties for a specified application.
6.2 Introduction
As advanced materials applications dictate increasingly rigorous composite performance, innovative technologies capable of providing properties beyond those of
traditional polymer matrix composites are becoming a necessity. Recently,
multifunctional composites providing structural integrity, as well as serving additional
functions have been widely investigated.199-200 Nano-materials in particular have been called upon to provide such capabilities.
125 Nano-particulate fillers enable property enhancement as a result of the large interface
available to the matrix.201-202 This interfacial area is only beneficial however, if contact
between the matrix and the particle is optimized. Several techniques to modify nano-
particle surfaces have been identified, and vary, based on the chemistry of the nano-
particle. In the case of graphite nano-flakes, the interfacial strength may be tailored by
oxidation of the graphene plane, or by reactive finish technology.203 However, as with any system, chemical modification results in material property trade-offs. In their pristine form, graphene platelets are characterized by a low surface energy, and therefore are poorly wetted by most polymer matrices. Consequently, poor graphite layer dispersion results, leading to reduced mechanical properties of the composite. However, the aromatic nature of un-oxidized graphene planes offer greatly enhanced transport properties, such as thermal and electrical conductivity. Such potential is reached when the graphene sheets are dispersed into individual platelets, which again is difficult in a situation where wetting is poor. Therefore, oxidation and functionalization is necessary
to improve dispersion, however this occurs at the expense of conductivity. As is true for
most nanocomposite systems, there exists a balance where wetting/chemical interaction is
strong, dispersion is acceptable, and both mechanical and transport properties are
enhanced. As research in this area is still relatively new, in comparison to layered silicate
work, there has yet to be a study which compares the property tradeoffs as a function of
oxidation, functionalization, and dispersion. This work addresses those issues.
In this chapter, graphite flakes functionalized by three separate methods were utilized.
Expanded graphite (EG) was oxidized along the edges of the basal plane by rapid heating
of a graphite intercalation compound. Functionalized graphene sheets (FGS), where
126 chemically oxidized graphite was split into individual graphene sheets through a rapid heating process. In the third approach, an epoxy coating was covalently bonded to the graphite surface through a coupling agent. This material was prepared by Adherent
Technologies, Inc., and will be referred to as ATI.
Therefore, the ATI and FGS graphites were functionalized within the graphene basal plane with epoxy and/or hydroxyl functionalities. Further functionalization was provided by the addition of excess amine curing agent to these nanocomposites. This technique allows covalent bonding between the graphite particles and the epoxy matrix
6.3 Experimental
Experimental details of the preparation and characterization of epoxy/ graphite nanocomposites are described.
6.3.1 Materials
Epoxy resin, Epon 826, was generously supplied by Resolution Performance Products.
Araldite DY3601, an aliphatic epoxy resin, and Jeffamine D230 curing agent were supplied by Huntsman Chemicals. Expanded graphite was supplied by Superior
Graphite, Inc. TG679 is a Graftech product which was modified by Adherent
Technologies, and will be referred to as ATI graphite. FGS was obtained through a grant with Princeton University.
6.3.2 ATI preparation
Expanded graphite, TG-679, (10.0 g) was combined with ATI-9307 coupling agent
(0.01g) in 50 mL of 2-butanone (MEK). The mixture was allowed to stir for twenty minutes to ensure homogeneity. The solvent was then removed by rotary evaporation.
127 The treated graphite was combined with EPON 828 (0.01 g) in 50 mL CH2Cl2. The mixture was allowed to stir for twenty minutes, and the solvent was removed under vacuum by rotary evaporation. The graphite was then placed under vacuum at 50°C to remove residual solvent. To activate the coupling agent, the graphite was heated in a vented oven at 210°C for twenty minutes. Once cool, the graphite was placed in an amber vial and purged with argon.
6.3.3 Nanocomposite preparation
Resin plaques of Epon 826 and DY3601 epoxy blends were prepared in 70:30 equivalent epoxy ratios with graphite content ranging from 0 wt% to 3 wt%. Epon 826
(18.4 g), DY3601 (7.875 g); and the calculated amount of graphite were mixed in a jar and sonicated at room temperature for 4 h. The curing agent, D230 (7.5 g), was added and the contents of the jar were poured into a 10.2 cm by 10.2 cm mold. The resin was degassed at 40°C for 3 h then cured at 75°C and 125°C for 2 h each. The preparation of nanocomposites prepared with excess amine curing agent followed the above procedure, using 7.9g D230 curing agent, 5% excess amine, or 8.25g D230, 10% excess amine.
6.3.4 Organic modification of graphite
The FGS and ATI nanoparticles were organically modified with methylene dianiline,
MDA. The MDA structure is shown in Figure 4.1. Either the FGS or ATI graphite (0.5 g) was dispersed in distilled water (85 mL) at room temperature, under nitrogen flow.
MDA (1.5 g, 7.6 mmol) was dissolved in 25 mL ethanol. The MDA/ethanol solution was
added drop-wise to the aqueous suspension under vigorous stirring. The mixture stirred
for 24 hours. The modified graphite was dried, washed repeatedly with warm ethanol,
128 and dried in a vacuum oven overnight at 90oC.204 These materials were characterized by
TGA to calculate the amount of MDA that had been grafted to each material.
6.3.5 Characterization
XRD and TEM experiments proceeded as described in Chapter 4. Tensile testing followed the procedure outlined in Chapter 5.
Electrical conductivity measurements were made by cutting samples of approximately
3.5 cm by 0.7 cm from larger preparative composites. Silver paint was used to apply electrical contacts to the ends of the sample, covering top, bottom, sides, and end. The thickness, width, and length dimensions of the composite were measured using a digital micrometer. The length used for calculations is the gauge length between the electrodes.
The samples were mounted using spring clips into a Keithley Model 8002 Restest High
Resistance Test Box to minimize stray currents. A constant voltage of 100.0 V dc was applied across the sample and current was measured using a Keithley Model 617
Electrometer.
BET measurements on graphite surfaces included preliminary sample degassing by evacuating to 100 mm Hg, followed by heating to 80oC and holding at temperature for 8
hrs. They were then analyzed for surface area and/or porosity using ultra high purity
nitrogen at LN2 temperature. Measurements were made using a Micromeritics
Instrument Corp., ASAP 2020 Surface Area and Porosity Analyzer.
X-ray photoelectron spectroscopy samples were analyzed on the VG Mk II
ESCALAB, using Mg Kα x-rays at 300 watts. The system consists of a turbo pumped
prep chamber (base pressure 10-9 torr), a fast load entry lock; and a main chamber which
maintains a base pressure of 1 x 10-10 using a diffusion pump with a cold trap. The
129 spectra were taken with the surface plane normal to the analyzer axis (90° take off angle) using an aperture of 1 mm x 1 mm. Survey scans were initially taken to identify all components, followed by higher resolution individual region scans using a pass energy of
20eV. The spectrometer was calibrated to yield the standard values of 75.13 eV for Cu3p and 932.66 eV for Cu2p3/2. Atomic concentrations were then calculated based on the individual regions.
Viscosity measurements were performed during mixing graphite with the epoxy
matrix at room temperature for times ranging from 15 minutes to 2 hours. The
measurement was made by submerging the tip of a hand held Brookfield Synchro-lectric
Viscometer into the resin and graphite mixture, and reading the measured viscosity. The
viscosity is a measure of the resistance necessary to rotate the viscometer tip.
6.4 Results
The experimental results detailing graphite dispersion and nanocomposite properties are
detailed in the following sections.
6.4.1 Characterization of Graphite Materials
The three graphite materials utilized in this work were functionalized by edge
oxidation following intercalation/vaporization of acid (EG), epoxy functionalized (ATI),
and chemical oxidation followed by rapid thermal treatment (FGS). The resulting
oxygen content varies significantly by functionalization technique. Table 6.1 summarizes
the results of XPS analysis, FTIR functional group characterization, and BET analysis for
each material.
130
Table 6.1. Summary of graphite nanoparticle characterization
Sample Density XPS BET FTIR (g/cm3) (% oxygen) (surface area) EG 2.25 1% 38 m2/g 3600 cm-1 (OH- stretch), 1700 cm-1 (C=O stretch) ATI 1.7 9% 17 m2/g 3300 cm-1 (weak, OH- stretch), 2900, 2800 cm-1 (C-H stretch), 1700 cm-1 (C=O stretch), 1100 cm-1 (C-O/C-O-C stretch) FGS 0.76 13% 640 m2/g 3400 cm-1 (OH-stretch), 1550 cm-1 (C=O stretch), 1100 cm-1 (C-O/C-O-C stretch)
131 The results of the FGS analysis indicate that chemical oxidation, followed by thermal
treatment, yields heavily oxidized, graphene sheets. Because the graphene sheets are
well separated, an enormous surface area is available for chemical or physical interaction
with the matrix resin.
In the case of ATI, the presence of the epoxy coating is evidenced by the XPS and
FTIR data. However, the available surface area is small, relative to FGS, because the
graphite aggregates remain intact, with functionalization occurring only on the outermost
layers. In this case, as with EG, in-situ graphite dispersion becomes an important processing parameter during the nanocomposite preparation. The physical
characterization of each graphite material is described below.
6.4.1.1 Expanded Graphite Characterization
Three primary factors contribute to the difficulty in dispersing EG into an epoxy
matrix. They include: incomplete graphite layer expansion, large platelet size, and low
surface energy. As discussed in the introduction, expanded graphite is prepared by
intercalation and subsequent vaporization of an acid, such as sulphuric and nitric acid.
However, characterization of the graphite intercalation compound and the expanded
graphite worm indicates that incomplete acid intercalation results in incomplete
expansion. Therefore, EG is composed of regions containing well expanded graphite,
and regions of closely packed graphite platelets. This is illustrated by the XRD pattern of
the graphite intercalation compound, Figure 6.1a, showing two diffraction peaks. One
o corresponds to the d002 spacing of natural flake graphite (2θ = 26.7 and d002 = 0.334 nm)
and the second corresponds to the d002 spacing of the intercalation complex, where 2θ =
o 25.63 and d002 = 0.348 nm. Therefore, on thermal shock, the intercalated graphite is 132
(a)
Acid Intercalated Graphite y it s n e t In
8 121620242832 2 Theta
(b)
Expanded Graphite y t i s n e t In
8 121620242832 2 Theta
Figure 6.1.(a-b) XRD patterns of (a) acid intercalated natural flake graphite, before
thermal expansion and (b) expanded graphite, following thermal expansion
133
Figure 6.2. SEM image of EG illustrating regions of closely spaced graphite layers, and
well separated graphite layers
134 expanded and results in an amorphous diffraction pattern, as shown in the XRD pattern of
Figure 6.1b. However, stacking of the unintercalated layers remains unchanged. This incomplete expansion is also observed by SEM, as shown in Figure 6.2.
The EG platelet size ranges from 5 to 20µm in length,123 which is at least 5 times greater than that of layered silicates or carbon nanotubes. Platelet size can be an issue because upon addition of EG to a resin, the viscosity of the resin increases. The increase in viscosity, along with the large platelet size, makes movement and dispersion of the graphene layers very difficult. Sonication is a useful technique to break apart the stacked graphene layers, but also serves to reduce the graphene platelet size. Therefore, sonicating EG in ethanol, for increasing lengths of time allowed characterization of both the reduction in aggregate size and platelet size, by XRD and SEM respectively. The decrease in particle size was tracked by SEM and is depicted in Figure 6.3. The reduction in aggregate size on sonication is monitored by XRD, and the diffraction
patterns for the three sonication times are shown in Figure 6.4. The intensity of the d002 peak corresponds to the number of graphene sheets contained within a graphite aggregate. As sonication time increases, the d002 peak intensity decreases, suggesting a reduction in aggregate size.
6.4.1.2 Characterization of Adherent Technologies Graphite
Dispersion difficulties are also anticipated with ATI graphite due to the close packing of the graphite layers. In this case, XRD characterization results in an intense d002 diffraction peak, indicating that the graphene layers are arranged in a regularly spaced, ordered structure. This is illustrated in Figure 6.5. The graphene stacking registry is maintained throughout functionalization, therefore epoxy monomer is present only on the
135
Figure 6.3. SEM images illustrate decrease in EG particle size on increased sonication
time
136
4 hours y
it 8 hours
ns 16 hours e Int
20 25 30 2 Theta
Figure 6.4. XRD pattern showing decreased peak intensity as EG aggregate size is
reduced
137
ity s n te In
010203040 2 Theta
Figure 6.5. XRD pattern of ATI graphite showing intense peak corresponding to graphite
layer stacking
138
(a) (b)
Figure 6.6. SEM images of (a) untreated graphite and (b) graphite covalently bonded with
epoxy monomer
139 outermost layers of the graphite aggregate. As a result, several graphene layers remain
unmodified within the aggregate and are arranged in closely packed layers, as with the
expanded graphite material. The difference however is in the particle size. Where a
typical graphene sheet within an EG aggregate has dimensions on the order of 5 µm to 20
µm in length and 2 nm to 5 nm in thickness, the dimensions of the ATI material has been
reduced through processing to dimensions of 100-200 nm in length and 10 to 20 nm in
thickness. As a result, dispersion of this material was achieved through sonication
techniques with little difficulty. The presence of the covalently bonded epoxy monomer
is clearly evident by SEM examination, and is shown in Figure 6.6.
6.4.1.3 Characterization of Functionalized Graphene Sheets (FGS)
Unlike the previous graphite materials, the XRD pattern of FGS, Figure 6.7, does not
exhibit a diffraction peak at 2θ = 26.7o, corresponding to the interlayer spacing of natural
flake graphite. There is however a broad amorphous peak ranging from approximately
2θ = 9o to 2 θ = 20 o. This suggests that FGS is not a fully exfoliated system, but contains
graphene planes stacked at distances ranging from 0.45 nm to 0.93 nm.
However, the oxidized nature of the material and the small particle dimensions greatly
faciliate dispersion in the epoxy matrix. A consequence of the highly oxidized nature is
that the graphene sheets become wavy, or less rigid. This is illustrated by the SEM image
in Figure 6.8.
The more aggressive oxidation/processing techniques tend to have a larger influence
over the flake dimensions. For example, sonicating the EG in epoxy reduces the number
of layers stacked in an aggregate, as well as reduces the flake size.205 Chemically coupling the reactive epoxy coating (ATI) considerably reduces the graphene flake size, 140
y it
ns EG e
Int FGS
8 121620242832 2 Theta
Figure 6.7. XRD pattern of FGS showing complete separation of the graphene layers
141
Figure 6.8. SEM image of FGS
142 relative to natural flake graphite. The FGS preparation yields graphene sheets, on the
order of 6 µm in length and less than 2 nm in thickness, greatly simplifying dispersion.
6.4.2 TGA Determination of Graphite Functionalization
TGA is useful in determining the degree of functionalization of FGS and ATI. The
TGA plot in Figure 6.9 shows weight loss as a function of temperature for the pure ATI graphite and the organically modified ATI. Figure 6.10 illustrates the same data for FGS.
The TGA curve of the ATI graphite alone shows a weight loss of approximately 4%, before the material rapidly degrades above 600oC. Following modification with methylene dianiline, the TGA plot shows ~8% weight loss between 240oC and 600oC, before rapid degradation, again above 600oC. This result indicates that MDA is present in 4% by weight of the sample, or 0.22 mmol MDA per 1 gram of ATI graphite.
The TGA plot of FGS graphite shows degradation beginning at 500oC. Following modification with MDA, degradation of the modified graphite occurs in three steps. The first two degradation steps may be attributed to degradation of the MDA modifier.
Between 160oC and 550oC, 54% by weight of the organically modified graphite is lost.
That corresponds to 5.9 mmol MDA per 1 gram of FGS. This amount of modification is approximately 25 times greater that of the ATI graphite material.
The degradation of MDA from FGS occurs in two steps. The majority of degradation
occurs early, with 35% weight loss observed between 160oC and 350oC. An additional
19% weight is then lost between 350oC and 550oC. The fact that two separate degradation steps occur may be due to a portion of the MDA (diamine) modifier covalently bonding to two graphene layers. This is more likely to occur within the FGS material, rather than the ATI. FGS is heavily oxidized, and the layers are well separated,
143
ATI, as received 629oC 96%
Organically modified ATI
628oC 92%
Figure 6.9. TGA plot of as received ATI and organically modified ATI
144
FGS, as received
500oC Organically modified FGS 91%
350oC 65% 550oC 46%
Figure 6.10. TGA plot of as received FGS and organically modified FGS
145 creating a very large surface area available for modification. ATI functionalization occurs only on the outermost layers of the graphite aggregate. Therefore, these will be the only layers modified, and bonding between multiple aggregates, in solution, would be less likely.
6.4.3 Dispersion in an Epoxy Matrix
The TEM images in Figure 6.11 illustrate that surface treatment has a large effect on the level of dispersion achieved in the epoxy matrix. The differences in dispersion clearly illustrate that graphene layer functionalization and separation facilitates dispersion. The variation in dispersion was anticipated, as well as the difference the extent of functionalization/ dispersion had on physical properties is significant.
6.4.4 Physical Property Analysis
Differences in the physical properties of the separate graphite particles yield significant differences in processing of the epoxy nanocomposites. The variation in density between graphite materials necessitates that greater loadings, by weight, of the more dense graphite samples are used. Table 6.2 lists the volume percent of each sample prepared. The volume percent offers insight into the processing difficulties that may be encountered for the various samples. For example, EG at a 3.0 wt% loading occupies the same volume as FGS at a 1.0 wt% loading. It is at this level that challenges in processing were observed, as viscosity increases and restricts resin flow. Table 6.3 lists the resin viscosity changes on mixing time for a sampling of nanocomposite compositions.
Aside from density variations, the surface chemistry also holds the potential to interrupt epoxy processing and cure. In the case of ATI and FGS, epoxy functionalities
146
Figure 6.11. Representative TEM images of EG, ATI, and FGS, respectively, dispersed in an epoxy matrix
147
Table 6.2. Calculated volume percent of graphite used for each corresponding weight
percent
EG- wt% EG- vol% ATI- wt% ATI- vol% FGS- wt% FGS- vol%
0.5 0.25 0.5 0.33 0.50 0.74
1.0 0.50 1.0 0.66 1.0 1.47
3.0 1.49
148
Table 6.3. Measured viscosity on increased mixing time
Mixing EG, 0.5 EG, 1.0 ATI, 0.5 ATI, 1.0 FGS, 0.25 Time wt% (cP) wt% (cP) wt% wt% wt% (cP) (cP) (cP) 0 min 1600 1600 1600 1600 1600
15 min 9000 14000 3000 7150 12250
30 min 9750 15500 4200 7650 13300
1 hour 12500 16500 6750 8150 14000
1.5 hours 13500 15000 7200 8550 12750
2 hours 14500 12500 7450 9250 13000
149 introduced onto the graphene surface, may react with the amine curing agent. The
presence of such functionalities can reasonably upset the stoichometry of the epoxy resin and affect the properties of the cured epoxy nanocomposite. Therefore, within the ATI and FGS series, nanocomposites containing increasing concentration of amine curing agent were prepared.
6.4.4.1 Glass Transition Temperature
The addition of graphite to the resin matrix was observed to reduce the glass transition
temperature, as outlined in Table 6.4, where the data presented represents one specimen
of each composition. While nanocomposites prepared with a thermoplastic matrix often
205 result in an increased Tg, the Tg of thermosetting nanocomposites sometimes drops, as the presence of the nanoparticle may interrupt the crosslinking density of the resin itself.
This is likely the case, in the graphite nanocomposites, as the reductions in Tg values become larger with increased graphite loading.
6.4.5 Electrical Conductivity
The balance between particle size, dispersion, and extent of oxidation dictates a number of epoxy/graphite nanocomposite properties. High electrical conductivity is expected under the ideal condition of minimal oxidation and maximum dispersion.
Unfortunately, a low degree of oxidation yields particle aggregation therefore
necessitating high loading for conductivity. Furthermore, increased oxidation reduces conductivity, again necessitating higher loadings. The electrical resistivity data is summarized in Table 6.5.
It is evident from the data that within each series, increasing the graphite content reduces the nanocomposite resistivity, i.e. improves conductivity. The transport of
150
Table 6.4. Glass transition temperature as measured by DSC
o o o Sample Tg ( C) Sample Tg ( C) Sample Tg ( C) (EG) (ATI) (FGS) 0% graphite 34 0% graphite 34 0% graphite 34
0.5 wt% 32 0.5 wt% 35 0.5 wt% 32 EG/ 0% ATI/ 0% FGS/ 0% excess excess excess amine amine amine 1.0 wt% 30 1.0 wt% 33 1.0 wt% 27 EG/ 0% ATI/ 0% FGS/ 0% excess excess excess amine amine amine EG ATI FGS 3.0 wt% 28 0.5 wt% 32 0.5 wt% 32 EG/ 0% ATI/ 5% FGS/ 5% excess excess excess amine amine amine 1.0 wt% 29 1.0 wt% 31 ATI/ 5% FGS/ 5% excess excess amine amine 0.5 wt% 29 0.5 wt% 29 ATI/ 10% FGS/ 10% excess excess amine amine 1.0 wt% 29 1.0 wt% 29 ATI/ 10% FGS/ 10% excess excess amine amine
151
Table 6.5. Electrical resistivity of graphite-epoxy nanocomposites
Sample Resistivity Sample Resistivity Sample Resistivity (EG) (ohm-cm) (ATI) (ohm-cm) (FGS) (ohm-cm) Pure EG121 1.0 E -4 Pure ATI Not available Pure FGS204 1.0 E -1 0% graphite Non- 0% graphite Non- 0% graphite Non- conductive conductive conductive 0.5 wt% 3.0 E +10 0.5 wt% 1.3 E +7 0.5 wt% 5.0 E +10 EG/ 0% (non- TG/ 0% FGS/ 0% (capacitive) excess conductive) excess excess amine amine amine 1.0 wt% 1.0 E + 10 1.0 wt% 8.8 E +5 1.0 wt% 1.0 E +6 EG/ 0% (non- TG/ 0% FGS/ 0% excess conductive) excess excess amine amine amine EG TG 679 FGS 3.0 wt% 3.0 E +8 0.5 wt% 4.6 E +6 0.5 wt% 7.0 E +10 EG/ 0% TG/ 5% FGS/ 5% (capacitive) excess excess excess amine amine amine 1.0 wt% 1.4 E + 5 1.0 wt% 5 E +9 TG/ 5% FGS/ 5% excess excess amine amine 0.5 wt% 2.2 E + 6 0.5 wt% 8 E +9 TG/ 10% FGS/ 10% excess excess amine amine 1.0 wt% 9.1 E +5 1.0 wt% 1.0 E +6 TG/ 10% FGS/ 10% excess excess amine amine
152 electric current through the insulating epoxy resin requires a continuous network of pi orbitals to allow movement of the charge through the composite. The point at which a continuous conductive network is formed throughout the composite is referred to as the percolation threshold. This is the loading where the composite will exhibit conductivity.
Poor graphite dispersion or increased oxidation requires increased loading to reach percolation. The EG samples contain poorly dispersed graphite. Therefore we see a decrease in resistivity only as filler loading increases above 3 %. FGS is well dispersed, but heavily oxidized, therefore these nanocomposites show capacitive behavior below 1 wt%. In both cases, EG and FGS, conductivity is reached at 1.5 vol%. The SEM images of Figure 6.12 show a similar extent of graphite dispersion in each of these samples.
The greatest increase in conductivity is observed from the ATI samples, where oxidation and dispersion is intermediate relative to the other materials. Within these systems, the electrical resistivity decreases by 5 orders of magnitude at 1 wt% loading, or
0.67 vol%. This is less than half the loading by volume necessary for conductivity within the EG or FGS nanocomposites. In comparing the conductivity of the EG and FGS nanocomposites, it is important to note that oxidation of FGS does negatively affect the electrical conductivity of the bulk material. As shown in Table 6.5, the conductivity of bulk FGS is an order of magnitude lower than that of EG. Additionally, several FGS nanocomposites are noted as exhibiting capacitive behavior, meaning that on exposure to electric current, the material becomes charged, but that charge is quickly dissipated. This is likely a result of either the good dispersion inhibiting flow of electric current, or it is due to graphene oxidation interrupting the flow of electric current.
153
Figure 6.12. SEM images of 3 wt% EG in epoxy and 1.0 wt% FGS of epoxy
154
Table 6.6. Results of EG nanocomposite tensile tests
Expanded Graphite
Material E (psi) σy(psi) εfailure (%) Toughness (psi)
0 wt% 1124 +/- 53 2568 +/- 6 45 +/- 8 1270 +/- 90
0.5 wt% 1731 +/- 34 2861 +/- 206 20 +/- 3 590 +/- 60
1.0 wt% 2167 +/- 38 3456 +/- 111 3 +/- 0 70 +/- 35
3 wt% 2648 +/- 14 2659 +/- 13 2 +/- 0 40 +/- 13
155
4000 3500 1.0% EG
) 3000 5.0% EG Neat Resin si 2500 p 2000 0.5% EG ess (
r 1500 t
S 1000 500 0 0 102030405060 Strain (%)
Figure 6.13. Stress-strain curves of EG nanocomposites
156 The addition of excess amine did not have an effect on the electrical conductivity of these
samples. This would be expected as the conductivity is dependent only on dispersion and
functionalization. The surface chemistry on the graphite is a factor in dictating the level
of dispersion attained, but the conductivity is not affected by the strength of the
interfacial bonding.
6.4.6 Mechanical Properties
The purpose of adding nanoparticles to a matrix is to improve one or more properties,
without negatively affecting others. Ideally, the reinforcement would provide both an
increase in strength and an increase in ductility. However, with most nanocomposite
samples, the nanoparticle is very rigid in comparison to the polymeric matrix. As a
result, the dispersion of these particles tends to lead to an increase in stiffness and
strength at the expense of ductility. The results of the EG tensile tests are listed in Table
6.6. Representative stress strain curves are shown in Figure 6.13. Of all the graphite
fillers, the stress-strain data for EG shows the least benefit in terms of ductility or yield
strength. The sample containing 1 wt% loading does appear to have a significantly
enhanced yield stress. The primary benefit of this material is the increase in modulus as loading increases. At 3 wt% loading, Young’s modulus is doubled relative to that of the neat resin.
It should be noted that the drastic reduction in ductility and toughness may not be a
significant drawback to application of this nanocomposite material. Likely, such a nanocomposite would be employed as the matrix of a continuous carbon fiber reinforced composite. In this application, the strain to failure of the carbon fiber is approximately
2%. Therefore, comparing the mechanical behavior of the nanocomposite to the neat
157 resin up to 2% strain is what is important, if composite failure is initiated by fiber fracture
rather than matrix dominated effects such as buckling or shearing. When that comparison
is made, the yield stress and Young’s modulus dominate the material characteristics, and
the EG nanocomposites clearly out perform the neat resin. The primary benefit of a
nanocomposite prepared from this material is the significant enhancement in modulus,
compared to the base resin. Incorporating a stiffer matrix into a carbon fiber reinforced
composite reduces fiber buckling on compressive load, thereby enhancing the
compressive strength of the composite. This is an attractive benefit in the field of
polymer matrix composites, as a weak compressive strength is a common drawback to
these types of materials.
Analysis of the ATI graphite nanocomposites requires additional consideration of
samples prepared with excess amine curing agent. With this filler, the surface of the
graphite contains epoxy functionalities covalently bonded to the surface of the graphite basal plane. The purpose of adding additional amine is to enable reaction between the graphite surface functionality and the epoxy resin, without disrupting the stoichometry of the resin itself. Therefore, the reasonable property comparison is between the neat resin
of stoichiometric ratio, and the nanocomposites containing 0% to 10% excess amine
curing agent. The results are listed in Table 6.7 and representative stress-strain curves are
shown in Figures 6.14a and 6.14b.
In comparing the nanocomposites to the neat resin of stoichiometric ratio, there is a
significant impact on the yield stress and modulus following graphite dispersion. The
most significant is noted in the sample containing 0% excess amine, i.e. the
stoichiometric ratio of epoxy to curing agent. The addition of excess amine, in an
158
Table 6.7. Results of ATI tensile tests, comparing all nanocomposites to the
stoichiometric neat resin
Adherent Technologies 0.5 wt% ATI Graphite 1.0 wt% ATI Graphite Material E σy(psi) εfailure Toughness E (psi) σy(psi) εfailure Toughness (psi) (psi) (psi) Neat Resin 1124 2568 45 1270 +/- 1124 2568 45 1270 +/- 0% excess +/- 53 +/- 6 +/- 8 90 +/- 53 +/- 6 +/- 8 90 amine 0% excess 1470 3077 51 1300 +/- 1690 3344 39 1130 amine +/- 90 +/- 219 +/- 4 100 +/- 180 +/- 146 +/- 1 +/- 30 5% excess 1272 2943 62 1450 1096 2534 60 1530 +/- amine +/- 15 +/- 70 +/- 3 +/- 50 +/- 101 +/- 51 +/- 4 70 10% excess 1000 2608 67 1670 +/- 1170 2684 63 1270 +/- amine +/- 80 +/- 30 +/- 4 50 +/- 89 +/- 66 +/- 1 20
159 0.5 wt% ATI Graphite
3500 0% excess amine 3000 Neat Resin 5% excess amine 2500 si) 10% excess amine p 2000 1500 ess ( r t
S 1000 500 0 0 20406080 Strain (%)
1% adherent technologies
4000
0% excess amine 3000 Neat Resin s s e
r 2000 10% excess amine 5% excess amine St 1000
0 0 5 10 15 20 25 30 35 40 45 50 55 Strain
Figure 6.14. Stress-strain curves of ATI nanocomposites in comparison to the
stoichiometric neat resin
160 attempt to strengthen the interface, offers little benefit to the strength and modulus of
these nanocomposites, but ductility and toughness is improved. At 0.5% graphite
loading, we see little change in E, εfailure, or σy, except where 0% excess amine is added.
In that case, E and σy increase by 15% and 27%, respectively. Correspondingly, this is the only composition at 0.5 wt% graphite loading where εfailure and toughness remain unchanged. Comparing the compositions containing 1.0 wt% ATI, to the base
(stoichiometric) resin, we see an increased strain to failure in all the nanocomposites, except the 0% excess amine sample. This corresponds to the increased modulus observed at this composition. However, the 0% excess amine composition showed the greatest improvement in modulus and yield stress, 31% and 20%, respectively, with little drop-off in strain to failure. Clearly, when comparing the nanocomposite properties to the base resin, those samples prepared with no excess amine exhibit the most benefit from the graphite, with respect to strength and stiffness.
This result is rather surprising, as the addition of excess amine allows covalent bonding between the matrix and the nano-filler. However, remembering that the quantity of functionalization on this graphite material is rather small, based on XPS and TGA data, it is clear that the excess amine is not necessary to optimize the interface of these nanocomposites.
Comparison of the properties of the nanocomposites prepared with excess amine with those of the stoichiometric neat resin holds the most meaning regarding benefits obtained from altering the graphite-resin interface. However, it is also of interest to compare these nanocomposites with the neat resins also prepared with the same amount of excess amine.
This data is listed in Table 6.8.
161 Comparing the samples with excess amine to the corresponding neat resin, we see that
adding 5% excess amine to either 0.5 wt% or 1.0 wt% nanocomposites yields reduced
strength and stiffness, but improved toughness. Furthermore, the addition of graphite to a
10% excess amine material offers only an improvement in toughness, relative to the
corresponding base resin.
The mechanical property data for FGS nanocomposites is listed in Table 6.9. In this
case, comparing the stoichiometric neat resin to all nanocomposite compositions, the
clear benefit to material strength and modulus of adding excess amine in the epoxy resin
composition is observed. Table 6.10 lists the change in nanocomposite properties relative
to the stoichiometric neat resin.
In comparison to the base, stoichiometric resin, the most significant property improvements are seen with 5% excess amine added to the composition. In all cases, we see a trade off between modulus and ductility. The nanocomposites prepared containing
5% excess amine offer a dramatic increase in modulus, but correspondingly reduce the ductility by a significant amount.
Unlike the ATI samples, comparing FGS nanocomposites to the stoichiometric base resin demonstrates a real benefit to the addition of excess amine. The FGS is functionalized to a far greater extent than ATI and offers a significant disruption to the stoichometry of the resin. Therefore, reaction between excess amine and FGS, allows covalent bonding between matrix and filler, but also allows the resin cure to proceed with the appropriate amount of curing agent.
Within the FGS/ 0% excess amine series (stoichiometric ratio), comparing 0.5 wt% and 1.0 wt% nanocomposites to that of the base resin of stoichiometric ratio, we see only
162
Table 6.8. Results of ATI tensile tests, comparing all nanocomposites to the
corresponding neat resin
Adherent Technologies 0.5 wt% ATI Graphite 1.0 wt% ATI Graphite
Material E (psi) σy(psi) εfailure Toughness E (psi) σy(psi) εfailure Toughness (psi) (psi) Neat Resin (5% excess 1792 3990 30 1530 +/- 1792 3990 30 1530 +/- amine) +/- 103 +/- 130 +/- 1 20 +/- 103 +/- 130 +/- 1 20 5% excess 1272 2943 62 1450 +/- 1096 2534 60 1530 +/- amine +/- 15 +/- 70 +/- 3 20 +/- 101 +/- 51 +/- 4 20 Neat Resin (10% excess 941 +/- 2662 53 1270 +/- 941 +/- 2662 53 1270 +/- amine) 20 +/- 137 +/- 2 60 20 +/- 137 +/- 2 60 10% excess 1000 2608 67 1670 +/- 1170 2684 63 1480 +/- amine +/- 80 +/- 30 +/- 4 70 +/- 89 +/- 66 +/- 1 20
163
Table 6.9. Results of FGS tensile tests, comparing all nanocomposites to the
stoichiometric neat resin
FGS 0.5 wt% 1.0 wt%
Material E (psi) σy(psi) εfailure Toughness E (psi) σy(psi) εfailure Toughness (psi) (psi) Neat 1124 2568 45 +/- 1270 +/- 1124 2568 45 1270 +/- Resin +/- 53 +/- 6 8 90 +/- 53 +/- 6 +/- 8 90 0% excess 1432 2995 33 +/- 1220 +/- 1129 2140 32 950 +/- amine +/- 247 +/- 95 4 70 +/- 360 +/- 0 +/- 6 70 5% excess 1878 3350 12 +/- 1955 3875 7 +/- 400 +/- amine +/- 51 +/- 50 1 830 +/- 10 +/- 128 +/- 275 2 80 10% excess 1441 2670 44 +/- 1380 +/- 1262 3175 15 660 +/- amine +/- 196 +/- 0 1 10 +/- 9 +/- 75 +/- 1 20
164
Table 6.10. Change in nanocomposite mechanical properties relative to stoichiometric
base resin
Sample ∆σy ∆E ∆εfailure
0.5 wt% FGS, + 16% + 27% within error Stoichiometric 1.0 wt% FGS, - 17% within error within error Stoichiometric 0.5 wt% FGS, + 30% + 67% - 73% 5% excess amine 1.0 wt% FGS, + 51% + 74% - 84% 5% excess amine 0.5 wt% FGS, + 4% + 28% within error 10% excess amine 1.0 wt% FGS, + 24% + 12% - 67% 10% excess amine
165 enhancement of mechanical properties in the 0.5 wt% FGS sample. The drop off in both
E and σy on increased loading is an indication that the FGS does interrupt the stoichometry of the resin.
With respect to the corresponding neat resins properties, the 5% excess amine
nanocomposite exhibits only a marginal improvement in mechanical properties, as shown
in Table 6.11. It should be noted here however, that mechanical properties of the base,
resin are exceptionally high. Within the 10% excess amine series, the yield stress and
modulus increase with increased FGS loading. For example, at 0.5% loading, E and σy increase by 70% and 6%, respectively. At 1.0 wt% loading, E and σy increase by 32%
and 29%, respectively.
6.5 Discussion
With regard to mechanical properties, addition of EG to the epoxy resin offered the
smallest improvement. The increase in strength is negligible and the increased modulus
is accompanied by a strain to failure that is reduced to a much greater extent than any
other nanocomposite studied in this work.
The poor mechanical properties of the EG nanocomposites can be attributed to poor
dispersion, resulting in low surface area of the graphite particles. However, the low level
of oxidation implies that even if the amount of surface area was increased, there would be
no covalent bonding and little intermolecular interaction with the epoxy resin. The
resulting weak interface, as evidenced by SEM images of the fracture surface, is
illustrated in Figure 6.15.
166
Table 6.11. Tensile test data of nanocomposites and neat resins prepared with excess
amine curing agent
FGS 0.5 wt% 1.0 wt%
Material E (psi) σy(psi) εfailure Toughness E (psi) σy(psi) εfailure Toughness (psi) (psi) Neat Resin 1792 3990 30 1530 +/- 1792 3990 30 1530 +/- (5% excess +/- 103 +/- 130 +/- 1 20 +/- 103 +/- 130 +/- 1 20 amine) 5% excess 1878 3350 12 830 +/- 1955 3875 7 +/- 400 +/- amine +/- 51 +/- 50 +/- 1 10 +/- 128 +/- 275 2 10 Neat Resin (10% excess 941 +/- 2662 53 1270 +/- 941 +/- 2662 53 1270 +/- amine) 20 +/- 137 +/- 2 60 20 +/- 137 +/- 2 60 10% excess 1441 2670 44 1380 +/- 1262 3175 15 660 +/- amine +/- 196 +/- 0 +/- 1 10 +/- 9 +/- 75 +/- 1 10
167 On the other hand, addition of ATI and FGS to the epoxy resin produced mixed results
in terms of mechanical property enhancements, in particular with respect to the addition
of excess amine to the resin mixture. The conflicting data should be expected as the two
materials differ in both oxidation and dispersion. Addition of ATI graphite to the
stoichiometric resin does provide for enhanced mechanical properties. Significant
improvements in the material strength and stiffness are achieved, with only minimal loss
of ductility. However, the addition of excess amine reduces the material properties.
Because the ATI process provides functionalization only to the outermost layers of the
graphite aggregates, the presence of the graphite does not disrupt the resin stoichometry
to the extent that the mechanical performance is negatively impacted. The partial
functionalization is also the reason why SEM images of the ATI-epoxy fracture surface
reveal areas of well bonded graphite, and areas with visible voids at the graphite- resin interface, as illustrated in Figures 6.16a and 6.16b.
The fracture surface of FGS nanocomposites show clearly visible voids in the
specimens prepared with a stoichiometric ratio of amine curing agent. However,
following the addition of excess amine to the composition, only well embedded graphite
is observed within the matrix, as illustrated in Figures 6.17a and 6.17b.
6. 6 Conclusions
The extent of graphite oxidation has a significant effect on the level of dispersion
attained within the epoxy nanocomposites. Increasing the extent of oxidation allows
improved wetting by the epoxy matrix, resulting in well separated graphene layers. The
168
(a)
(b)
Figure 6.15. Representative SEM images of the fracture surface of (a) 0.5% EG in epoxy,
and (b) 3.0% EG in epoxy.
169
Figure 6.16. Fracture surface of ATI nanocomposites containing 0.5 wt% graphite and
either 10% excess amine (top), or 0% excess amine (bottom)
170
Figure 6.17. Fracture surface of FGS nanocomposites containing either 0% excess amine
(top), or 10% excess amine (bottom)
171 processes required to attain increased levels of oxidation also leads to reduced graphite
particle size, which aids in dispersion.
However, increased oxidation can have a negative impact on electrical conductivity,
since this can lead to the introduction of defects within the graphite structures, reducing
electronic conjugation. In this case, the highly oxidized, but well dispersed FGS affords
nanocomposites with conductivity behavior similar to that of EG nanocomposites, where
the EG nanoparticles are poorly dispersed but are only mildly oxidized. The material
imparting the greatest conductivity to the epoxy matrix is ATI, where an intermediate
level of oxidation and functionalization was attained.
Oxidation of ATI and FGS provides a means for further tailoring of the graphite
epoxy interface. In these systems, the excess amine curing agent which is added to the
epoxy resin allows for covalent bonding between graphite and the epoxy matrix. The
excess amine is unnecessary in the ATI series, due to the low level of epoxy
functionalization on the graphite layer surface. The FGS does, however, benefit from the
added curing agent. The presence of epoxide groups on the surface of the FGS
“unbalances” the stoichometry of the epoxy formulation. Therefore, the addition of
excess amine is necessary to “re-balance” the formulation, as well as create a strong
interface. Strengthening the interface translates into greatly enhanced mechanical
properties.
It is not possible to say from this data that one graphite is “better” than another.
However, the processes available to oxidize graphite nano-particles are growing in number and improving in efficiency. Controlling the level of oxidization within the graphite nanoparticle is becoming routine and approaching industrial scales. Therefore,
172 knowledge of how the degree of functionalization affects the processing and properties should become a prerequisite to designing a polymer-graphite nanocomposite.
173
CHAPTER VII
CONCLUSIONS AND FUTURE WORK
This work identified methods to optimize the interactions between a nanoparticle and resin system. It was demonstrated that the surface chemistry of a nanoparticulate filler exerts significant influence over the processing and properties of polymer based nanocomposites. The many factors which contribute to the mechanical and thermal behavior of polymer nanocomposites containing either layered silicate clays or nano- graphite materials were described, such as preferential contact between the particulate and the rigid section of an epoxy blend, or the extent of functionalization of graphite nanoparticles.
The nano-particles described were dispersed in a high temperature polyimide and a series of epoxy resins. The data showed that regardless of the nano-particle or the polymer matrix, the nature of the filler surface chemistry dictated the properties that were attained. In all cases, understanding the chemistry and composition of the filler-matrix interface allowed optimization of the bulk nanocomposite.
7.1 Layered Silicate Clay Nanocomposites
The work with layered silicate clays demonstrated a preferential affinity of the clay layers for aromatic compounds, relative to aliphatic based materials. This was influential
174 in the behavior of both high temperature polyimide nanocomposites, and epoxy-based
nanocomposites.
Within the high temperature polyimide series, the apparent affinity between an
aromatic modifier and the silicate surface corresponded with an increased oligomer melt
viscosity and a reduced polymer crosslink density. This was detrimental to the thermal
oxidative stability of the bulk nanocomposite. This work utilized a novel approach to
silicate modification by ion exchange with an aliphatic component, in addition to an
aromatic. The XRD showed an 0.25 nm increase in d-spacing, suggesting that the
modifiers were extending away from the silicate surface. Furthermore, the increase in
silicate d-spacing following co-modification of the aliphatic and aromatic components
improved the extent of clay layer separation that was observed through TEM
characterization. The modification system yielded reduced oligomer melt viscosity and
allowed crosslinking to proceed as in the unfilled polyimide. This resulted in over a 30%
reduction in material weight loss on long term, high temperature aging.
Within the epoxy blend nanocomposites, the preference of the clay for aromatic over
aliphatic components led to inhomogeneous dispersion of the clay throughout the blend,
as was evidenced by SEM images. Such behavior had not been previously noted in a
blend material. In systems where an excess aromatic material was utilized, the clay
tended to reside in close contact with the aromatic rich regions of the blend. This
resulted in only marginal enhancements to the nanocomposite strength, stiffness, Tg, and
CTE, as the silicate was essentially reinforcing the already rigid regions of the blend.
Therefore, the clay was forced into the flexible, aliphatic component by a pre-swelling technique, and the silicate reinforcement was directed to areas of the blend where the
175 silicate strength and stiffness would be most beneficial. As noted in Chapter 5, the pre-
swelling technique has been used for purposes of enhanced dispersion, however not for
directing placement of the silicate layers within the matrix. The result was a dramatic
improvement in the mechanical properties of the material. The pre-swelling process
additionally increased the Tg and reduced the CTE of the bulk nanocomposite.
7.2 Graphite Nanocomposites
The work with graphite nanocomposites demonstrated the material property trade-offs
observed when dispersing graphite nanoparticles into an epoxy matrix. As work with EG
is relatively new in the literature, there has not been a basic study comparing such
property trade offs as resulting from various degrees of particle functionalization and
dispersion. Primarily, this work determined that a certain degree of oxidation or
functionalization was necessary to allow the graphite to be wetted by an epoxy matrix.
Such wetting led to a strong interface as well as aided in dispersing the graphite
throughout the matrix. However, materials of increasing levels of oxidation negatively
impacted the electrical conductivity of the material. The increased oxidation also
affected the stoichometry of the epoxy material. The mechanical properties of the most
heavily oxidized particle, FGS, were optimized by the addition of excess amine curing
agent to the epoxy system. Unfortunately, electrical conductivity of these systems
required loadings greater than 1 wt%. Such high loadings introduced processing
difficulties due to the rapid increase in resin viscosity.
The ATI graphite was moderately functionalized and therefore did not affect the
stoichometry of the epoxy system. In this case, optimum mechanical performance was
176 achieved without excess amine curing agent. The electrical conductivity of the material was also greatly improved at only 0.5 wt%.
7.3 Future Work
There remains a need for continuous nanocomposite development for high temperature applications, as well as for toughened epoxy systems. The focus within the high temperature polymer community has shifted to resins that may be processed by resin transfer molding (RTM). This places several restrictions on the influence of nanoparticles on polymer processing. Primarily, the increase in viscosity that results from the nano-particulate fillers should not exceed the maximum viscosity allowable for
RTM processing. Furthermore, it is desirable that the silicate be well dispersed, i.e. exfoliated, within the imide oligomer melt. This may minimize any filtering of the clay nanoparticles on a carbon fiber perform, and ensure homogenous silicate dispersion throughout the composite. Additionally, the goal use temperature of polyimides considered for high temperature applications tends to increase. Therefore, identification of suitable organic modifiers for polyimides should continue to be investigated.
The interest in epoxy-clay nanocomposites continues to be in toughening the matrix material for cryogen storage tank applications. This is difficult as, generally, the increase in strength and stiffness that accompany silicate dispersion tends to reduce the resin ductility. However, this work has demonstrated that ensuring clay mobility in the matrix, and reinforcing of the toughening component of a blend may offer significant benefit to the material properties.
Future work with graphite nanocomposites will address truly optimizing the multifunctional nature of the graphite nanoparticle. This can be achieved, for example,
177 through alignment of graphite nano-particles. Alignment of the particles provides the
opportunity to direct material properties to provide a specified function, such as increased composite compressive strength, or directed conductivity.
178
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