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NATURAL AND SYNTHETIC MINERAL SILICATES AS FUNCTIONAL
NANOPARTICLES IN POLYMER COMPOSITES

By

Hua Shao

A DISSERTATION
Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of
DOCTOR OF PHILOSOPHY
Chemistry

2009

ABSTRACT

NATURAL AND SYNTHETIC MINERAL SILICATES AS FUNCTIONAL
NANOPARTICLES IN POLYMER COMPOSITES

By
Hua Shao

A new strategy is described for the substantial enhancement of the barrier
properties for both a therrnoset epoxy polymer and a thermoplastic polyolefin by
sandwiching a novel self-supported clay fabric film between thin polymer sheets.
The success of this strategy is attributed to the high orientation of clay
nanolayers in the paper-like clay fabric films and to the filling of the micro— or sub-
micro sized voids between imperfectly tiled clay platelet edges by the polymer
chains.

Thermoplastic polyolefin-clay fabric film composites were fabricated by
hot-pressing the clay films between two sheets of high density polyethylene
(HDPE) films. The sandwiched composites exhibit more than a 30-fold decrease
in 02 transmission rate with respect to the pure HDPE film. lmpregnating the self-
supported clay papers with epoxy pre-polymers successfully leads to therrnoset
composite films with more than 2-3 orders of magnitude reduction in 02
permeability in comparison to the pristine epoxy matrix.

Owing to the promising use of synthetic Mg-saponite (denoted SAP) as
epoxy polymer reinforcing agents, we investigated the cost-effective synthesis of
SAP by replacing urea with sodium hydroxide as base source. Co-crystallization
0f new zeolite phases, such as garronite (denoted GIS) and cancrinite (denoted

CAN). occurred along with SAP upon increasing the alkalinity of the reaction

mixture. This finding represents the first example of the preparation of a
CAN/SAP phase mixture. Moreover, pure-phase cancrinite with rod-like
morphology up to several pm in length was synthesized under Mg-free conditions.
Also, the Si/Al ratio within the synthesis gel has an influence on the chemical
composition and textural properties of pure CAN crystals.

Microporous cancrinite is a promising candidate for reinforcing epoxy
polymers, considering that CAN represents a substantial fraction of the mixed
CAN-SAP phase formed during the synthesis of saponite. Therefore, the reaction
conditions (e.g. alkalinity, reaction temperature, and duration) were further
expored in order to optimize the formation of small crystals of CAN (~ 100 nm)
with large external surface area and high pore volume. These textural features
facilitate the homogeneous dispersion of cancrinite particles in epoxy matrices. In
addition, a new phase sodalite (denoted SOD) was co-crystallized along with
CAN in some cases, depending on reaction conditions.

The synthetic zeolites CAN and SOD act as reinforcing agents for
therrnoset epoxy polymers owing in part to their small particle size, large surface
area and high pore volume. The tensile strength, modulus and toughness of an
epoxy matrix are simultaneously enhanced without the need for organic surface
modification of zeolites. This finding represents the first example for which the
mechanical properties of glassy epoxy composite are all improved through the

use of a microporous zeolite as a functional inorganic nanoparticle.

ACKNOWLEDGEMENT

First, I would like to express my appreciation to Professor Thomas J.
Pinnavaia for his thoughtful and kind guidance for the last five years. I learned a
lot of things from him that I need to be an independent researcher. I feel
extremely lucky to be his student.

Also, I would like to thank former and current group members. Dr. Park
had great patience in teaching me how to use the MOCON instrument and to
prepare epoxy-clay film composites. Dr. Xue helped me a lot by showing me how
to synthesize saponite and make glassy epoxy nanocomposites. Dr. Sun has
always been willing to help me when I got troubles especially my early years in
the US. Dr. Kim, Dr. Zhang and Dr. Wang are very nice to me and they are really
diligent and smart people. I also enjoyed the experience of working with Joel for
the Ford project back in 2007. My close friend Gwen has always shared good
and bad things happening to me.

I could not have finished my PhD study without the love and support from
my parents, Mr. Jingchao Shao and Mrs. Xiaoyun Zhou. I would also like to

dedicate this humble thesis to my beloved husband, llhwan An.

TABLE OF CONTENTS

LIST OF TABLES ................................................................................................. ix

LIST OF FIGURES ............................................................................................... x

CHAPTER 1

AN OVERVIEW OF POLYMER-BASED NANOCOMPOSITES1
1.1 Clay ...................................................................................................... 1
1.2 Polymer-Clay Nanocomposites ............................................................ 3
1.3 Barrier Properties of Polymer-Clay Nanocomposites ......................... 10
1.4 Polymer-Carbon Nanotube Nanocomposites ..................................... 14
1.5 Polymer-Zeolite Nanocomposites ....................................................... 20
1.6 Research Goals and Significance ...................................................... 26
References ............................................................................................... 30

CHAPTER 2

THERMOPLASTIC POLYOLEFlN-CLAY FABRIC FILM COMPOSITES WITH
HIGH BARRIER PROPERTY TOWARDS OXYGEN AND PETROLEUM FUEL

COMPONENTS .................................................................................................. 34
2.1 Introduction ........................................................................................ 34
2.2 Experimental ...................................................................................... 37
2.2.1 Materials ............................................................................... 37
2.2.2 Preparation of Homoionic and Mixed-ion Clay
Fabric Films .......................................................................... 37
2.2.3 Hot-pressing the Organoclay Film with HDPE Sheets .......... 37
2.2.4 Characterization and Testing of Clay Films and
HDPE-Clay Film Composites ................................................ 38
2.3 Results and Discussion ...................................................................... 39
2.3.1 Composition and Structural Properties of Homoionic and
Mixed-ion Clay Fabric Films ................................................. 39
2.3.2 lnterfacial Adhesion and Oxygen Permeability of HDPE-
Clay Fabric Film Composites ................................................ 44
2.3.3 Permeability of HDPE-Clay Fabric Film Composites
toward Representative Petroleum Fuel Components ........... 49
2.4 Summary ............................................................................................ 54
References ............................................................................................... 56
CHAPTER 3
GAS BARRIER PROPERTIES OF TRANSPARENT EPOXY-CLAY FABRIC
FILM COMPOSITES ........................................................................................... 58
3.1 Introduction ........................................................................................ 58
3.2 Experimental ...................................................................................... 61
3.2.1 Materials ............................................................................... 61

3.2.2 Preparation of Homoionic and Mixed-ion Clay

Fabric Films .......................................................................... 61
3.2.3 Preparation of Specimens for Testing .................................. 62
3.2.4 Characterization and Testing of Clay Films and
Epoxy-Clay Film Composites ................................................ 62
3.3 Results and Discussion ...................................................................... 64

3.3.1 Composition and Structural Properties of Homoionic and
Mixed-ion Montmorillonite Clay Fabric Films and Their
Epoxy Composites ................................................................ 64
3.3.2 Composition and Structural Properties of Homoionic and
Mixed-ion Fluorohectorite Clay Fabric Films and Their

Epoxy Composites ................................................................ 78
3.3.3 Oxygen Barrier Properties of Epoxy-Clay Fabn‘c Film
Composites ........................................................................... 80
3.4 Summary ............................................................................................ 83
References ............................................................................................... 85
CHAPTER 4
SYNTHESIS AND CHARACTERIZATION OF CANCRINITE, SAPONITE AND
THEIR PHASE MIXTURES ................................................................................ 87
4.1 Introduction ........................................................................................ 87
4.2 Experimental ...................................................................................... 89
4.2.1 Materials ............................................................................... 89
4.2.2 Saponite Synthesis ............................................................... 89
4.2.3 Synthesis of Phase Mixtures of Saponite and Cancrinite ..... 90
4.2.4 Cancrinite Synthesis ............................................................. 90
4.2.5 Characterization ................................................................... 91
4.3 Results and Discussion ...................................................................... 92
4.4 Summary .......................................................................................... 110
References ............................................................................................. 113
CHAPTER 5
INFLUENCE OF THE ALKALINITY AND TEMPERATURE ON THE SYNTHESIS
OF CANCRINITE-SODALITE SYSTEM ........................................................... 115
5.1 Introduction ...................................................................................... 115
5.2 Experimental .................................................................................... 117
5.2.1 Materials ............................................................................. 117
5.2.2 Synthesis of Nitrate Cancrinite ........................................... 118
5.2.3 Synthesis of Phase Mixtures of Nitrate Cancrinite and
Nitrate Sodalite ................................................................... 118
5.2.4 Characterization ................................................................. 118
5.3 Results and Discussion .................................................................... 119
5.4 Summary .......................................................................................... 134
References ............................................................................................. 1 36
CHAPTER 6

vi

EPOXY -ZEOLITE COMPOSITES WITH ENHANCED TENSILE

PROPERTIES ................................................................................................ 138
5.1 Introduction ................................................................................... 138

6.2 Experimental ................................................................................. 139

6.2.1 Materials .......................................................................... 139

6.2.2 Cancrinite and Sodalite Synthesis ................................... 139

6.2.3 Preparation of Epoxy-Zeolite Composites ....................... 140

6.2.4 Characterization .............................................................. 140

6.3 Results and Discussion ................................................................. 141

6.4 Summary ....................................................................................... 151
References .......................................................................................... 153

vii

LIST OF TABLES

Table 1.1 Classification of some clay minerals ..................................................... 2
Table 1.2 Electrical and mechanical properties of the films ................................ 18
Table 1.3 Permeation properties of carbon/ZSM-5 membranes at 298K ............ 25
Table 2.1 Basal spacings of homoionic and mixed-ion montmorillonite and
fluorohectorite clay fabric films ............................................................................ 41
Table 2.2 Oxygen permeability data for HDPE—clay fabric film composites ........ 48
Table 3.1 Basal spacings of homoionic and mixed-ion montmorillonite and
fluorohectorite clay fabric films ............................................................................ 65
Table 3.2 Oxygen permeability data for epoxy-clay fabric film composites ........ 81

Table 4.1 Textural properties of saponite, cancrinite and CAN/SAP phase
mixtures ............................................................................................................ 105

Table 5.1 Synthetic conditions together with the textural properties of reaction

products ............................................................................................................ 123
Table 5.2 Relative intensities of four diffraction peaks for sodalite and cancrinite
phase mixtures .................................................................................................. 125
Table 6.1 Textural properties of synthetic zeolites ............................................ 142

Table 6.2 Tensile properties of glassy epoxy composites reinforced by nitrate
cancrinite and sodalite ...................................................................................... 144

viii

LIST OF FIGURES

Figure 1.1 Structure of 2:1 phyllosilicates. ........................................................... 2

Figure 1.2 Schematic illustrations of three types of polymer-clay nanocomposites:
(a) conventional composites, (b) intercalated nanocomposites, and (c) exfoliated
(delaminated) nanocomposites. The corresponding X-ray diffraction (XRD)
patterns for composites at low-angle region (1-10°) are also shown ..................... 4

Figure 1.3 Schematic illustration of the types of interfacial interactions occurring
in polymer-organoclay nanocomposites, including direct binding (adsorption) of
the polymer to the basal siloxane oxygens (type A), “dissolving” of the onium ion
chains into the polymer matrix (type B), and polymer binding to hydroxylated
edge sites (type C) ................................................................................................ 6

Figure 1.4 TEM image of the MA-functionalized-PP/2C18-montmorillonite (MMT)
nanocomposite structure at 6 wt% loading, showing the coexistence of
intercalated clay tactoids (A) and disordered exfoliated clay layers (B) ................ 7

Figure 1.5 Effect of laponite concentration on the mechanical properties of
thermoplastic elastomers. a, Representative stress-strain (0-8) curves of thin-film
nanocomposites containing between 0 and 20 wt% laponite. b, Initial elastic
modulus, E, increases with increasing laponite concentration, <19. c, Toughness,
U03 (defined as the area under the 0-5 curve up to 5:03, increases quadratically
with CD. The error bars in b, c are equal to one standard deviation (minimum of
five samples) plus the contribution based on the instrument imprecision ._ ............ 9

Figure 1.6 (a) A tortuous path model for a diffusing gas through a polymer filled
with highly oriented impermeable disks; (b) Individual clay nanolayer with aspect
ratio, a = UW ...................................................................................................... 11

Figure 1.7 (a) Effect of sheet orientation on the relative permeability of exfoliated
nanocomposites. The illustrations show the definition of the direction of preferred
orientation (n) of the silicate layers and the actual orientation (p), and three
values of the order parameter 8= -‘/2, 0, and 1. (b) Effect of incomplete exfoliation
on the relative permeability. The illustrations show the effect of having one-, two-,
and four-sheet aggregates dispersed throughout the matrix ............................... 13

Figure 1.8 SEM micrograph showing a network of SWNT aggregates on the
surface of a fractured nanotube-epoxy composite .............................................. 14

Figure 1.9 (a) Reaction scheme for the functionalization of SWNTs and (b)
Integration of the functionalized SWNTs into an epoxy matrix ............................ 15

Figure 1.10 Chemical structures of the polymer matrix Parmax and poly(p-
phenylene ethylene)s (PPEs) used for the surface modification of carbon
nanotubes ........................................................................................................... 17

Figure 1.11 (a) Optical microscope image for composite containing 0.05 wt%
SWNT and 2 wt% clay under partial polarized light condition, and (b) Schematic
illustration of clay assisted dispersion of SWNTs ................................................ 20

Figure 1.12 Gas sorption isotherms at 35°C for unmodified zeolite L (ZL) and
modified zeolite L with APTES (APTES-2L): (a) nitrogen and (b) oxygen .......... 23

Figure 1.13 High-resolution TEM images of carbon/ZSM-5 composite membrane
with 9.1 wt% of ZSM-5 (denoted PIZZ): (a) ZSM-5 embedded in membranes, (b)
carbon phase in membranes .............................................................................. 25

Figure 1.14 Comparison between the framework topologies of cancrinite and
sodalite. The vertices represent the positions of Si and Al in the framework and
the edges represent the positions of bridging oxygen atoms .............................. 27

Figure 2.1 X-ray diffraction (XRD) patterns of (a) Na"-montmorillonite clay
powders, (b) Na*-montmorillonite clay fabric film and (c) Li*-fluorohector'rte clay
fabric film ............................................................................................................ 40

Figure 2.2 X-ray diffraction (XRD) patterns of (a) Na‘lBETmo-MMT, (b)
Na”/BET150-MMT, (c) Na‘lCTAHz-MMT and (d) cholinemo-MMT clay fabric films...
............................................................................................................................ 42

Figure 2.3 Therrnogravimetric analysis (TGA) curves for cetyltrimethylammonium
bromide powder, sodium montmorillonite and CTA" partially exchanged MMT
clay fabric films ................................................................................................... 44

Figure 2.4 (a) Schematic representation of orientation of the clay nanolayers at
the “air" and “glass” side of the clay fabric film and peeling off the HDPE coating
from the “air" side; (b) XRD of HDPE sheet removed from the “air" side of the
Na‘lcholinezs-MMT film ....................................................................................... 46

Figure 2.5 Permeation curves of HDPE-organoclay composite films and pure
HDPE polymer towards toluene, isooctane and ethanol vapor at 40°C .............. 51

Figure 2.6 XRD patterns of Na+/CTA*42-MMT clay fabric film before and after a
toluene permeation test ...................................................................................... 53

Figure 3.1 X-ray diffraction (XRD) patterns of (A) Na+ICTA+42-MMT clay fabric
film and (B) epoxy-clay film composite made from (A) ........................................ 67

Figure 3.2 X-ray diffraction (XRD) patterns of (A) homostructured mixed-ion
Na‘lGTABo-MMT clay fabric film and (B) epoxy-clay film composite made from (A)
............................................................................................................................ 69

Figure 3.3 X-ray diffraction (XRD) patterns of diprotonated diamine Jeffamine
02000 exchanged montmorillonite fabric films at different onium ion equivalents ..
............................................................................................................................ 70

Figure 3.4 Schematic representation of the orientations of the diprotonated
diamine Jeffamine D2000 onium ions intercalated in the mixed Na+-onium ion
smectite clay: (A) the diamine adopts an extended chain conformation in the
gallery with low d-spacing of ~ 16 A; (B) the diamine adopts a folded chain
conformation in the gallery with high d-spacing of ~ 40 A ................................... 71

Figure 3.5 X-ray diffraction (XRD) patterns of (A) heterostructured Na*/D2000*50-
MMT clay fabric film, (B) epoxy-clay film composite made from (A), only aged
under ambient conditions, (C) epoxy-clay film composite made from (A), further
cured at 125°C for 3h, (D) heterostructured Na*/02000*2o-MMT clay fabric film,
(E) epoxy-clay film composite made from (D) ..................................................... 72

Figure 3.6 TEM images of thin section of an epoxy-D200020 PGW composite film
showing the c-axis direction (cross-section): (a) low-magnification image showing
spacial distribution of clay aggregates; (b) exfoliated clay nanolayers in
composite film and (c) stacked clay layers in clay tactoids in composite film ...... 73

Figure 3.7 Optical images of (a) Epoxy-0200020 PGW film composites with a
thickness of 65 pm and (b) Epoxy-0200040 PGW film composites with a
thickness of 78 pm .............................................................................................. 75

Figure 3.8 X-ray diffraction (XRD) pattern of epoxy coating released from the
“air” side of Na*/02000*2o-MMT clay fabric film .................................................. 76

Figure 3.9 X-ray diffraction (XRD) patterns of homoionic organic and mixed
inorganic-organic ion fluorohectorite films ........................................................... 79

Figure 3.10 Optical images of (a) Epoxy-020006] FH film composites with a
thickness of 52 pm and (b) Epoxy-0200020 FH film composites with a thickness
of 53 um .............................................................................................................. 80
Figure 3.11 SEM micrograph of the cholinewo-FH fabric film ............................. 81
Figure 4.1 XRD patterns showing different crystalline product formation as a

function of increasing alkalinity at 90°C. The ratios represent the stoichiometry of
SizAlegzNaOH in the starting solutions .............................................................. 92

xi

Figure 4.2 TEM image of synthetic SAP formed at 90°C after a reaction time of 1
day ...................................................................................................................... 93

Figure 4.3 XRD patterns of crystalline products formed upon decreasing the
Mg(N03)2 content of the reaction mixture. For each product, the molar ratios of
Si:Al:Mg:NaOH in the synthetic solutions are labeled, as well as the reaction time
at 90°C ................................................................................................................ 94

Figure 4.4 TEM images of CAN/SAP phase mixture formed from a reaction
stoichiometry of SizAleg:NaOH=3.6:0.4:3:30 at 90°C after a reaction time of 4
days, and the corresponding ratios of SilAl=9, Mg/(Si+AI)=0.75, and OH'
/(Si+AI+Mg)=4.3 .................................................................................................. 96

Figure 4.5 TEM images of CAN/SAP phase mixtures formed at 90°C after a
reaction time of 3 days. (a) and (b) are for the CAN/SAP mixture made from
starting materials with the composition Si:Al:Mg:NaOH=3.6:0.4:1.5:30, and the
corresponding ratios of SilAl=9, Mg/(Si+Al)=0.375, and OH'/(Si+AI+Mg)=5.45; (c)
and (d) are for the CAN/SAP mixture made from SizAleg:NaOH=3.6:O.4:0.75:30,
and the corresponding ratios of SilAl=9, Mg/(Si+Al)=0.188, and OH'
/(Si+Al+Mg)=6.32 ................................................................................................ 97

Figure 4.6 TEM images of single-phase cancrinite formed from a reaction
stoichiometry of SizAleg:NaOH=3.6:0.4:0:30 at 90°C after a reaction time of 7
days, and the corresponding ratios of SilAl=9, Mg/(Si+Al)=0, and OH'
/(Si+Al+Mg)=7.5 .................................................................................................. 98

Figure 4.7 2"Si and 27Al MAS NMR spectra of (a) SAP formed after 1 day
reaction time (3.6:0.4:3:5), (b) CAN/SAP phase mixture formed after a reaction
time of 4 days (3.6:0.4:0.75:30) and (0) pure CAN made after 3 days
(3.6:0.4:0:30). The ratio in the parentheses represents the molar ratio of
Si:Al:Mg:NaOH in the starting solution. The asterisks indicate spinning sidebands
.......................................................................................................................... 100

Figure 4.8 TEM image and X-ray energy dispersive spectrum (XEDS) of
cancrinite crystals in the CAN/SAP phase mixture formed from a reaction mixture
with the composition SizAleg:NaOH=3.6:0.4:0.75:30 after 4 days. Unlabeled
peaks are from copper grid and carbon film. Quantitative analysis result is
summarized in the inserted table and note the Si/Al=1.25 ................................ 102

Figure 4.9 FTIR spectra of (a) pure saponite after 1 day of aging from initial
solutions with a SizAlegzNaOH stoichiometry at 3.6:0.4:3:5, and (b) CAN/SAP
phase mixture after 4 days of aging from initial solutions with a Si:AI:Mg:NaOH
stoichiometry at 3.6:0.4:0.75:30 ........................................................................ 103

Figure 4.10 N2 adsorption-desorption isotherms of pure saponite and cancrinite,
and their phase mixtures. The ratios represent the molar ratios of SizAlegzNaOH

xii

in the reaction mixtures. The reaction time for pure saponite is one day and all
the other products are made after a reaction period of 3 days .......................... 106

Figure 4.11 XRD patterns of cancrinite products formed as a function of the Si/Al
ratio within the starting system after 3 days of reaction time under Mg-free
condition ........................................................................................................... 107

Figure 4.12 2S’Si MAS NMR spectrum of pure cancrinite synthesized from the
educts with the composition of Si:Al:NaOH=1.6:0.4:30 under Mg-free condition
.......................................................................................................................... 108

Figure 4.13 TEM images of cancrinite crystals formed at Si:Al:NaOH=1.6:0.4:3O
under Mg-free condition .................................................................................... 109

Figure 5.1 XRD patterns of reaction products as a result of varying alkalinity in
the starting solutions after a reaction time of 3 days at 90°C. The ratios represent
the Si:Al:NaOHzH2O stoichiometry of the reaction system. (S=sodalite,
C=cancrinite) ..................................................................................................... 120

Figure 5.2 TEM images of (a), (b) and (c) ideal cancrinite crystals formed at 3
different alkalinities of x=75, 37.5 and 18.8, respectively; (d) cancrinite in
CAN/SOD phase mixture formed at x=9.38; (e) and (f) rod-like cancrinite and
cotton-ball like sodalite in CAN/SOD phase mixture formed at x=4.69. All the
samples were prepared at 90°C after a reaction time of 3 days ........................ 121

Figure 5.3 FTIR spectra of selected products formed at varying alkalinity in the
starting solutions after 3 days reaction at 90°C. The ratios represent the
Si:Al:NaOHzH2O stoichiometry in the reaction system ...................................... 126

Figure 5.4 29Si MAS NMR spectrum of pure cancrinite precipitated from the
educts with composition of Si:Al:NaOHzH2O=1 :1 :18.8:125 at 90°C after 3 days .....
.......................................................................................................................... 128

Figure 5.5 N2 adsorption-desorption isotherms of (a) pure cancrinite products
and (b) CAN/SOD phase mixtures, formed at different alkalinities after 1 day or 3
days of reaction period. The insert provides BJH pore size distributions
determined from the adsorption branches of the N2 isotherms of selected
products ............................................................................................................ 129

Figure 5.6 XRD patterns of products formed at different reaction temperatures.
(S=sodalite, C=cancrinite) ................................................................................. 131

Figure 5.7 TEM images of sodalite crystals prepared at 25°C after 3 days at high
and low magnifications ...................................................................................... 132

Figure 5.8 FTIR spectra of products as a result of changing reaction temperature

xiii

.......................................................................................................................... 133

Figure 5.9 N2 adsorption-desorption isotherms of samples formed as a function
of reaction temperature after 3 days reaction time. The insert provides the BJH
pore-size distribution determined from the adsorption branch of the N2 isotherm
of SOD/CAN@25°C phase mixture ................................................................... 134

Figure 6.1 XRD patterns of synthetic zeolites after a reaction time of 3 days ..142

Figure 6.2 N2 adsorption-desorption isotherms of synthetic zeolites after a
reaction time of 3 days ...................................................................................... 143

Figure 6.3 Representative stress-strain curves for glassy epoxy nancomposites
reinforced by (a) CAN-18.8-90°C and (b) SOD-18.8—25°C at the wt% zeolite
loadings indicated ............................................................................................. 145

Figure 6.4 Optical image comparing the transparency of (a) pristine epoxy
polymer and epoxy composites prepared from (b) CAN-18.75-90°C, (0) CAN-
37.5—90°C and (d) SOD-18.8-25°C. All of the composites are at 5 wt% zeolite
loading .............................................................................................................. 148

xiv

CHAPTER 1

An Overview of Polymer-Based Nanocomposites

1.1 Clay

Smectite clays are a class of layered silicate minerals characterized by
various properties such as a very small particle size, large ion-exchange and
adsorption capacities, a low hardness index, and a high particle aspect ratio.1
The elementary polyhedral building blocks of layered silicates are tetrahedra and
cctahedra of oxygen ions. The coordinating cations in the center of the
tetrahedra are essentially Si, Al, or Fe3", among which Si is the most frequent.
Octahedrally coordinated cations are usually Al, Mg, Fe3+ and Fe”.

The structural unit formed by linking one octahedral sheet to one
tetrahedral sheet results in a 1:1 layer, whereas the unit formed by sandwiching
one octahedral sheet with two tetrahedral sheets affords a 2:1 layer, as shown in
Figure 1.1.2 The exposed surfaces of 1:1 layers consist of oxygen atoms on the
tetrahedral layer surface, and —OH groups on the octahedral layer side. Only
bridging oxygen atoms are exposed on the basal surfaces of 2:1 layers. A brief
summary of the major 1:1 and 2:1 clay groups is shown in Table 1.1.1

1:1 and 2:1 layers are not always neutral. Very often, as a result of
isomorphic metal ion substitution in the lattice, the layers bear a net negative
charge that is balanced by hydrated cations situated inside the galleries. The

galleries cations (Na+ or K”) can be displaced by alkylammonium ions through

   

A ~—-Tetrahedral

~Inm _ “ ‘
E v «n—- Octahedral
g <—-— Tetrahedral

   

 

 

 

 

0 0 O G 0 0 a
Exchangeable Cations“ 0 Al, Fe, Mg, Li
" OOH
.0
OLi, Na, Rb, Cs
Figure 1.1 Structure of 2:1 phyllosilicates.
Table 1.1 Classification of some clay minerals
Structure Mineral
Charge Formula
type group
1:1 layer kaolinite ~0 [AI4](SI4)O10(OH)3
2:1 layer talc ~0 [Mge](Si3)020(OH)4

smectite 0.2-0.6 M"*x'yH2O[(AI, Mg, Fe, Zn)4]((Si, AI)8)020(OH)4
vermiculite 1.2-1.3 M92+x[(Mg, Ca, K, Fe2*)6]((Si, AI, F63+)3)020(0H)4

 

ion-exchange reaction to fine-tune the surface chemistry of clay layers, which
converts the normally hydrophilic silicate surface to an organophilic one. Organic
modifiers can lower the surface energy of the nanoclay particles and improve the

wetting behavior with the polymers. In some cases, cationic surfactants can

provide functional end groups that can react with the polymer matrix3 or catalyze
the polymerization of monomers“’6 to improve the strength of interfacial
interactions between the inorganic silicates and the polymer matrix.

The most commonly used layered silicates for the preparation of polymer-
clay nanocomposites are smectite clays, which belong to the 2:1 phyllosilicates
(Figure 1.1). Typical members of the family are montmorillonite (MMT), hectorite
and saponite. The substitution of Al3+ by Mg” results in negative charges on the
octahedral sheet in layers of montmorillonite. The negative charge centers are
counterbalanced by hydrated sodium or calcium cations in the interlayer between
2:1 silicate sheets. The layer thickness is around 1nm, and the lateral dimensions

may vary from 30 nm to several microns or lager, providing a high aspect ratio.

1.2 Polymer-Clay Nanocomposites

Based on the dispersion of clay particles in the polymer matrix, polymer-
clay nanocomposites can be divided into three categories as illustrated in Figure
1.27. For the “conventional composites”, there’s no intercalation of polymer
chains into the clay layers and the clay platelets remain as a tactoid structure.
Therefore, the X-ray diffraction (XRD) pattern shows the same d-spacing as the
pristine clay. Due to the limited interfacial interaction, phase segregation normally
occurs in this case. For the “intercalated nanocomposites”, polymer molecules
penetrate into the clay gallery and correspondingly expand the gallery height.
The gallery height expansion is reflected by an increase in the X-ray d-spacing

along the layer stacking direction. In the case of “exfoliated or delaminated

polymer-clay nanocomposites", individual clay nanolayers are separated by the
polymer, as verified by the disappearance of the reflection peak in the low-angle
region of X-ray pattern. Exfoliated composites optimize the interfacial interaction
between the silicate layers and the polymer matrix and result in simultaneous

improvements in mechanical, thermal, flame resistance and barrier properties.

(a) Conventional

            
       
      
  
 
    
   
 

 
  
  

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26

Figure 1.2 Schematic illustrations of three types of polymer-clay nanocomposites:
(a) conventional composites, (b) intercalated nanocomposites, and (c) exfoliated
(delaminated) nanocomposites. The corresponding X-ray diffraction (XRD)

patterns for composites at low-angle region (1-10°) are also shown.

Nanocomposite studies have been carried out on various polymer
systems, including polyurethane, polyimide, polyamide, polystyrene,
polypropylene, epoxy, etc.25'&25 In the case of thermoplastic polyurethane3,
nanocomposites containing 5 wt% layered silicate organoclay exhibit a 125%
increase in tensile stress, 100% increase in elongation, and 78% increase in
tensile modulus along with a 130% increase in tear strength due to the
reinforcement provided by the full exfoliation of nanoclay layers.

Epoxy-nanoclay systems have been widely studied with regard to the
mechanism of clay tactoids intercalation and exfoliation. Earlier works"6 have
shown that the relative intra- and extragallery polymerization rates, along with the
epoxy curing temperature, play an important role in controlling the degree of clay
particle exfoliation. The exfoliation is facilitated by acidic primary alkylammonium
ions in the galleries that can catalyze the intragallery epoxy polymerization. But
Jana8 has observed that catalytic intragallery polymerization, though expedited
the exfoliation process, was not necessary for exfoliation. The entropic elastic
force exerted by the cross-linked epoxy molecules inside the clay galleries was
found to be responsible for the delamination of clay layers.

The synthetic methods and material properties of polypropylene-clay
nanocomposites have been widely studied due to the very large commercial
importance of polypropylene (PP). In terms of thermodynamics, the dispersion of
clay platelets (intercalation or exfoliation) was energetically favorable when the
polymer-silicate surface interactions were maximized and polymer-alkyl

d ’2627

surfactants interactions were minimize as illustrated in Figure 1.328.

 

    

 

Clay Layer

 

 

Figure 1.3 Schematic illustration of the types of interfacial interactions occurring
in polymer-organoclay nanocomposites, including direct binding (adsorption) of
the polymer to the basal siloxane oxygens (type A), “dissolving" of the onium ion
chains into the polymer matrix (type B), and polymer binding to hydroxylated
edge sites (type 0).28

Therefore, the incorporation of any polar or polarizable functional groups in
polypropylene polymer will promote PP compatibility with organoclays.“"2‘*23 For
instance, Manias and co-workers23 prepared maleic anhydride functionalized
polypropylene (PP-r—MA) nanocomposites by melting processing. A TEM image
of the composite containing 6 wt% organo-MMT clearly shows that there are
intercalated tactoids and disordered exfoliated silicate layers coexisting in the

nanocomposite structure (Figure 1.4).

 

Figure 1.4-TEM image of the MA-functionalized-PP/ZC18-montmorillonite (MMT)
nanocomposite structure at 6 wt% loading, showing the coexistence of

intercalated clay tactoids (A) and disordered exfoliated clay layers (B).23

The goal of MA functionalization was to improve the interactions between
the polymer and layered silicate surfaces so as to become thermodynamically
more favorable compared to the polymer-surfactants interactions. The same
effect can be achieved by decreasing the interactions between the polymer and
alkyl surfactants, which subsequently will render the PP/clay contacts favorable.
This approach can be realized by introducing appropriate organic modifiers, such
as semi-fluorinated surfactants.25 Specifically, octadecylammonium exchanged

montmorillonite was further modified by reaction with a semi-fluorinated

alkyltrichlorosilane surfactant (CF3-(CF2)5-(CH2)2-Si-Cl3), which was grafted to the
clay surface through an interaction of the trichlorosilane groups with hydroxyls in
the cleavage plane of the montmorillonite. An estimated 25% and 40% of the
resulting clay layers are exfoliated in unfunctionalized PP nanocomposites made
via melt-intercalation processing and through extrusion in a twin-screw extruder,
respectively.

It has been realized recently that unmodified synthetic inorganic clays can
disperse uniformly in polymer matrix and thus provide mechanical reinforcement
and improvement in thermal properties. Xue29 reported that synthetic porous Mg-
saponite3° improves the tensile properties of both rubbery and glassy epoxy
matrices without the need for the organic modification of the clay surface. The
improvements in tensile strength, modulus and toughness are superior to those
of orgaoclays, a result that is ascribed to the irregular stacking and textural
porosities of saponite nanolayers. The polymer fills the interparticle pores,
thereby resulting in favorable interfacial interactions and reducing the tendency
toward microcrack formation. Using a novel solvent-exchange method, McKinley
and co-workers successfully dispersed the fully exfoliated synthetic laponite
nanolayers in a thermoplastic elastomer, composed of polyurethane and
polyether block copolymer, and preferentially reinforced the polyurethane hard
microdomains.15 The stronger adhesion of the hydrophilic, unmodified laponite
with the more polar hard segments rather than with the soft segments allows the
laponite to segregate into the hard microdomains during processing, thus

resulting in the increase in the stiffness, toughness and heat distortion

 

315

12

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Yield point

 

 

  
   

 

 

 

 

 

 

 

 

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Figure 1.5 Effect of laponite concentration on the mechanical properties of

thermoplastic elastomers. a, Representative stress-strain (o-s) curves of thin-film

nanocomposites containing between 0 and 20 wt% laponite. b, Initial elastic

modulus, E, increases with increasing laponite concentration, <D. c, Toughness,

U23 (defined as the area under the 0-8 curve up to e=0.3, increases quadratically

with (D. The error bars in b, c are equal to one standard deviation (minimum of

five samples) plus the contribution based on the instrument imprecision.15

temperature of the material. The dramatic influence of laponite on the elastic
modulus of polyurethane is shown in Figure 1.5”. The initial modulus of the
elastomer increases significantly with increasing laponite concentration as this
region is governed by the polyurethane crystallinity. Specifically, the sharp
increase in modulus at CD > 6 wt% (Figure 1.5 b), is a consequence of the
percolative laponite network. After the yield point, the region of moderate slope is
indicative of soft-segment deformation and the steep upturn is characterized by
the strain-induced crystallization of the soft domains. The addition of laponite
does not affect the slopes of these two regions, which can be explained by the

fact that there is no interaction between the smectite clay and soft microdomains.

1.3 Barrier Properties of Polymer-Clay Nanocomposites
The mechanism of barrier improvement for polymer-clay nanocomposites
was first explained by Nielsen’s tortuous path model (Figure 1.6), wherein the
exfoliated impermeable nanoclay layers are highly oriented and retard the
transport of diffusing species, such as oxygen and water molecules.31 The
tortuosity factor (I) is defined by the ratio of the actual distance that a diffusing
gas must travel to the shortest distance that it would have traveled in the
absence of the layered silicate and is expressed in terms of the length (L), width
(W), and volume fraction of the clay (CD) as
r = 1+(L/2W) (D = 1+ 0 43/2 (Equation 1.1)
where a is the aspect ratio of the silicate clay. Relative permeability is defined by

the ratio of the permeability of the composite to that of the pristine polymer:

10

POIPp = 1 /[1+(U2W) ¢] = 1 /(1+ (1 613/2) (Equation 1.2)
where Pc and Pp represent the perrneabilities of the composite and pure polymer,

respectively. It can be seen that the relative permeability PJPp equals 1/ r.

 

 

 

 

 

 

 

Figure 1.6 (a) A tortuous path model for a diffusing gas through a polymer filled
with highly oriented impermeable disks; (b) Individual clay nanolayer with aspect

ratio, a = UW.

In contrast to expectations, only a 2-4-fold reduction in permeability is
observed in practice. Nielson's model is based on a key assumption of full
exfoliation of clay tactoids and clay nanolayers highly oriented along the direction

perpendicular to the film surface. Neither condition is achieved in practice.

11

Therefore, the effect of clay platelet dispersion and orientation on the barrier
properties of polymer-clay nanocomposites has been studied in detail”:33

With regard to the effect of sheet orientation, the layer order parameter is
defined as

S = ‘/2(3cos26-1) (Equation 1.3)
where 9 represents the angle between the direction of preferred orientation (n)
and actual sheet orientation (p) as shown in Figure 1.7 (a)32. This function can
range from 1 (0=0), indicating perfect orientation of platy clay layers, to -‘/2
(9=1tl2), indicating perpendicular or orthogonal orientation, and a value of 0 for
random orientation. The tortuosity factor is modified further to include the platelet
orientation order 8. The resulting relative permeability is given by
POIPp = 1 l[1+(U2W)<I>(2l3)(S+‘/2)] (Equation 1.4)
The above expression reduces to Equation 1.2 when 8:1 and converges
approximately to the permeability of the pristine polymer when 8= -/2.

The effect of increasing aggregation while keeping a parallel platelet
orientation is illustrated in Figure 1.7 (b)32. For single clay layers dispersed in the
polymer matrix, the width is ~ 1 nm which allows for the maximum aspect ratio. In
contrast, consider a case where aggregates composed of two silicate layers and
an intercalated polymer layer are dispersed throughout the matrix. Now, the
width increases to ~ 3 nm assuming that the polymer layer is 1 nm thick.
Aggregates of four silicate sheets with three layers of intercalated polymer chains
lead to an effective width of ~ 7 nm, thus dramatically increasing the relative

permeability.

12

 

WI: 3nm

 

 

l‘ncreaslng delamination (exfonation)

Figure 1.7 (a) Effect of sheet orientation on the relative permeability of exfoliated
nanocomposites. The illustrations show the definition of the direction of preferred
orientation (n) of the silicate layers and the actual orientation (p), and three
values of the order parameter S= -‘/z, 0, and 1. (b) Effect of incomplete exfoliation
on the relative permeability. The illustrations show the effect of having one-, two-,

and four-sheet aggregates dispersed throughout the matrix.32

1.4 Polymer-Carbon Nanotube Nanocomposites

Carbon nanotubes (CNTs) have attracted special attention as an ideal
filler for polymer composites due to their exceptional mechanical properties (e.g.
high modulus, stiffness and extensibility),34'35 and high electrical36'37 and thermal
conductivity.38'39 Currently, a variety of nanotube types is available, such as
single-walled carbon nanotubes (SWNTs), double-walled carbon nanotubes
(DWNTs), and multi-walled carbon nanotubes (MWNTs). When nanotubes are
considered for reinforcement applications, it is important to address bundle
formation because of intertubular van der Waals forces (Figure 1.8).40 In order to
facilitate CNTs dispersion/exfoliation, surface functionalization of the nanotubes

is necessary in order to form strong interfacial interactions between the nanotube

41-45

and the polymer matrix.

 

Figure 1.8 SEM micrograph showing a network of SWNT aggregates on the

surface of a fractured nanotube-epoxy composite.40

14

o o o o 0
II II II II Heat II (a)
HOCCH2CHZCOOCCHZCHZCOH 70—» HOCCHZCHZ-
" 2

_ o . o _
to“? II ”6%. ll V6?- ||

§g+ HOCCHZCHZ- —» ."-{-CHZCH2C-OH n 393'? i3 {CHZCHZC-NH-X-NHZ
‘30. ‘90. 2. diamine €01 n

SWNT

Where "X" represents ——O— —O—

I—O—I

(b)

   

OH
CHz-CH-CHZW
CO-NH-X-h<
H? 2 CHz-CH—CHZW
WHZC'HC'HZCEN x NHCIIJCH CH H
WHZC-HC-HZC 2 2
OH
on I
CHz-CH'CHZW
o-NH-x-N<
CH2*CI:H-CH2‘W
OH

Figure 1.9 (a) Reaction scheme for the functionalization of SWNTs and (b)

Integration of the functionalized SWNTs into an epoxy matrix.“

Wagner and co-workers reported a significant enhancement of mechanical
properties of poly(vinyl alcohol) (PVA) composites with the addition of covalently

functionalized SWNTs with multiple surface hydroxyl groups. The improvements

15

are attributed to the strong hydrogen bonding at the polymer-nanotube
interface.45 In one study, a fully integrated nanotube-epoxy polymer composite
with covalent chemical bonding between the matrix and functionalized SWNTs is
reported.42 Functionalized nanotubes were prepared through the reaction of
diamines with alkylcarboxyl groups attached to the SWNTs sidewalls as shown in
Figure 1.9 (a). The terminal amino groups attached to the nanotubes readily
react with epoxy resin and thus serve as an additional curing agent. Therefore,
nanotubes are covalently integrated into the matrix and become part of the cross-
linked structure, rather than just a separate filler (Figure 1.9 (b)). At a loading of
only 1 wt% functionalized SWNTs, the resulting epoxy composites demonstrate
significant increase in tensile strength, Young’s modulus and elongation-at-break
by more than 25%, 30% and 30%, respectively. This also suggests higher
toughness with respect to the neat epoxy. These observations support the
theoretical and molecular simulation“6 predictions that load transfer from the
matrix to the nanotubes can be effectively increased by the formation of covalent
bonding at the interface. The simulations predict that the shear strength required
for load transfer can be enhanced by more than one order of magnitude with the
introduction of chemical bonds involving less than 1% of nanotube carbon atoms.

Although surface functionalization enables particle-matrix interactions to
be controlled through a variety of organic groups attached to the carbon
nanotubes, and as a consequence, to design nanotube-reinforced polymer
composites, covalently functionalized carbon nanotubes are not good candidates

for the improvement of composite conductivity, because strong tube-polymer

16

adhesion produces weaker nanotube-nanotube contacts, which weakens the

three-dimensional nanotube netvrrork.“7'49

 

I
C12HEU£OC12H5 g

F PPE 3
PPE 2 <5

Figure 1.10 Chemical structures of the polymer matrix Parmax and poly(p-

phenylene ethylene)s (PPEs) used for the surface modification of carbon

nanotubes.47

For the first time, Chen and co-workers suggested non-covalently

functionalized SWNTs as being crucial for achieving a simultaneous

17

Table 1.2 Electrical and mechanical properties of the films47

 

 

Film Electric conductivity Young’s modulus Tensile strength
sample (s-m'1 ) (GPa) (MPa)
A <10'11 3.9:02 14418
B 0.42 5.3 1: 0.8 155 i 9
C 0.56 5.9 :I: 0.7 211 1: 7
D 28.6 5.0 :I: 0.3 248 i 8

 

improvement in mechanical and electrical properties of SWNT-polymer
nanocomposites.47 They used Parmax (Figure 1.10), poly(benzoyI-1,4-phenylene)
-co-(1,3-phenylene), as the polymer matrix in this study because of its
extraordinary mechanical strength and stiffness, inflammability, and high-
temperature stability. They used three poly(p-phenylene ethylene)s (PPEs)
(Figure 1.10) to non-covalently functionalize SWNTs and disperse SWNTs
homogeneously in polymer matrix. Studies have shown that the backbone of
PPE modifiers can bind irreversibly to carbon nanotube surfaces by 1!:-
stacking.”51 Four different composite film samples were prepared for this study:
sample A is pure Parmax and samples B, C and D are Parmax composites with
the addition of 2 wt% PPE 1, PPE 2 and PPE 3 functionalized SWNTs,
respectively. As shown in Table 1.2, in comparison to sample B, sample C
demonstrates a similar enhancement in electrical conductivity, but a more
significant increase in tensile strength. The weak van der Waals’ interactions
exist for both samples, while the large, branched side chains in PPE 2 may
increase mechanical interlocking. Furthermore, the PPE 3/Parmax interface

exhibits the most dramatic effects on the tensile strength and electrical

18

conductivity of the PPE-SWNT—Pan'nax composite, suggesting the best adhesion
between PPE 3 and Parmax matrix as compared to that between PPE 1, 2 and
Parmax. This is achieved through the 1m: interactions between ferrocenyl groups
in PPE 3 and Parmax.

More recently, a new approach to make conductive nanotube-filled
polymer composites is to introduce clay into single-walled carbon
nanotube/epoxy composites to assist nanotube dispersion without harming
mechanical property.48 It was found out that with only 0.05 wt% SWNT, electrical
conductivity was increased by more than four orders of magnitude in combination
with 0.2 wt% clay. This improvement in conductivity was accompanied by a slight
increase in storage modulus. In addition, the percolation threshold for the
nanocomposites was reduced from 0.05 wt% without clay to 0.01 wt% with the
addition of clay. The most likely reason why clay helps SWNT network formation
is that the nanotubes seem to interact more strongly with the negatively-charged
montmorillonite layers than the epoxy polymer. This tube-clay interaction can be
observed using partially polarized light in an optical microscope in Figure 1.11 (a):
the epoxy matrix is in gray color, whereas clay is seen as bright clusters and
SWNT is black. Almost all of the clay tactoids are surrounded by a SWNT
network and the nanotubes are not observed in open areas that have little or no
clay. Clay-assisted dispersion of carbon nanotubes is illustrated in Figure 1.11
(b). The limited improvement in composite mechanical properties might be

caused by the inadequate exfoliation of the clay nanolayers.

19

 

Figure 1.11 (a) Optical microscope image for composite containing 0.05 wt%
SWNT and 2 wt% clay under partial polarized light condition, and (b) Schematic

illustration of clay assisted dispersion of SWNTs.48

1.5 Polymer-Zeolite Nanocomposites
Zeolites are crystalline microporous materials that are widely used as
catalysts for applications like the cracking of heavy oil fractions, cumene

production, alkane hydroisomerization and fine chemicals synthesis.52 Since the

20

19803, there has been a great deal of interest in the preparation, characterization
and application of polymer-zeolite composites.

Al-ghamdi and Mark53 studied reinforcements of poly(dimethylsiloxane)
(PDMS) by zeolite 3A with pore size of 3 A and zeolite 13X with pore diameter of
10 A, both with a cubic morphology and particle size of 1-5 pm. These two
zeolites were simply blended with PDMS and found to improve the mechanical
properties of the PDMS elastomers. But the effect was more dramatic for zeolite
1 3X with the larger pore size. Further characterization of such zeolite-reinforced
PDMS elastomers by Wen and Mark54 concluded that there was negligible
penetration of the polymer into the zeolite pores, as suggested by the small-
angle neutron scattering (SANS) results of the pure and zeolite-filled PDMS
networks. Also, the superior reinforcement achieved by zeolite 13X might be
attributed to its smaller particle size, rather than its larger pore size.

In the cases discussed above, the polymer is directly introduced into the
zeolite confined structure. A different approach is to introduce monomers and
subsequently polymerize monomers within channels to obtain the final composite
materials.5557 In the first application of this approach, Frisch and Xue55 added
monomeric styrene to a desired amount of zeolite 13X in various ratios, followed
by the addition of a benzoyl peroxide initiator for the polymerization of the styrene.
In the presence of ethylene glycol dimethacrylate (EGDM) as a cross-linker, the
cross-linked polystyrene (PS)/zeolite composite formed a fully interpenetrating
polymer network (IPN) as no significant amount of cross-linked polystyrene

chains was removed by extensive solvent extractions. In contrast, the linear PS

21

chains and zeolites were considered to form a pseudointerpenetrating polymer
network (PIPN) since half of the linear polymer could be removed after solvent
extraction. According to the differential scanning calorimetry (DSC)
measurements, an extracted sample of PIPNs exhibited no glass transition
temperature Tg because the linear polymer chains situated in the internal pores
of zeolite 13X, adopt a more extended, one dimensionally oriented conformation.
This interpretation is consistent with Giannelis’s observation that intercalated PS
chains in silicate clay layers do not show a T9 characteristic of the pure
polymer.58

Polymer-zeolite composites may provide economical, high performance
gas separation membranes if interfacial voids at the molecular sieve/polymer
interface can be eliminated.”63 Marand and co-workers fabricated mixed matrix
membranes (MMMs) composed of aminopropyltriethoxysilane (APTES) surface
modified zeolite l. (ZL) covalently linked to a polyimide matrix containing
carboxylic acid groups along the backbone.61 Unlike zeolites employed in other
MMM studies, zeolite L containing one-dimensional channel with an aperture of
7.1 A can separate gas molecules by selective adsorption in addition to size
exclusion. The channels contain cationic sites, which can interact with negatively
charged or polar molecules.64 As illustrated in Figure 1.12, the sorption capacity
of the APTES-2L is significantly lower than that of the pristine ZL, suggesting the
possibility of pore blockage by the coupling agents. 02 and N2 permeability
increases 13% and 24%, respectively, through the MMM when compared to the

pure polyimide membrane. However, this results in a slight decrease in the

22

 

— - d
P P O M A
n I n A A 1

Amount Adsorbed (mglgi
A

N
I A

 

 

I
2L I
I
I
I
I
I
I
. ' mesa.
e . , e .
'0' . e ‘ .
-I .. 0 °
'-
I
sooioooisoozomzsooaoooasoowoouoo

 

Pressure(mbar)

 

Amount Adsorbed (mg/g)
3

 

 

 

 

(b)

Figure 1.12 Gas sorption isotherms at 35°C for unmodified zeolite L (ZL) and
modified zeolite L with APTES (APTES-2L): (a) nitrogen and (b) oxygen.61

selectivity factor for O2/N2 from 4.8 to 4.5 because N2 is expected to have higher

sorption in zeolite L.

Yong et al. proposed a new method to fill the space between the polymer

and zeolite particles by introducing 2,4,6-triaminopyrimidine (TAP) to form

0 . 560 r1000'1500j2000125‘00‘30'oofi3500'4obofi4500

Pressure (mbar)

23

hydrogen bonds between the two phases instead of the surface treatment of
zeolites.62 The void-free Pl/zeolite 13 X/T AP membrane showed higher
permeability for He, N2, 02, CO2, and CH4 with little compromise in
pennselectivity as compared to the Pl/T AP membrane. The Pl/zeolite 4A/T AP
membrance showed lower permeability but higher selectivity for all gases. These
results could be explained by the different pore sizes in 13X and 4A zeolites. The
7.4 A pore size of zeolite 13X, appears to be large enough for the facile
penetration of all gases compared to the Pl/T AP matrix. However, the 3.8 A
diameter of zeolite 4A is very close to the kinetic diameter of the gas molecules,
which may provide the molecular-sieving effect.

A novel carbon/ZSM-5 nanocomposite membrane with high perfonnace
for oxygen/nitrogen separation is prepared by incorporating nano-sized ZSM-5
into the polyimide precursor.63 Carbon membranes have an ultra-micropore
structure with pore diameters of 3-5 A, close to the diameters of O2 and N2. Thus,
they are good candidates for air separation. But there’s a strong trade-off
relationship between the gas flux and selectivity because of the disordered pore
structure and diffusion resistance in membranes. Zeolite ZSM-5 features uniform
pores (5.5 A) and can endure high temperature treatment. Therefore, the
channels in zeolite will remain intact to some degree and function as continuous
transport pathways for gas species as shown in Figure 1.13. Table 1.3 shows the
permeation properties of the pure carbon membrane and the carbon
nanocomposite membranes incorporated with different zeolite loadings.63 It’s

clear that the perrneabilities of O2 and N2 increase substantially at higher zeolite

24

loadings, indicating that porous zeolite ZSM-5 helps to minimize the resistance in
the gas diffusion process. In addition, the composite membranes retain good
O2/N2 selectivity from 12 to 14.4, which can be mainly attributed to the carbon
matrix, because ZSM-5 is a non-selective zeolite with a pore size greater than

oxygen and nitrogen.

 

Figure 1.13 High-resolution TEM images of carbon/ZSM-5 composite membrane
with 9.1 wt% of ZSM-5 (denoted PIZZ): (a) ZSM-5 embedded in membranes, (b)

carbon phase in membranes.63

Table 1.3 Permeation properties of carbon/ZSM-5 membranes at 298K63

 

 

 

Flux/Barrersa Selectivity
Membrane 02 N2 O2/N2
PI (polymeric membrane) 0.16 0.04 4.2
Pi (carbon membrane) 2.21 0.16 13.8
PIZ1 (4.8 wt% ZSM-5) 70.0 5.85 12
PIZ2 (9.1 wt% ZSM-5) 431.0 30.0 14.4
PIZ3 (16.7 wt% ZSM-5) 499.0 39.6 12.6

a 1 Barrer = 1><10'1o cm3 (STP) cm/cm2 s cmHg

 

25

1.6 Research Goals and Significance

In contrast to the exponential reduction in permeability expected on the
basis of the Nielsen’s model for composites embedded with oriented lamellar
platelets, a more typical 2- to 4-fold reduction in gas permeability is usually
achieved by polymer-clay nanocomposites owing to the partial exfoliation and
disalignments of clay nanolayers in the polymer. This thesis demonstrates a new
strategy that substantially enhances the barrier properties of both thennoset
epoxy polymer and thermoplastic polyolefins by sandwiching novel self-
supported clay fabric films between polymer sheets. The paper-like clay films are
prepared by casting suspensions of the clay onto glass plates, allowing the
suspensions to evaporate and then peeling away the film. Van'ous homoionic
‘organic and homostructured and heterostructured mixed organic-inorganic ion
montmorillonite (MMT) and fluorohectorite (FH) clay fabric films were prepared
by ion-exchange reaction of naturally occurring Na*-MMT and synthetic Li"-FH
with a variety of organic modifiers in order to afford adequate interfacial adhesion
with the polymer matrices. Thermoplastic polyolefin-clay fabric film composites
were fabricated by sandwiching the resulting clay films between two sheets of
high density polyethylene (HDPE) films through a hot-pressing process. The
sandwiched composites exhibit more than 30-fold decrease in 02 transmission
rate with respect to the pure HDPE film. In the case of therrnoset system,
impregnating the self-supported clay papers with epoxy pre-polymers
successfully leads to composite films with more than 2-3 orders of magnitude

reduction in 02 permeability in comparison to the pristine epoxy matrix.

26

Another objective of this thesis is to investigate the synthesis of a new
phase mixture of saponite and cancrinite. Saponite (denoted SAP) belongs to 2:1
structured smectite clay, whereas cancrinite (denoted CAN) is a porous zeolite
with the natural composition Na6Ca[AISiO4]6003(H2O)2. The cancrinite
framework topology is shown in Figure 1.14. Synthetic Mg-saponite has potential
applications as reinforcement agents for themoset epoxy polymers29 and as acid
catalysts for alkylation65. However, the available synthetic routes are cost
intensive.65 In an effort to develop a more efficient and cost-effective synthesis of
VMg-saponite, sodium hydroxide instead of urea is used as the base source. Pure
delaminated SAP platelets were synthesized at low alkalinity at 90°C.
Interestingly, upon increasing the alkalinity, cancrinite was co-crystallized
together with saponite. Also, pure-phase cancrinte with dimensions ranging from
several hundred nm to several pm in length was achieved by varying synthetic

conditions.

Sodalite

 

Figure 1.14 Comparison between the framework topologies of cancrinite and
sodalite. The vertices represent the positions of Si and Al in the framework and

the edges represent the positions of bridging oxygen atoms.

27

It has been found that synthetic Mg-saponite can improve the tensile
properties of both rubbery and glassy epoxy polymers without the need for the
organic modification of the clay surface.29 Microporous cancrinite may also be
promising candidates for reinforcing epoxy polymers, considering that cancrinite
can be co-crystallized along with Mg-saponite during the saponite synthesis.
Therefore, the reaction conditions (e.g. alkalinity, reaction temperature, and
duration) are explored in the present work with the objective to form small
crystals of cancrinite (~ 100 nm) with large external surface area and high pore
volume. These particle properties will facilitate the uniform dispersion of
cancrinite particles in epoxy matrices. Moreover, as part of the exploration of the
synthesis of cancrinite with a small particle size, we find that co-crystallization of
sodalite, another mineral silicate with zeolite structure (c.f. Figure 1.14), takes
place in some cases.

Only a few studies have been reported on zeolite-epoxy composites,66459
wherein the focus was on the effect of natural zeolite particles on the curing
kinetics and mechanical properties of the epoxy system. With increasing zeolite
content, tensile strength decreased dramatically, while Young’s modulus
increased, which is the result typical of conventional composites.66

The present work aims to use synthetic cancrinite and sodalite as
reinforcing agents for thermoset epoxy polymers owing to their small particle size,
large surface area and high pore volume. The results show that the tensile
strength, modulus and toughness of the composites are simultaneously

enhanced without the need for an organic surface modification of zeolites. This

28

finding represents the first example for which the mechanical properties of glassy
epoxy composite are all improved through the use of a microporous zeolite as a
functional inorganic filler. Also, the properties of epoxy nanocomposites
reinforced by microporous cancrinite and sodalite are compared with those of
nanocomposites made from organo-clay nanoparticles, carbon nanotubes and

mesoporous silicas.

29

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32

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13,459.

33

CHAPTER 2

Thermoplastic Polyolefin-Clay Fabric Film Composites with High Barrier

Property towards Oxygen and Petroleum Fuel Components

2.1 Introduction

Since the late 19803 much research has focused on developing methods
to incorporate fully exfoliated smectite clays into polymers to enhance
mechanical and barrier properties.1 However, it has been a scientific challenge to
achieve the exfoliation of the clay nanolayers and even intercalated
morphologies in polymer matrices have been difficult to obtain, particularly for
thermoplastic polyolefins. The highly hydrophobic character of polyolefins
prevents the development of adequate physical binding to the aluminosilicate
surfaces of smectite clays. Two methods are often used to improve the interfacial
interactions between two phases: (i) the surface chemistry of clays is fine-tuned
by ion-exchange reaction with cationic organic surfactants such as
alkylammonium ions, and (ii) polar maleic anhydride (MA)—grafted polyolefins are
introduced in the formulations as compatibilizers?‘1O MA-grafted polyolefins are
typically characterized by very complicated molecular structures due to many

side reactions,”12

including severe chain degradation during the free radical
grafting process, as well as many impurities. An alternative approach is to use
reactive chain transfer agents to prepare end-functionalized polyolefins

containing a terminal functional group (OH, NH2, COOH, anhydride, etc.), while

34

maintaining a well-controlled polymer molecular weight and narrow molecular
weight and composition distributions.7'13

Clay platelet particles enhance the barrier properties of polymers
according to a tortuous path model developed by Nielsen.14 In the model the
impermeable clay nanolayers retard the transport of diffusing species such as
oxygen or water vapor. The magnitude of barrier improvement strongly depends
on the aspect ratio, volume fraction, and the state of dispersion and orientation of
the platy inorganic particles. In contrast to the expected orders of magnitude
reduction in permeability, a 2-4-fold reduction in permeation rates is more typical
owing to the limited exfoliation and dis-alignment of the individual clay layers.
The oxygen transmission rate of the epoxy matrix was reduced 1.25-4 times at 1-
5 vol% of organoclay loading."5 Long alkylammonium ions enhance the polymer
intercalation but increase the oxygen permeability coefficient, most likely due to
the phase separation at the interface between the polymer and the layered
silicates. Composites made from ethoxylated short chain alkylammonium-
montmorillonite exhibit the best oxygen barrier properties. It has been suggested
that the functional -OH groups in the ethoxylated alkylammonium ions may react
with the epoxy resin, energetically favoring the partial exfoliation of clay platelets
with the entropically unfavorable small d-sapcings. A recent study16 on the barrier
properties of olefinic nanocomposites showed a 300-fold improvement in barrier
properties was achieved due to the strong interfacial interaction between the wax
and the organoclay surface. It was also found out that phase separation

happened with the large difference between the freezing points of the polymer

35

and organoclay, leading to the decrease in the magnitude of barrier
improvements.

This research is focused on applying clay fabric film nanotechnology to
thermoplastic polymer systems for automotive applications, particularly for the
manufacturing of automotive fuel tanks for both conventional and flex-fuel

automotive vehicles. Our recent study":18

on the oxygen barrier properties of
epoxy-clay fabric film composites proposed a new strategy for achieving high
barrier perfon'nace by impregnating the self-supported clay papers with epoxy
pre-polymer. We successfully achieved 2-3 orders of magnitude reduction in 02
permeability in comparison to the pristine polymer. The present investigation of
oxygen barrier properties of high density polyethylene (HDPE)-clay fabric film
composites exhibits more than 30-fold decrease in 02 transmission rate with
respect to the pure HDPE film. The magnitude of barrier improvements of
polyolefin system is lower than that of the epoxy system, which can be explained
by the fact that polyethylene chains can only block interparticle voids on the
external surface of the clay films, unlike the epoxy system wherein the epoxy
monomer and curing agent may diffuse into the interlayers of the clay film, as
well as block the voids formed between clay particle edges upon curing.
Measurements carried out in collaboration with Central Laboratory of Ford Motor
Company have shown that the transmission rates of the HDPE-clay fabric film

composites can be reduced to about one-sixth with regard to three petroleum fuel

components: toluene, isooctane and ethanol.

36

2.2 Experimental
2.2.1 Materials

PGW Na+-montmorillonite (Na+-MMT) with a cation exchange capacity
(CEO) of 100 mqu 100 g was provided by Nanocor Inc.. A synthetic Li*-
fluorohectorite clay (Li"-FH) with a CEC similar to that of PGW montmorillonite
was purchased from Corning Inc.. All other chemicals were purchased from
Aldrich Chemical Co. and used without further purification.
2.2.2 Preparation of Homoionic and Mixed-ion Clay Fabric Films

Homoionic inorganic Na*-montmorillonite and Li+-fiuorohectorite clay films
were made by preparing a 5 wt% homogeneous aqueous colloidal suspension of
clay powders, then casting the suspension on a flat glass substrate, followed by
drying in air for 1 day at room temperature. Continuous and self-supported clay

fabric films were carefully peeled off the glass plate.

Homoionic organo-clay films and mixed inorganic-organic clay films were
prepared by ion exchange of aqueous inorganic clay suspension with the desired
amounts of onium ions. The ion-exchange reaction was carried out for 24-48 h at
ambient temperature. The exchanged clay were isolated by centrifugation and
washed with deionized water 3-4 times until free of Na”, then followed by casting
and drying in air at room temperature.

2.2.3 Hot-pressing the Organoclay Film with HDPE Sheets

An organoclay film approximately 8 cm x 8 cm was placed between two

sheets of HDPE film (6 inch x 6 inch) and the assembly placed at the center of

two flat Taflon-coated steel sheets used to prevent sticking. A sandwich-type

37

composite film was subsequently made in a Tetrahedron Press lnstmment using
a hot-pressing procedure, wherein the sheets were heated to 270-330°F at a rate
of 30°F/min under an applied pressure of 4 sz, followed by cooling down to the
room temperature at the same rate. The dwell time at the maximum temperature
and pressure was normally 6 min.
2.2.4 Characterization and Testing of Clay Films and HDPE - Clay Film
Composites

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 2003
diffractometer equipped with Cu Ka X-ray radiation and a curved crystal graphite
monochromator, operating at 45 KV and 100 mA. Samples of the clay films were
prepared by mounting the films on glass X-ray sample holders.

Therrnogravimetric analyses (T GA) of the clay films were performed using
a Cahn TG System 121 Analyzer. The clay films were cut into small pieces, which
were heated to 800°C at a rate of 5°C/min under air flow. The compositions of
clay films were determined by CHN chemical analysis and TGA.

Oxygen permeability measurements were performed on a Mocon Ox—tran
2/60 oxygen permeability instrument with 100% oxygen as test gas. The
specimens tested were the pristine HDPE polymer with an average thickness of
~0.06 mm (~2.36 mil), and the HDPE - clay fabric film composites with an

average thickness of ~0.14 to ~0.16 mm.

The fuel permeation tests of pristine polymeric and composite materials
were carried out by the Central Laboratory of Ford Motor Company. The

permeation measurements refer to three different test fuels: ethanol, isooctane

38

and toluene at 40°C.

2.3 Results and Discussion

2.3.1 Composition and Structural Properties of Homoionic and Mixed-ion
Clay Fabric Films

The formation of self-supported smectite clay films depends on the ability
of clay particles to disperse homogeneously with a high degree of exfoliation in
aqueous suspensions. Figure 2.1 shows X-ray patterns of Na*-montmorillonite
and Li”-fluorohectorite clay films and pristine Na*-montmorillonite powders,
respectively. As expected, the interlayer spacings (001 peak, 12.3 A) are the
same for the powder and oriented film form of the clay. However, it’s clear that
001 peak of the montmorillonite film sample exhibits more than ten times higher
intensity than that of the powder, indicating that the clay nanolayers are highly
oriented along the direction perpendicular to the surface of the film. This
orientation of the layers is realized by the restacking of exfoliated individual clay
platelets in the colloidal suspensions when casted on the glass plate. The 001
peak of the Li+-fluorohectorite film is 8 times more intense than that of Na*-
montmorillonite film as shown in Figure 2.1. The higher degree of platelet
orientation may be caused by the large aspect ratio of the fluorohectorite
platelets, which is over ten times that of montmorillonite particles and facilitates
more efficient nanolayer alignments. Homoionic inorganic Na"-montmorillonite
and Li”-fluorohectorite clay fabric films are self-supported, smooth, and semi-

transparent. Moreover, the Li"-flurohectorite fabric film is very robust paper-like

39

material and much more flexible than the Na"-montmorillonite film.

_._-.__.___-_._..._.__._ ___ __ ,__O-____!

 

 

Intensity (a.u.)

 

p—d
N
DJ
>0

 

 

 

 

 

 

 

 

 

$ 2...” *2 u L r “1...“ z. (a)
1 3 5 7 9 ll 13 15
26 (degree)

Figure 2.1 X-ray diffraction (XRD) patterns of (a) Na*-montmorillonite clay
powders, (b) Na”-montmorillonite clay fabric film and (c) Li+-fluorohectorite clay

fabric film.

Na*-MMT and Li*-FH clays are hydrophilic due to the inorganic cations
situated on the basal surface of the layered silicates. However, the polyethylene
POIYmer is exclusively hydrophobic. In order to match the polarity of the clay with
the polymer matrix, ion-exchange of inorganic clays with cationic surfactants is

necessary. Homoionic organic and heterostructured mixed inorganic-organic ion

40

clay fabric films using different onium ions were prepared to afford adequate
interfacial adhesion with the thermoplastic polyolefin. The X-ray basal spacings
for the clay films containing betaine, choline, and cetyltrimethylammonium (CTA+)
ions as the organic modifiers are provided in Table 2.1. Representative XRD

patterns of some organo-MMT fabric films are shown in Figure 2.2.

Table 2.1 Basal spacings of homoionic and mixed—ion montmorillonite and

fluorohectorite clay fabric films

 

clay sample organic onium ion basal spacings

 

designation modifier (%) a dam (A) des°"pt'°"

Na”-MMT (PGW) - - 12.3 continuous film
Na+lBET1oo-MMT betaine 100 14.3 continuous film
Na‘lBET15o-MMT betaine 150 14.3, 18.7 continuous film
Na‘lBET2oo-MMT betaine 200 18.7 continuous film
choline1oo-MMT choline 100 14.1 continuous film
Na‘lCTA“42-MMT CTA 42 27.8, 15.1 continuous film
Li“-FH - - 12.1 continuous film
choline1oo-FH choline 100 13.8 continuous film
Li‘lCTAHg-FH CTA 40 27.1, 13.5, 12.2 continuous film

 

a The percent onium ion-exchange represents the extent to which the ion-exchange sites of the
clay have been replaced by onium ions; the reported values were determined by elemental or
TGA analysis.

Betaine (CH3)3N+CH2COO' is a zwitterionic surfactant and doesn’t replace
inorganic cations because of its electrically neutral structure. Instead, betaine co-
exists with Na+ or Li+ ions in the clay layers. lntercalated betaine arranged a
monolayer and bilayer structure in the gallery at 100% and 200% equivalent
amounts, respectively. Phase segregation happened at a loading of 150%
eQUivalents, wherein both monolayer and bilayer orientations of betaine

molecules were observed in terms of two basal spacings. The long-chain

41

 

 

 

 

 

”14.111

15.111

 

 

 

 

 

Intensity (a.u.)

 

-' “ -- (d)

(b)

___- ___ ___ _ l 1 _ _,_L________,__ 2“,- _L_2_‘ 1 j (a)

12 14 16

 

 

 

 

 

8 10
20 (degree)

Figure 2.2 X-ray diffraction (XRD) patterns of (a) Na"/BET1oo-MMT, (b)
Na‘lBET150-MMT, (c) Na*/CTA*42-MMT and (d) choline1gg-MMT clay fabric films.

cetyltrimethylammonium (CTA‘) ion results in a large basal spacing of 27.8 A for
the mixed-ion clay film containing 42 mol% CTA” and 58 mol% Na+ ions (denoted
as Na‘lCTA+42-MMT). This spacing can be attributed to the clay gallery
interlayered by paraffin-like structured onium ions. However, for the additional
peak corresponding to the lower basal spacing of 15.1 A, the onium ions adopt a
monolayer orientation. Therefore, segregation of onium ions occurs in the clay

film due to the different orientation of alkyl surfactants in the clay interlayers.

42

Homoionic-choline exchanged montmorillonite and fluorohectorite films
have only one d-spacing of about 14 A. The segregation of the organic and
inorganic exchange cations into separate galleries is observed for the 40% CTA+
exchanged Li"-fluorohectorite film. Basal spacings of 27.1 A and 13.5 A suggest
the galleries are interlayered by paraffin-like and monolayer stmctured
alkylammonium ions, respectively. The diffraction peak at 12.2 A can be
attributed to the inorganic phase exclusively intercalated by Li+ ions. The
segregation of the ions is driven by the differences in the solvation properties of
the two very different types of cations.

Compared with the pristine Na+-MMT clay film, betaine modified Na*-MMT
is much more flexible, which is not observed with other partially or fully
exchanged organo-MMT films. Betaine penetrates into the clay interlayers and
adopts a monolayer or bilayer orientation in clay films. At the same time, betaine
molecules might act as a binder as well and occupy the spaces formed between
the platy clay particle edges. The edges of clay platelets are positively charged
and can undergo electrostatic interaction with COO‘ groups in betaine. Therefore,
the attractive flexible property of BET/Na*-MMT film can be attributed to a unique
nanostructure, in which densely packed clay layers are reinforced by the soft
betaine molecules between the clay platelets.

Figure 2.3 showed the thermogravimetric analysis (T GA) curves of Na*-
MMT and Na‘lCTA"42-MMT clay films. In both cases, the weight loss in the range
of 25 to 130°C is due to the evaporation of water adsorbed on the clay surfaces

or in the clay interlayers. Dehydroxylation of the inner 2:1 lattice hydroxyl groups

43

takes place around 600°C. The thermal decomposition of (CH3)3N*(CH2)150H3
starts around 220°C. Therefore, the HDPE melt-mixing temperature won’t cause
the Hofmann degradation of alkylammonium ions. The initial thermal degradation
temperature of CTAB powder is 170°C. The increased thermal stability of onium
ions in the clay can be explained by the strong electrostatic interaction with

negatively charged silicate layers.

 

 

 

 

 

 

 

 

 

110 ———--‘-——-- me,- ————e—... . - -AM _ . W71
90 ~ wmm film
M
A 70 “ Na+/CTA+42-NIMTfiIm
e\":
E
'3'? 50 ~ ,
3 j l
30 i I
I
I
10 H ]
CTAB
-10 t - -~ . . , ---.__._ _I
0 200 400 600 800 1000

Terrperature (°C)

Figure 2.3 Therrnogravimetric analysis (T GA) curves for cetyltrimethylammonium
bromide powder, sodium montmorillonite and CTA+ partially exchanged MMT

clay fabric films.

2.3.2 lnterfacial Adhesion and Oxygen Permeability of HDPE-Clay Fabric

44

Film Composites

Homoionic inorganic Na*-MMT and Li+-FH clay films exhibit no interfacial
adhesion to the thermoplastic polyethylene since the polar Na” cations situated
on the silicate basal planes inhibit the physical adsorption of non-polar -
(CH2CH2),.- chains to the outer surface of the clay film. Improvement in the
interfacial adhesion is achieved by the fully or partially exchanged organo-MMT
films with alkylammonium ions.

In fact, the clay films have two different external surfaces that result from
the fabrication process. One side of the clay film is in contact with the glass
substrate and the other surface contacts the air interface. The “glass” side of the
clay film is smoother than the “air" side, suggesting that the evaporation of water
during the drying procedure causes some disorientation of the individual clay
nanolayers in the upper part of the clay film (air side) and the flat glass substrate
may help the efficient alignment of nanoclay platelets in the lower part of the clay
film as well (Figure 2.4 (a)). Concerning disorientation, platelike particles can
aggregate in a disoriented edge-to-face manner and not only in a face-to-face
manner.

It’s interesting that two sides of the montmorillonite clay films exhibit
different degrees of interfacial interaction with the HDPE polymer coatings.
Peeling test shows the smooth “glass” sides of MMT clay films always adhere
better to HDPE sheets than the rough “air" sides. The XRD of the HDPE sheet
removed from the “air" interface of the Na+lcholine25-MMT film exhibits the

characteristic 001 peak of organo-MMT (Figure 2.4 (b)). Which indicates the weak

45

 

 

 

 

— _
— _ —
— — —
. — — _ _ T
"glass" Side

0’) . - . ._. ......... WWI
l
A l
:i l
5‘. l
a :
e l
8 i
s l
l
l

l
I 14.0 A
l

 

 

4 6
20 (degree)

Figure 2.4 (a) Schematic representation of orientation of the clay nanolayers at
the “air” and “glass” side of the clay fabric film and peeling off the HDPE coating
from the “air” side; (b) XRD of HDPE sheet removed from the “air" side of the

Na‘lcholine25-MMT film.

46

interface is at clay-clay interface instead of clay-polymer interface. It is well
known that the individual MMT platy particles are held together by the
electrostatic interaction. Therefore, pulling apart clay layers should not be easy.
However, at the “air" side, the disorientation of the individual clay nanolayers
decreases the magnitude of ionic bonding between the layers since some clay
layers only interact in part with edges of adjacent clay platelets as shown in
Figure 2.4 (a). On the other hand, the clay platelets at the “glass” side of the clay
paper are well oriented face-to-face, which optimizes the electrostatic attraction
between the stacked platelets. Another factor resulting in the weak clay-clay
interface is when the HDPE coating was peeled off the “air” side of the clay film,
the clay layers would move more relatively in the horizontal direction and in this
direction it is much easier to overcome the ionic bonding to take them apart than
in a vertical direction.

In comparison to organo-MMT composites, the interface failure at clay-
clay interface was more obvious for the HDPE/organo-FH fabric film composites,
wherein clay was clearly seen to be sticking to both sides of the HDPE sheet
coating when peeled apart. It might be caused by the bent edge of fluorohectorite
nanolayer taking less percent of space in the whole plate due to its higher aspect
ratio.

From the SEM images, the clay films possess micro or sub-micro voids
formed between particle edges and bent platelets. These voids generate
extremely high 02 permeability (>6000 cc-mil/mZ-day). However, betaine modified

clay films are extremely impermeable towards oxygen gas (< 0.1 cc-mil/mz-day),

47

which may be explained by its unique nanostructure, wherein betaine molecules

bind to the positively charged edges of clay layers and occupy the voids formed

between particle edges.

Table 2.2 Oxygen permeability data for HDPE-clay fabric film composites

 

 

. thickness of thickness of HDPE- 0 en
cfg’rgabgfiigr clay fabric clay fabric film penggbility
9° film (mm) composite (mm) (cc-miI/(mZ-day»
pristine HDPE 0.060 1832.3
Na‘IBEng-MMT 0.049 s01
choline1oo-MMT 0.040 0.150 80.0
Na‘lcholine25-MMT3 0.025 0.084 76.3
Na‘lCTA*42-MMT 0.035 0.160 54.4
Na‘lCTA*42-MMT° 0.035 0.097 130.0
Li*-FH 0.048 0.160 29.2
cholinemo-FH 0.048 0.160 207.5
Li*/CTA"4o-FH 0.030 0.141 883.6

 

a The HDPE coating on the “air” side of the clay film has been peeled off, therefore, only one
layer of HDPE coating exists on the “glass” side of the clay film.

Hot-pressing the clay films between a pair of thermoplastic HDPE sheets
results in little or no intercalation of polyethylene into the clay gallery because the
basal spacing doesn’t change upon sandwiching clay film. Instead, the high
density polyethylene coating blocks the surface voids on the outer region of the
clay film, which dramatically improves the barrier property of the films. The O2
permeability values of some composite samples were summarized in Table 2.2.
For the HDPE composites prepared from organo-MMT clay films, the O2
permeability is decreased 100- and 30-fold compared to the parent clay film and

the pristine HDPE sheet (1832 cc-mil/mz-day). reSpectively. Even after the

48

removal of the HDPE coating from the weak “air” interface of organo-MMT film,
the resulting composite retains the barrier property similar to the original
composite. However, the organo-FH/HDPE composite films are more permeable
than the corresponding montmorillonite composites with regard to oxygen. This
may be attributable to the much larger aspect ratio of fluorohectorite platelets.
The more bendable FH nanolayers may form folded sheets, resulting in more
voids in the PH clay film and higher oxygen transmission rate. Interestingly, Li‘-
FH/HDPE composites shows better barrier property toward oxygen (29.2
cc-mil/mz-day) than MMT composites.
2.3.3 Permeability of HDPE-Clay Fabric Film Composites toward
Representative Petroleum Fuel Components

An effort was made to test the clay fabric films for automotive applications,
particularly for potential applications in automotive fuel tanks. Measurements
carried out in collaboration with Central Laboratory of Ford Motor Company have
shown that high density polyethylene composite made from 150 mol % betaine
modified Na+-MMT clay fabric film is more than 200 times lower in permeability
toward representative petroleum fuel components, toluene and isooctane and
more than 50 times lower toward ethanol in comparison to the pristine polymer,
which is consistent with the extremely high oxygen barrier performance of
betaine montmorillonite films. However, this nanocomposite film is not applicable
due to the lack of a chemically stable interface between the clay film and the
HDPE matrix. The weak interface facilitates the condensation of the fuel

components at the clay-polymer interface, causing the failure of the interface and

49

the formation of blisters that fill with liquid fuel. Results from Ford Motor
Company show that BET15o-MMT/HDPE composite maintained low level of
permeation towards organic solvents in spite of the formation of blisters.
However, the blisters compromise the structural integrity and strength of the
composite.

In order to avoid blister formation, organic alkylammonium ions were used
to modify the inorganic clay. The resulting HDPE-organoclay film composite was
expected to provide adequate interfacial adhesion. The permeation rate curves of
composite materials over the aging time from 1 day to 35 days are shown in
Figure 2.5. Choline and CTA+ exchanged MMT and FH clay film composites are
not as impermeable as betaine MMT composite to petroleum components.
However, no blisters are formed in these four samples, proving the adequate
interfacial interaction at the clay-polymer interface. Although organo-FH/HDPE
composites exhibit higher oxygen permeability than MMT composites as
discussed before, they show less permeation than montmorillonite counterparts
toward petroleum components. Particularly, homoionic choline fluorohectorite
composite is the most impermeable to all three fuel components and the
transmission rate is reduced to about one-sixth in comparison to the pure HDPE
polymer.

Interestingly, toluene and ethanol permeation curves even show declining
Permeation rates in the case of CTA+ exchanged montmorillonite and
fluo rohectorite composites, which is rather unusual. Typical permeation curves

are evaluated based on the initial slope and the rate at equilibrium (flat). Most of

50

Figure 2.5 Permeation curves of HDPE-organoclay composite films and pure

HDPE polymer towards toluene, isooctane and ethanol vapor at 40°C.

51

a. TOLUENE Permeation, 40 deg C.

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

2500 r i —- ~—
' + cholineIOO-MMT
2000 _j:—f".\___ __ gm _ +Na/CTA42-MMT 1
\‘\\ + choline 1 OO-FH
@1500 .2” ___#___f:.g_ f. + Li/CTA40-FH
E» { \\\ -~ Pristine HDPE
57 r“ ‘\
'g 1000 q_ ”WW: \ \\_
OD \\
500 av" —;—;1— #*:—— —*F\ —A
Thaw—fl: _ + _ i _ f _a
0 +‘ ' T i r l'“ ' “““ ‘r T r —" '—“
0.0 5.0 10.0 15.0 20.0 25.0 30.0 35.0 40.0
Days
b. ISOOCTANE Permeation, 40 deg C.

1000 ~~--——--—--—--- ~-~---— - ~'~- ~ - a}
900 rr~ A 4M —‘l
800 $2 _, ,jfé cholineIOO-MMT 1

A 700 JLfl W V _ __ ____ ___ +Na/CTA42-MMT 4
NE +cholineIOO-FH .
g» 600 l/fi . ‘ I +Li/CTA40-FH ’1
§ " ""‘ " +Pristine HDPE ;
'~ A I
'E ‘.
00
0 I i T i l l _. i
0.0 5.0 10.0 15.0 20.0 25.0 30.0 35.0 40.0
Days
6. ETHANOL Permeation, 40 deg C.
300 2.2--- -- ___.__________ LA, I
250 a ___./Arm~v_4\\_____m --,-
if, \\ i
{€200 1-—-—-—- __ *——- +cholinelOO-MMT A
_ -— T‘ +Na/CtA42-MMT
5,50 /__/\__2 ___. _FW +choline100-FH __,
§ +Li/C’I‘A40-FH
“g +Pristine HDPE
80

 

 

 

 

0.0 5.0 10.0 15.0 20.0 25.0 30.0 35.0 40.0
Days

52

the curves are flat, which would suggest an equilibrium condition. The unusual
decrease in permeability with time might be due to the redistribution of long-chain
alkyl ammonium ions CTA+ during the penetration of organic molecules through
the composite film, which may block the free volume between the bent clay
layers and particle edges. The X-ray diffraction pattern of Na‘ICTA*42-MMT film
after the toluene permeation test clearly exhibits a decrease in the intensity of
001 reflection peak characteristic of organo-clay intercalated by paraffin-like
structured CTA“ (see Figure 2.6), suggesting a swelling of the onium ion galleries
and a change in the orientation state of this long alkyl chain onium ion in the
gallery. The swelling of the galleries may contribute to a reduction in void volume

between platelets.

 

 

 

 

15.1 A

’3 l

3

E‘

3:» 27.8 A

E:

Na+/CTA+42-MMT film
__...-""~
After penetrated by toluene vapor
0 2 4 6 8 IO 12 14 16
20 (degree)

Figure 2.6 XRD patterns of Na‘lCTA*42-MMT clay fabric film before and after a

toluene permeation test.

53

2.4 Summary

Self-supported clay fabric film nanotechnology was applied to the
thermoplastic polyolefin system by sandwiching the clay film between two sheets
of high density polyethylene (HDPE) in order to achieve high barrier property
towards small molecules such as oxygen, as well as organic solvents included in
the petroleum fuel (e.g., isooctane, toluene, ethanol). In order to match the
compatibility of the polyolefin with layered silicate, homoionic organic and
homostructured and heterostmctured mixed inorganic-organic ion clay films were
prepared by full or partial ion-exchange reaction of smectite clay Na*-
montmorillonite (MMT) and Li"-fluorohecton'te (FH) with long or short chain
quaternary ammonium ions: cetyltrimethylammonium (CTA”)
(CH3)3N’(CH2)15CH3, and choline (CH3)3N+CH2CH2OH, and zwitterionic
compound betaine (CH3)3N"CH2COO'. The pristine clay films are very permeable
to oxygen due to the micro- to sub-micro size voids formed between clay
nanolayer edges. However, the hot-pressed HDPE-clay composite films exhibit
more than 100 and 30 times of decrease in 02 permeability with respect to the
parent clay films and the pristine polymer, respectively. The enhancement in
barrier properties is attributable to blocking the voids between clay platelet edges
on the external surface of clay film by polyethylene chains. Betaine modified
montmorillonite film itself is extremely impermeable to oxygen owing to its unique
nanostructure, wherein betaine molecules clung to the positively charged edges
of clay layers and occupy the voids formed between particle edges. Moreover,

high density polyethylene composite made from 150 mol % betaine modified

Nai-MMT clay fabric film is more than 200 times lower in permeability toward
representative petroleum fuel components, toluene and isooctane and more than
50 times lower toward ethanol, in comparison to the pristine polymer. However,
this nanocomposite film is not applicable due to the lack of a chemically stable
interface between the layered silicate and the polymer. On the other hand,
organic alkylammonium ion modified clay films afford adequate interfacial
adhesion. In the case of homoionic choline fluorohectorite composite, the
transmission rate of all three organic solvents is reduced to about one-sixth in

comparison to the pure polymer.

55

References
(1) Pinnavaia, T. J., Beall, G. W., Eds. Polymer-Clay Nanocomposites; John
Wiley & Sons: New York, 2000.

(2) Jacquelot, E.; Espuche, E.; Gerard, J.-F.; Duchet, J.; Mazabraud, P. J. Polym.
Sci, Part B: Polym. Phys. 2006, 44, 431.

(3) Zhong, Y.; Janes, D.; Zheng, Y.; Hetzer, M.; Kee, D. D. Polym. Eng. Sci. 2007,
47, 1 101.

(4) Arunvisut, S.; Phummanee, 8.; Somwangthanaroj, A. J. Appl. Polym. Sci.
2007, 106, 2210.

(5) Picard, E.; Gauthier, H.; Gerard, J.-F.; Espuche, E. J. Colloid Interface Sci.
2007, 307, 364.

(6) Deshmane, C.; Yuan, Q.; Misra, P. D. K. Mater. Sci. Eng. A 2007, 460-461,
277.

(7) Wang, Z. M.; Nakajima, H.; Manias, E.; Chung, T. C. Macromolecules 2003,
36,8919.

(8) Galgali, G.; Ramesh, C.; Lele, A. Macromolecules 2001, 34,852.

(9) Nam, P. H.; Maiti, P.; Okamoto, M.; Kotaka, T.; Hasegawa, N. Usuki, A.
Polymer 2001, 42, 9633.

(10) Hotta, 8.; Paul, D. R. Polymer, 2004, 45, 7639.

(11) Lu, 8.; Chung, T. c. Macromolecules 1999, 32, 2525.

(12) Priola, A.; Bongiovanni, R.; Gozzelino, G. Eur. Polym. J. 1994, 30, 1047.
(13) Chung, T. C. Prog. Polym. Sci. 2002, 27, 39.

(14) Nielsen, L. E. J. Macromol. Sci. (Chem) 1967, A1 , 929.

( 15) Suter, U. W.; Osman, M. A.; Mittal, V.; Morbidelli, M. Macromolecules, 2004,
37, 7250.

( 16) Chaiko, D. J.; Leyva, A. A. Chem. Mater. 2005, 17, 13.

(17) Triantafyllidis, K. S.; LeBaron, P. C.; Park, l.; Pinnavaia, T. J. Chem. Mater.
2006, 18,4393.

56

(18) Shao, H.; Pinnavaia, T. J. Unpublished data, 2006.

57

CHAPTER 3

Gas Barrier Properties of Transparent Epoxy-Clay Fabric Film Composites

3.1 Introduction

The use of inorganic nanoparticles as additives to enhance polymer
performances has been established. The technical advances provide numerous
commercial opportunities in areas ranging from advanced aerospace systems to
commodity plastics.1 Small amounts (1-5 wt%) of highly anisotropic nanoparticles,
such as layered silicates, carbon nanotubes or mesoporous silicas. provide
polymer property enhancements that are comparable to those achieved by
conventional loadings (15-40 wt%) of traditional fillers. Layered silicate clays, in
particular smectite clays, have been widely used to improve the mechanical,
thermal and barrier properties of thermoset and thermoplastic polymers. Owing
to their high aspect ratio, ion-exchange capacity, good crystallinity and chemical
robustness, smectite clays provide improvements that include high modulusz,
increased tensile strength“ and heat distortion temperatures, and decreased gas
permeability:H1 and thermal expansion coefficient.3'4

Polymers are used extensively in packaging applications, especially in
food and drug packaging, where transparency and gas barrier properties are
important criteria. Clay platelet particles (e.g., montmorillonite, fluorohectorite)
enhance the barrier property of polymers according to a tortuous path model

developed by Nielsen.12 In this model, the impermeable clay nanolayers retard

58

the transport of diffusing species such as oxygen or water vapor. The magnitude
of improvement in barrier properties strongly depends on the aspect ratio,
volume fraction, and the state of dispersion and orientation of the clay platelets.
Epoxy-clay nanocomposites have been extensively studied with regard to the
mechanism of clay tactoid intercalation and exfoliation, and the enhanced
mechanical properties.”22 However, few researches have been reported
regarding the gas barrier properties of epoxy nanocomposites.23”26

The oxygen transmission rate of the epoxy matrix was reduced to 1.25- to
4-fold at 1-5 vol% organoclay loading.23 Long alkylammonium ions on the
exchange sites of the clay enhance polymer intercalation, but contribute to an
increase in the oxygen permeability coefficient probably due to the phase
separation at the interface between the polymer and the layered silicates, which
leads to a decrease in polymer density at the interface, resulting in higher “free
volume” and consequently a higher transmission rate. In contrast to expectation,
the composites made from organo-montmorillonite with ethoxylated
alkylammonium ions exhibit the best oxygen barrier properties. The functional -
OH group of the N-CH2CH2OH moiety in the short alkyl chain onium ions may
react with the epoxy resin, energetically favoring the partial exfoliation of clay
particles with the entropically unfavorable small d-spacings in epoxy matrix. This
is verified by the existence of single nanolayers shown in the TEM cross-section
images. However, only intercalated tactoids are observed for nanocomposites
made from montmorillonite with non-functional long alkyl groups on the onium

exchange cations, leading to even ~2 times higher oxygen transmission rate than

59

that of the neat polymer. Therefore, exfoliated aluminosilicate layers build a
better barrier toward the diffusing gas molecules, whereas the polymer
interlayered tactoids do not contribute much to the permeation barrier
performance. However, the best barrier property enhancement is only a 4-fold
reduction in the oxygen permeability in comparison to the neat epoxy polymer
owing to the limited exfoliation and dis-alignments of the individual clay layers. A
recent study26 of the oxygen barrier properties of epoxy-clay fabric film
composites proposed a new strategy based on impregnating the self-supported
clay papers with epoxy polymer. These clay fabric films provide 2-3 orders of
magnitude reduction in O2 permeability. However, the adhesion of the epoxy
matrix to the clay film was poor, thus limiting the usefulness of the films.

In our study, in order to improve the interfacial adhesion of clay films to the
thermoset epoxy matrix, various homoionic organic and homostructured and
heterostructured mixed inorganic-organic ion clay films were prepared by full or
partial ion-exchange reactions of smectite clay Na*-montmorillonite (MMT) and
Li*-fluorohectorite (FH) with organic onium ions: cetyltrimethylammonium (CTA‘)
(CH3)3N*(CH2)1sCH3, choline (CH3)3N*CH2CH2OH, glycidyltrimethylammonium
(GTA‘) (CH3)3N*CH2C{-lB-l2 and diprotonated primary 0, 1.11—diamine
H2NCH(CH3)CH2[OCH2CH(CH3)]XNH2 (denoted Jeffamine 02000 with x = 33.1).
The objective of the present work is to prepare homoionic organic and mixed
inorganic-organic ion smectite clay fabric films and to examine the interfacial
adhesion between the epoxy coating and the clay films. In addition to the very

high oxygen barrier properties, some epoxy-clay fabric film composites also

60

provide attractive optical properties.

3.2 Experimental
3.2.1 Materials

The epoxide resin used for the formation of the glassy epoxy — clay
composite films was a diglycidyl ether of bisphenol A (DGEBA), more specifically,

Shell EPON 826 with an epoxy equivalent weight of ~182 g/equivalent:

O
CHQCH-CH2{O-©~((::Ei©—O-CHZE:-CszO—ETOEOCH2C/SLH2
Cl-l3 " CH3

n = 0 (90 %), 1 (10 %) (Scheme 3-1)
The curing agent was a polyoxypropylene diamine of the type
NH2CH(CH3)CH2[OCH2CH(CH3)]xNH2 (x=2.6), which was provided by Huntsman
Chemicals under the tradename Jeffamine 0230. Jeffamine 02000 (x=33.1) was
used for the preparation of heterostructured mixed organic-inorganic ion clay
films. PGW Nat-montmorillonite (Nai-MMT) with a cation exchange capacity
(CEC) of 100 mqu 100 g was provided by Nanocor Inc.. A synthetic Li‘-
fluorohectorite clay (Li*-FH) with a CEC similar to that of PGW montmorillonite
was obtained from Corning Inc.. All other chemicals were purchased from Aldrich
Chemical Co. and used without further purification.
3.2.2 Preparation of Homoionic and Mixed-ion Clay Fabric Films
Homoionic inorganic Na*-montmorillonite and Li"-fluorohectorite clay films
were made by preparing a 5 wt% homogeneous aqueous colloidal suspension of

clay powders, then casting the suspension on a flat glass substrate, followed by

61

drying in air for 1 day at room temperature. Continuous and self-supported clay
fabric films were carefully peeled off the glass plate.

The homoionic organic clay films and the mixed inorganic-organic clay
films were prepared by ion exchange of an aqueous inorganic clay suspension
with the desired amount of onium ions. Usually, the ion-exchange reactions were
carried out for 24-48 h and the exchanged clay were isolated by centrifugation
and washed with deionized water for 3 to 4 times until free of Na". The washed
clay was used to form a self-supported film as described above.

3.2.3 Preparation of Epoxy-Clay Fabric Film Composites

Pristine glassy epoxy polymer specimens (disks) were prepared by mixing
stoichiometric amounts of epoxy monomer (EPON 826) and curing agent
(Jeffamine 0230) at 50°C for ~30 min. Liquid mixture was then outgassed at
room temperature for 20 min and cured on an aluminum disk mold at 75°C for 3
h and then at 125°C for another 3 h under N2 flow. The formed epoxy disks were
polished to an average thickness of ~0.2 mm (~8 mil).

For the preparation of the epoxy-clay composite film sepecimens, the self-
supported clay fabric films were dip-coated with a stoichiometric mixture of
outgassed EPON 826 and Jeffamine 0230. The wetted films were then
suspended vertically to remove excess uncured resin. After being allowed to age
in air for 1 day, the partially cured composite was heated in an oven at 125°C for
3 h in nitrogen to complete the cure.

3.2.4 Characterization and Testing of Clay Films and Epoxy - Clay Film

Composites

62

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 200B
diffractometer equipped with Cu Ka X-ray radiation and a curved crystal graphite
monochromator, operating at 45 KV and 100 mA. Samples of the clay films or the
epoxy — clay composite films were prepared by mounting the films on glass X-ray
sample holders.

Therrnogravimetric analyses (T GA) of the clay films were performed using
a Cahn TG System 121 Analyzer. The clay films were ground to produce
powdered clays, which were heated to 800°C at a rate of 5°C/min under N2 flow.

Scanning electron microscope energy dispersive spectroscopy (SEM-EDS)
was performed on a JEOL 6400V microscope with LaBg emitter, and the EDS
detector was an INCA x-sight. The film samples were mounted on aluminum
stubs and coated with carbon before the analysis. The compositions of the clay
films were determined by SEM-EDS, CHN chemical analysis and TGA analysis.

Transmission electron microscopy (T EM) images were obtained on a
JEOL 2200FS field emission microscope with a ZrO/W Schottky electron gun and
an accelerating voltage of 200 kV. Thin sections (~100 nm thick) of the epoxy-
clay composite films were put on naked copper grids.

Oxygen permeability measurements were performed on a Mocon Ox-tran
2/60 oxygen permeability instrument with 100% oxygen as test gas. The
specimens tested were the pristine epoxy polymer disks with an average

thickness of ~0.2 mm (~8 mil), and the epoxy - clay fabric film composites with

an average thickness of ~0.05 to ~0.10 mm.

63

3.3 Results and Discussion
3.3.1 Composition and Structural Properties of Homoionic and Mixed-ion
Montmorillonite Clay Fabric Films and Their Epoxy Composites

The preparation of homoionic inorganic Nai-montmorillonite and Li“-
fluorohectorite clay fabric films was described in Chapter 2, along with their X-ray
patterns (Figure 2.1). Na*-montmorillonite and Li*-fluorohectorite clay fabric films
are self-supported, smooth, and semi-transparent. Moreover, Li+-fluorohectorite
fabric film is very robust paper-like material and much more flexible than Nai-
montmorillonite film. The difference in flexibility is in accord with the difference in
platelet aspect ratio which is ~2,000 for fluorohectorite vs. ~200 for
montmorillonite. The difference in film flexibility is maintained for the mixed
inorganic-organic cation exchange forms of the two clays.

Epoxy-clay fabric film composites were prepared by impregnating the clay
film with a stoichiometric mixture of the epoxy resin Epon 826 and curing agent
Jeffamine 0230. The impregnated films were partially cured at room temperature,
followed by full curing at 125°C for 3 hours. For the homoionic inorganic Na+-
montmorillonite and Li*-fluorohectorite clay films, both showed very poor
interfacial adhesion to the thermoset epoxy matrix. This can be explained by the
fact that large amounts of polar inorganic ions situated on the basal surfaces of
layered silicates impede the physical adsorption and intercalation of epoxy
polymer chain to the clay, thus resulting in very weak interfacial interaction

between these clay films and polymer.

64

Table 3.1 Basal spacings of homoionic and mixed-ion montmorillonite and

fluorohectorite clay fabric films

 

clay sample

organic

onium

basal spacings

 

designation modifier ion (%) a dom (A) description
Nai-MMT (PGW) - - 12.3 continuous film
cholinemo-MMT choline 100 14.1 continuous film
Na‘lCTA+42-MMT CTA 42 27.8, 15.1 continuous film
Na‘lGTA’kg-MMT GTA 60 14.1 continuous film
Na‘lGTA‘so-MMT GTA 80 14.1 continuous film
Na“/GTA*1oo-MMT GTA 100 14.0 powdered, no
film
Na"/02000*2g-MMT 02000 20 30.0, 15.9 continuous film
Na*/02000*4o-MMT 02000 40 42.9, 16.3 continuous film
Na+/02000*50-MMT 02000 60 41.3, 16.4 continuous film
Na"/02000*go-MMT 02000 80 - easily fractured
film
DZOOO‘tm-MMT 02000 100 42.4 powdered, no
film
Li*-FH - - 12.1 continuous film
choline1gg-FH choline 100 13.8 continuous film
Li‘lCTAHo-FH CTA 40 27.1, 13.5, continuous film
12.1
Li*/D2000*6,7-FH 02000 6.7 72.4, 12.2 continuous film
Li*/02000*15-FH 02000 15 57.0, 12.2 continuous film
Li*/D2000"2o-F H 02000 20 62.2, 12.2 continuous film
Cholinegg/02000"10- choline, 90, 10 71.0, 13.8 continuous film
FH 02000
ChOIIneao/D2000+20- choline, 80, 20 71.0, 13.8 continuous film

FH D2000

a The percent onium ion-exchange represents the extent to which the ion-exchange sites of the
clay have been replaced by onium ions; the reported values were determined by elemental or
TGA analysis.

Homoionic organic clay fabric films and homostructured and
heterostructured mixed inorganic-organic ion clay fabric films using different
onium ions were prepared to evaluate interfacial adhesion to a thermoset epoxy
matrix. The basal spacings for the films are summarized in Table 3.1. The

improvement in the interfacial adhesion is found for the composites made from

65

homoionic organo montmorillonite film intercalated exclusively by choline ion
(CH3)3N+CH2CH2OH (denoted cholinemo-MMT). The adhesion may be attributed
to the hydroxyl group in the choline cation. Due to the 3-membered ring strain in
the epoxide group of epoxy monomer, even the neutral —OH group in choline is a
good nucleophile to attack the epoxide ring, forming an ether linkage with the
epoxy monomer. The OH-epoxy interaction can be the driving force for the pre-
solvation of the clay by epoxy pre-polymer and curing agent. Thus, the
interaction of the choline and epoxy monomer is not only confined to the outer
region of the clay film but also exists in the internal region of the film. The
covalent bond between the choline cations and the epoxy resin contributes to the
improvement in adhesion.

The mixed ion clay film containing 42 mol% cetyltrimethylammonium
(CTA’) cations on the exchange sites of Na*-montmorillonite also exhibited good
adhesion to the epoxy polymer. The long-chain alkylammonium ion CTA" results
in a large basal spacing at 27.8 A in the parent clay film (as shown in Figure 3.1),
which facilitates swelling the clay gallery by the uncured precursor. XRD patterns
show that the parent clay film has two d-spacings. The peak at 27.8 A can be
attributed to the clay gallery interlayered by paraffin-like structured onium ions.
However, for the gallery with a lower basal spacing of 15.1 A, the onium ions
adopt a monolayer orientation. Therefore, phase segregation of clay layers exist
in the clay film due to the different orientation of alkyl surfactants in the clay
interlayer. For the epoxy composites, the disappearance of the diffraction peak

with the larger interlayer spacing indicates that these clay layers are highly

66

exfoliated in the thermoset epoxy matrix upon curing. The reduction in the
intensity of the diffraction peak with regard to the smaller basal spacing suggests
some of these clay particles are swelled by the polymer. Even though only
unfavorable physical van der Waals interaction exists between the epoxy polymer
and cationic surfactants CTA“, the dispersion of clay layers in the epoxy matrix
maximize the number of contacts of the two phases, therefore, resulting in some

degree of adhesion.

 

15.1 A

 

Intensity (a.u.)

 

 

1.5 3.5 5.5 7.5 9.5 11.5 13.5
20 (degree)

Figure 3.1 X-ray diffraction (XRD) patterns of (A) Na‘ICTA”42-MMT clay fabric

film and (B) epoxy-clay film composite made from (A).

67

Self-supported clay films exhibit two distinguishable surfaces due to the
method used to prepare the films. One surface contacts the glass substrate and
the other surface contacts the air interface. The “glass” side of the clay film is
smoother than its “air" side, suggesting that the evaporation of water causes
some disorientation of the individual clay nanolayers in the upper part of the clay
film (air side). Also, the flat glass substrate may help to align the clay platelets in
the lower part of the clay film. It’s interesting to note that two sides of the clay
films exhibit different degree of interfacial interaction to the epoxy polymer in both
epoxy-clay fabric film composites discussed above. Peeling test shows that the
smooth “glass” side of the clay film adheres better to the epoxy than the rough
“air” side. This interesting phenomenon will be further discussed later.

In order to further improve the adhesion between the clay film and epoxy
coating, functional onium ion exchanged clay films were prepared, wherein the
organic modifiers play a functional role in linking the clay film to the polymer
matrix by forming covalent bonds at the interface of the two phases. Alkyl
ammonium ions terminated with either epoxide group or amino group are good
candidates as functional organic modifiers.

A series of glycidyltrimethylammonium (GTA‘) -PGW clays were prepared
at three different levels of GTA“ exchange, namely 60%, 80% and 100%.
Continuous films were obtained at 60% and 80% exchange, but homoionic GTA+-
MMT clay at 100% exchange did not afford a continuous film. The basal spacing
of 14.1 A tells us that this short chain onium ion orients in a monolayer fashion in

the gallery. The decrease in the intensity of the 001 diffraction peak of the

68

epoxy/GTA‘ao-PGW composite film (as shown in Figure 3.2) suggests that some
of the clay galleries were swollen by the epoxy polymer. The small gallery height
in the parent homostructured mixed-ion clay film might impede the transport of
curing agent and epoxy monomer into the clay galleries. However, the
glycidyltrimethylammonium ions can form covalent bonds with curing agent 0230
by ring-opening polymerization. This is expected to be the driving force for the

solvation of clay layers by the polymer precursors.

 

 

 

 

Intensity (a.u.)

 

(A) ’
, (B)

 

 

 

 

1 3 5 7 9 11 13 15
20 (degree)

Figure 3.2 X-ray diffraction (XRD) patterns of (A) homostructured mixed-ion
Na’lGTA*go-MMT clay fabric film and (B) epoxy-clay film composite made from
(A).

69

 

 

 

‘424 A
l
... 41.3 A
=2 16.3 A
s
g 21.2 A
‘8
g ii 20.6 A
j “ 132000*.oo-Pow
42.9 . - - - - - -
15.9 A

 

 

r
—

Na+/DZOOO+6o-PGW .

  
 

Na+/D2000+4o-PGW

 

    

 

 

.1

1.5 3.5 5.5 7.5 9.5 11.5 13.5
20 (degree)

 

Figure 3.3 X-ray diffraction (XRD) patterns of diprotonated diamine Jeffamine

02000 exchanged montmorillonite fabric films at different onium ion equivalents.

Diprotonated diamine 02000 cation in mixed ion montmorillonite functions
concomitantly as a clay surface modifier, an acid catalyst for the intragallery
polymerization of epoxy monomers, and the curing agent in the thermoset epoxy
matrix. Figure 3.3 provides the XRD patterns for a series of diprotonated
Jeffamine 02000 onium ion exchanged PGW films. Replacement of up to 60% of

the Na“ ions by 02000 cations are still capable of forming continuous

70

montmorillonite clay films (see Table 3.1). However, homoionic organoclay
020001oo-MMT forms a powder rather than a continuous film upon casting the
clay suspension on glass. For the mixed-ion DZOOO“-Na+ montmorillonite films at
60%, 40% and 20% onium ion equivalent levels, two basal spacings are
observed, indicating the existence of two phases in the clay film. The larger d-
spacing of ~40 A can be attributed to clay galleries intercalated by polyamine
adopting a folded chain conformation as shown in Figure 3.4, part B. However,
some diamines adopt a more extended chain orientation, which result in a lower
d-spacing at ~16 A. The DZOOan-PGW clay film shows a second order harmonic
dooz at 20.6 A. There is no evidence of a separate inorganic-rich phase (d-

spacing 12.3 A) in any of the mixed-ion clay films.

 

(A) - - - -

 

 

Na” Na“ NH3+ NH3+
(B) - - - -
"'3+ NH2," Na” H3+
Hf Na“ ”'43+ NH3+

 

Figure 3.4 Schematic representation of the orientations of the diprotonated
diamine Jeffamine 02000 onium ions intercalated in the mixed Na+-onium ion
smectite clay: (A) the diamine adopts an extended chain conformation in the
gallery with low d-spacing of ~ 16 A; (B) the diamine adopts a folded chain

conformation in the gallery with high d-spacing of ~ 40 A.

71

 

41.3A

 

 

 

 

 

 

 

 

16.4 A
f; 20.6 A
3
.E‘
.2
(A) ___H
(B)
L— "l" - --—---—— J l i . (C) ___":
1.4 3.4 5.4 7.4 9.4 11.4 13.4
29 (degree)
l
15.9 A l
"‘ l
=3 15.2 A 5
3 :
.é‘
E»
.5
30.0 A
26.5 A

(D)

I?) w

 

 

 

_. _______.l_____ _._____s_2_ __ L

1.5 3.5 5.5 7.5 9.5 11.5 13.5
29 (degree)

Figure 3.5 X-ray diffraction (XRD) patterns of (A) heterostructured Na+/02000*60-
MMT clay fabric film, (B) epoxy-clay film composite made from (A), only aged
under ambient conditions, (C) epoxy-clay film composite made from (A), further
cured at 125°C for 3h, (D) heterostructured Na*/DZOOO+2o-MMT clay fabric film,
(E) epoxy-clay film composite made from (D).

72

 

Figure 3.6 TEM images of thin section of an epoxy-0200020 PGW composite film
showing the c-axis direction (cross-section): (a) low-magnification image showing
spacial distribution of clay aggregates; (b) exfoliated clay nanolayers in

composite film and (c) stacked clay layers in clay tactoids in composite film.

73

Figure 3.5 provides representative XRD patterns for Epoxy-0200060 and
Epoxy-0200020 PGW film composites. After being cured in air, the clay particles
with a basal spacing of 41.3 A are highly separated by the intragallery
polymerization of epoxy precursors, resulting in a decrease in the intensity of the
diffraction peak characteristic of clay nanolayers intercalated by folded-chain
diamine 02000. Upon full curing at higher temperature, these clay layers
remained highly exfoliated, as verified by the disappearance of the diffraction
peak in the small-angel region of the XRD pattern. Evidence for clay layer
swelling by polymer in the DZOOOzo-PGW film is provided by the reduction in the
intensity of the 29 peak at ~30 A. However, clay platelets with lower spacings of
16.4 A and 15.9 A do not undergo an expansion in basal spacing upon polymer
composite formation. Thus, these clay particles are not interlayered by the
polymer. Both exfoliated and intercalated clay structures can be observed from
TEM images of thin-sectioned epoxy/ D2OOOZO-PGW composite film shown in
Figure 3.6, in agreement with the XRD pattern.

The resulting epoxy-02000 PGW composite films have attractive optical
properties, as illustrated by the optical images shown in Figure 3.7, parts a and b.
The composite film made from DZOOOzo-PGW is totally transparent and very
smooth. The “glass” side of the film shows good adhesion to the epoxy matrix.
The adhesion at the “air" side of the film is less robust. Epoxy composites from
40% and 60% 02000 onium ion exchanged PGW have unfavorable “wetting
marks” on the surface, which decrease the clarity of the composites. Moreover,

the 40% and 60% 02000 composites are very brittle, even more brittle than the

74

parent clay paper, which is consistent with the high modulus of the epoxy matrix.
The fragile nature of the composites formed from 0200040 and DZOOOGO-PGW
clay film is associated with the “wetting mark” imperfections but the interfacial
adhesion exhibits the best result so far and the epoxy polymer coating cannot be

peeled off either side of the clay film.

a - b
A Epoxy-D. )Oi’ho PGW Film Con oosites

A - ‘fi‘ié‘fir 5:472“! Frill”: (.lz'iém OSlleS ,
. Epoxy E , ' ’ p Thick'res“; 78 pm

91...} L-
~'<;:' k_,f‘._tx:§f~f, (3.; “tn

Figure 3.7 Optical images of (a) Epoxy-0200020 PGW film composites with a
thickness of 65 pm and (b) Epoxy-0200040 PGW film composites with a

thickness of 78 pm.

The parent mixed-ion 02000-Na+ montmorillonite clay films are opaque,
which can be attributed to the high degree of light scattering caused by the
efficient alignment of clay platelets. However, the epoxy composite films are
optically transparent. The improvement in transparency might originate from the
reduction of the oriented stacking of clay nanolayers due to the exfoliation of clay
layers by the polymer. Optical clarity of the nanocomposites depends on layered

silicate orientation, as the amount of scattering and haziness is relative to size of

75

the particle and wavelength of light. Alignment of the layers perpendicular to the
surface will decrease the effective particle size, by a factor associated with the
plate aspect ratio. However, a more likely reason for the improved transparency

for the composite films is the matching of the refractive index for the clay and the

epoxy polymer.

 

 

 

 

 

Intensity (a.u.)

 

 

 

 

4 6
29 (degree)
Figure 3.8 X-ray diffraction (XRD) pattern of epoxy coating released from the

“air" side of Na+/D2000+20-MMT clay fabric film.

Adhesion of the clay film to the epoxy polymer is facilitated by the
formation of covalent bonds at the interface of the two phases. The large contact
areas between the two phases due to clay layer swelling by the epoxy matrix

also promote better adhesion. With good polymer-clay adhesion, the weakness

76

at the interface may be caused by clay-clay interface, not by clay-epoxy interface.
This hypothesis is verified by the X-ray diffraction pattern (Figure 3.8) and SEM-
EDS analysis of the epoxy layer released from the “air” side of a DZOOOzo-PGW
film. The reflection peak at 15.8 A is characteristic of clay layers intercalated by
diamines with a folded chain conformation, thus verifying that the clay particles
undergo delamination when the film is peeled apart. A SEM-EDS analysis of the
epoxy removed from the “air” side of 0200020-PGW film showed the molar ratio
of Na : Mg : Al : Si is 1.67 : 1.27 : 5.30 : 11.45, in agreement within experimental
error with the unit cell composition of montmorillonite and further proving that the
weakness at the interface is caused by the clay-clay interface.

The individual MMT nanolayers are held together by the electrostatic
interaction. Therefore, pulling apart clay layers should not be easy. However, at
the “air" side, the disorientation of the individual clay nanolayers decreases the
magnitude of ionic bonding between the layers since some clay layers only
interact in part with the edges of other clay platelets as illustrated in Figure 2.4 (a)
of Chapter 2. On the other hand, the clay platelets at the “glass” side of the clay
paper are well oriented face-to-face, which optimizes the electrostatic attraction
between the stacked platelets. Another factor that weakens the clay-clay
interface is when the epoxy coating was peeled off the “air” side of the clay film,
the clay layers move relatively in a horizontal direction and in this direction it is
much easier to overcome the ionic bonding to take them apart in comparison to

taking them apart in a vertical direction.

77

3.3.2 Composition and Structural Properties of Homoionic and Mixed-ion
Fluorohectorite Clay Fabric Films and Their Epoxy Composites.

Homoionic inorganic Li+-fluorohectorite clay films are very robust paper-
like materials with high flexibility that can be associated with the large aspect
ratio of fiuorohectorite. The aspect ratio of fluorohectorite is ten times higher than
that of montmorillonite. Also, this large particle size results in the more efficient
alignment of clay nanolayers in fluorohectorite paper, as verified by a much more
intense 001 reflection peak compared with montmorillonite.

In order to obtain good interfacial interaction with an epoxy matrix, various
homoionic organic and mixed inorganic-organic ion fluorohectorite clay films
were prepared through ion-exchange reaction with desired onium ions (Table 3.1).
Homoionic choline exchanged fluorohectorite film only has one d-spacing of 13.8
A as shown in the XRD pattern of Figure 3.9. The segregation of the organic and
inorganic exchange cations into separate galleries exist in the 40% CTA+
exchanged Li*-fluorohectorite film. Basal spacings of 27.1 A and 13.5 A suggest
the onium ion galleries are interlayered by alkylammonium ions with paraffin-like
and monolayer orientations, respectively. The diffraction peak at 12.1 A can be
attributed to the inorganic phase exclusively intercalated by Li+ ions. The
segregation of the ions is driven by differences in the solvation properties of
these two very different types of cations. The inorganic galleries of the
heterostructure are hydrophilic and readily solvated by water, whereas the

organic galleries are lipophilic and avoid hydration.27

78

 
 

l
’ J 12.1 A
Li+/CTA+4o-FH

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

’3 ___ -
65‘.
.§‘
8
g L cholineloo-FH
12.2 A ‘
71.0 A . +
‘ --~.._ 011011116ng2000 2_o-FH
‘\ LLi+lD2000+2_o-FH Epoxy composite
62.2 A + +
k ILL/D2000 zo-FH
72.4 A
x L Li+/D2000+6.7-FH A
O 5 10 15
29 (degree)

Figure 3.9 X-ray diffraction (XRD) patterns of homoionic organic and mixed

inorganic-organic ion fluorohectorite films.

The partial replacement of Li+ ions with diprotonated diamine D2000 also
results in the segregation of two types of cations into separate galleries, affording
a mixture of an organic and inorganic clay phases. XRD patterns of Li‘lD2000*6_7-
FH and Li+/D2000+2mFH (Figure 3.9) show that the diffraction peak of the
inorganic phase at 12.2 A is more intense than that of organic phase at 72.4 A
and 62.2 A, resulting from the galleries interlayered by diprotonated D2000 with a

folded chain conformation. The disappearance of the peak for the organic phase

79

in Li‘lD2000‘zo-FH lepoxy composites after epoxy impregnation indicates that
those clay layers are highly swollen by the epoxy matrix. However, the inorganic
phase is not intercalated by polymer.

Li‘lDZOOO‘ey-FH fabric film can afford a highly transparent composite with
‘ epoxy polymer (Figure 3.10, part a). However, the “wetting marks” on the

composite films made from Li"/02000+20-FH decrease the optical transparency

(Figure 3.10, part b).

     

Figure 3.10 Optical images of (a) Epoxy-020005] FH film composites with a
thickness of 52 pm and (b) Epoxy-0200020 FH film composites with a thickness

of 53 pm.

3.3.3 Oxygen Barrier Properties of Epoxy-Clay Fabric Film Composites.

The bulk density of Na+-montmorillonite film (2.0 g/cm3) is smaller than the
intrinsic density of Na”-MMT tactoids (2.6 g/cm3),24 indicating that approximately
23 vol % of the clay fabric film is occupied by the voids formed between bent clay

layers or clay platelet edges. This imperfect tiling of the layers results in the high

80

oxygen permeability of pristine Na"-montmorillonite film (> 6000 cc-mil/mz-day).
SEM image of cholinemo-FH film (Figure 3.11) shows the micro-size voids formed

between platy fluorohectorite particles.

 

Figure 3.11 SEM micrograph of the choline1oo-FH fabric film.

Table 3.2 Oxygen permeability data for epoxy-clay fabric film composites

 

Thickness of Thickness of epoxy- Oxygen

 

Sample clay fabric film clay composite film permeability
(mm) (mm) (cc-mil/mz-day)
Neat epoxy 0.20 98.9
Cholinemo-PGW/Epoxy 0.050 0.074 110.7
D200060-PGW/Epoxy 0.071 0.103 440.8
D200020-PGW/Epoxy(1)a 0.038 0.049 6.70
D200020-PGW/Epoxy 0.038 0.061 < 0.1

 

a (1) represents the epoxy coating on the “air" side of the clay film has been peeled off, therefore.
only one layer of epoxy coating exists on the “glass” side of the clay film.

Oxygen transmission rates of some epoxy-clay composite films are
summarized in Table 3.2. In comparison with the neat epoxy. more than 3 orders
of magnitude reduction in Oz permeability was achieved by D20002o-PGW/epoxy
composite film, wherein the epoxy precursors diffuse efficiently into the clay

fabric film, energetically favored by the covalent bond formation between the

81

functional diprotonated D2000 onium ions on the clay exchange sites and the
epoxy resin. After being cured, the crossed-linked epoxy polymer fills voids
between clay platelet edges, which lead to the enhancement in the gas barrier
property. Even after epoxy coating at the “air” side of DZOOOzo-PGW is peeled off
the composite film, the remaining composite having one side of epoxy coating at
the “glass” side retains the barrier property and exhibits 1 order and 3 orders of
magnitude decrease in Oz permeability with regard to the neat epoxy and the
parent clay film, respectively.

In the preparation of the epoxy-clay composite film, the D2000 onium ions
present in the clay films are not counted as being supplementary to the amino
groups of curing agent Jeffamine D230. The fraction of epoxide groups capable
of reacting with the diprotonated D2000 onium ions on the clay exchange sites is
limited by stoichiometry to approximately ~3% in the DZOOOzo-PGW film.
Therefore, this small amount of D2000 can be disregarded as contributing to the
curing process. However, in the case of DZOOOw-MMT/epoxy composite film
wherein the “wetting mark” imperfection exists, the moles of onium ions present
in the clay can react with around 15% of epoxy monomers. This large excess of
amino groups may deteriorate the barrier property of the epoxy polymer. Thus,
D200060-MMT/epoxy composite film is even 4 times more permeable than the
neat polymer, whereas, oxygen permeability is reduced almost 15-fold compared
to the parent clay film, which is again attributed to the organic components
occupying the space between the clay particles. The resulting epoxy composite

from choline onium ion fully exchanged montmorillonite film does not enhance

82

the barrier property of the parent clay film as dramatically as 0200020-
PGW/epoxy. This can be rationalized by the unfavorable small interlayer spacing
of cholinemo-MMT (14.1 A) which impedes the penetration of epoxy pre-polymers
into the film interior. Therefore, epoxy polymer chains can only block the voids in

the outer region of the clay film.

3.4 Summary

We have developed epoxy-clay fabric film composites with high 02 gas
barrier properties by impregnating self-supported clay fabric films with epoxy pre-
polymers.26 In order to improve the interfacial adhesion between the oriented
clay particles and the epoxy matrix, homoionic organoclays and homostructured
and heterostructured mixed inorganic—organic ion clay films were prepared.
Naturally occurring Na*-montmorillonite (MMT) and synthetic Li*-fluorohecton'te
were exchanged with cetyltrimethylammonium (CTA+), choline
(CH3)3N*CH2CH20H, glycidyltrimethylammonium (GTA‘) and diprotonated
primary (1, w—diamine H2NCH(CH3)CH2[OCH2CH(CH3)]XNH2 (denoted Jeffamine
D2000 with x = 33.1). In comparison with the parent inorganic clay films, which
show no adhesion toward the epoxy matrix, CTA‘, choline and GTA+ clay films
present substantially improved interfacial adhesion. The large interlayer spacing
of ~27 A for CTA+ organoclay (42% onium ion exchange) and the functional —OH
and epoxide groups in choline and GTA‘, respectively, may facilitate the
solvation of clay nanolayers by epoxy precursors and promote adhesion. Clay

films partially exchanged by di-protonated Jeffamine D2000 provide the best

83

adhesion to epoxy polymer and impart the high optical transparency. The di-
protonated organocations play a dual role of organic modifier and curing agent
for the thermoset epoxy matrix. Thus, the clay fabric films links to the epoxy
matrix by forming covalent bond at the interface of the two phases. An
interesting finding based on XRD and SEM-EDS measurements is that polymer
adhesion improves with the increasing degree of clay platelet orientation at the
surface of the fabric film. Also, fabric films assembled from fluorohectorite plates
are very robust paper-like materials, whereas most of the organo-MMT films are
very fragile. The difference in durability is correlated with the aspect ratio of the
platelets. The pristine clay films are very permeable to oxygen due to the micro-
to sub-micro size voids formed between clay nanolayer edges. However, the
epoxy-clay fabric composite films show 3-4 orders of magnitude decrease in Oz
permeability with respect to the parent clay films. The enhancement in barrier
properties is attributed to the efficient diffusion of epoxy monomer and curing
agent into the clay film and to the blocking of the voids between clay platelet

edges upon curing.

84

References

(1 ) Vaia, R., AMPTIAC Newsletter 2002, 6, 17.
(2) Jana, S. C.; Pattanayak, A. Polymer 2005, 46, 5183.

(3) Zhu, Z.; Yang, Y.; Yin, J.; Wang, X.; Ke, Y.; Qi, Z. J. Appl. Polym. Sci. 1999, 73,
2063.

(4) Zhu, 2.; Yang, Y.; Yin, J.; Wang, X.; Ke, Y.; Qi, Z. Polymer 1999, 40, 4407.

(5) Yano, K.; Usuki, A.; Okada, A.; Kurauchi, T.; Kamigaito, O. J. Polym. Sci, Part
A: Polym. Chem. 1993, 31, 2493.

(6) Grunlan, J. C.; Grigorian, A.; Hamilton, C. B.; Mehrabi, A. R. J. Appl. Polym.
Sci. 2004, 93, 1102.

(7) Saura, J. J.; Gavara, R.; Cabedo, L.; Gimenez, E. Polymer 2004, 45, 5233.
(8) Pinnavaia, T. J.; Lan, T.; Kaviratna, P. D. Chem. Mater. 1994, 6, 573.

(9) Khayankam, O.; Magaraphan, R.; Schwank, J. W. J. Appl. Polym. Sci. 2003,
89, 2875.

(10) Yano, K.; Usuki, A.; Okada, A. J. Polym. Sci, Part A: Polym. Chem. 1997,35,
2289.

(11) Chang, J.; Park, K. M.; Cho, D. Polym. Eng. Sci. 2001, 41, 1514.
(12) Nielson, L.E. J. Macromol. Sci. (Chem) 1967, A1 , 929.

(13) Pinnavaia, T. J.; Wang, M. S. Chem. Mater. 1994, 6, 468.

(14) Pinnavaia, T. J.; Shi, H.; Lan, T. Chem. Mater. 1996, 8, 1584.

(15) Pinnavaia, T. J.; Triantafillidis, C. S.; LeBaron, P. C. Chem. Mater. 2002, 14,
4088.

(16) Pinnavaia, T. J.; Triantafillidis, C. S.; LeBaron, P. C. J. Solid State Chem.
2002, 167, 354.

(17) Pinnavaia, T. J.; Lan, T. Chem. Mater. 1994, 6,2216.

(18) Pinnavaia, T. J.; Lan, T.; Kaviratna, P. D. Chem. Mater. 1995, 7,2144.

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(19) Giannelis, E. P.; Messersmith, P. 8. Chem. Mater. 1994, 6, 1719.

(20) Pinnavaia, T. J.; Lan, T.; Kaviratna, P. D. J. Phys. Chem. Solids 1996, 57,
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(21) Pinnavaia, T. J.; Wang, 2. Chem. Mater: 1998, 10, 1820.
(22) Jana, S. C.; Park, J. H. Macromolecules 2003, 36, 2758.

(23) Suter, U. W.; Osman, M. A.; Mittal, V.; Morbidelli, M. Macromolecules, 2004,
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(24) Osman, M. A.; Mittal, V.; Lusti, H. R. Macromol. Rapid Commun. 2004, 25,
1145.

(25) Kim, J.; Hu, 0.; Woo, R. S. C.; Sham, M. Composites Science and Tech.
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(26) Pinnavaia, T. J.; Triantafyllidis, K. 8.; LeBaron, P. 0.; Park, I. Chem. Mater.
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86

4.1

CHAPTER 4

Synthesis and Characterization of Cancrinite, Saponite

and Their Phase Mixtures

4.1 Introduction

The synthesis of smectite clay minerals has drawn much attention in
recent years in order to provide derivatives of high purity and adjustable
composition."6 Vogels and co-workers first reported the synthesis of saponites at
90°C from a silica-alumina gel and a solution containing urea and an M2*-nitrate
(M2*=Mg, Zn, Ni, Co).6 Compared to naturally occurring smectite-type clays,
synthetic saponites exhibit improved catalytic activity and selectivity as acid
catalysts for alkylation due to the small and uniform size of clay platelets
(typically less than 50 nm), disordered nanolayer stacking, and controllable
chemical compositions that result in excellent transport properties of organic
compounds. Recently, it has been found by Xue7 that synthetic Mg-saponite
based on Vogels’s methods improves the tensile properties of both rubbery and
glassy epoxy matrices without the need for the organic modification of the clay
surface.

Cancrinite (CAN) is found in nature as a porous zeolite with the
composition Nai;Ca[AISiO.1]6COa(H20)2.8 The cancrinite framework is easily built
from ABAB stacking of hexagonal arrays of 6-membered rings of SiO4 and NO,

tetrahedra along the c-axis (see Figure 1.14).9'10 Small 11-hedral cages (s-cages)

87

and a big 12-ring channel along the [001] direction are the results of this
arrangement of sheets. lsomorphous positioning of aluminum in tetrahedral
coordination results in negative charges on the framework counterbalanced by
cations held within the cavities and channels. The big open channels host Na"
and Ca2+ cations, as well as 0032' and other anions, while the s-cages contain
chains of alternating charge-balancing cations and water molecules.”12

A few researches have been reported on the synthesis of cancrinite under
mild hydrothermal conditions. Preparation of large amounts of pure-phase
canrinite can be difficult due to the formation of disordered cancrinite and an
intermediate phase between cancrinite and sodalite, and the co-crystallization of
sodalite.9'1°"3"19 Pure nitrate cancrinite crystals less than 100 nm in length were
prepared by Navrotsky et al.20'21 based on a simple synthetic route at 90°C,
without the formation of intermediate phase or sodalite. This can be attributed to
the difference in concentrations of the starting solutions and the fact that no
crystalline precursors (e.g. kaolinite, zeolite A) were used in their synthesis
process, unlike previously reported synthesis.9"°'13‘"

We report herein the investigation of the synthesis of Mg-saponite
(denoted SAP) using sodium hydroxide instead of urea as base source, which
represents an economic benefit. Pure delaminated SAP platelets can be
achieved at low NaOH concentration at 90°C. Interestingly, with increasing

alkalinity, new phases were formed along with the co-crystallization of saponite.

Garronite (GIS framework type zeolite) first emerged together with saponite when

88

the NaOH concentration was doubled. Then cancrinite (denoted CAN) was co-
crystallized along with saponite upon further increasing the alkalinity.

Another purpose of the present work is to investigate the low-temperature
synthesis of single-phase cancrinite. Therefore, we also investigated the impact
of Mg” concentration on saponite and cancrinite phase formation, as well as the
effect of Si/Al ratios in the initial liquor on the average crystal dimensions of
cancrinite material. With decreasing Mg amount, the crystallinity of cancrinite is
improved. Also, a higher fraction of cancrinite appears in the product mixtures.
Upon eliminating Mgz” entirely from the synthesis gel, singe—phase cancrinite with
rod-like morphology up to several um in length was synthesized. Decreasing the
Si/Al ratio tends to facilitate the growth of smaller CAN crystals (~ 100 nm). The
chemical composition of cancrinite is also dependent on the Si/Al ratio within the

synthesis gel.

4.2 Experimental
4.2.1 Materlals

Sodium silicate containing 27 wt.% Si02 and 14 wt.% NaOH and
AI(NO3)-9H20 were purchased from Aldrich Chemical Co.. Mg(NOa)2.6H20 was
purchased from Columbus Chemical Industries, Inc. All other chemicals were
purchased from Aldrich Chemical Co. and used without further purification.
4.2.2 Saponite Synthesis

Synthetic saponite was prepared at 90°C according to previously

described methods6 using water glass solution (27 wt% silica, 14 wt% NaOH),

89

AI(N03)3-9HZO, Mg(N03)2-6H20 and NaOH as the source of base. The molar
ratio of Si : Al : Mg : NaOH (not including NaOH in water glass solution) is 3.6 :
0.4 : 3.0 : 5 per 400 moles of water.

In a typical synthesis, 5.50 g of sodium silicate solution (0.02475 mol silica)
and 1.031 9 (0.00275 mol) of AI(NO3)-9HZO were mixed in 49.5 g of deionized
water at room temperature for 1.5 h. Then 1.37 9 (0.03425 mol) of NaOH was
introduced to the white suspension, followed by another 1.5 h of stirring at room
temperature. The mixture was then brought to 90°C and heated under stirred
condition for 30 min, followed by the introduction of 5.28 9 (0.0206 mol) of
Mg(NOa)2-6H20. The mixture was refluxed under stirring at 90°C for 24 h in a 250
mL flask. After that, the product was centrifuged, triple washed with deionized
water, and dried at 80°C under N2 flow.

4.2.3 Synthesis of Phase Mixture of Saponite and Cancrinite

The synthesis steps are similar to those for saponite synthesis, except that
various amounts of sodium hydroxide and magnesium nitrate are used to
investigate the influence of concentration on the saponite to cancrinite product
ratio. The reactions were performed at different time periods ranging from 24 h to
168 h.

4.2.4 Cancrinite Synthesis

Synthetic cancrinite was prepared at 90°C using water glass solution (27
wt% silica, 14 wt% NaOH), Al(N03)3-9H20, and NaOH. The molar ratio of Si :
Al : NaOH (not including NaOH in water glass solution) is x : 0.4 : 30 per 400

moles of water for x = 3.6, 2.4, 1.6, 0.8, 0.4. The mixture was refluxed under

90

stirring at 90°C for reaction periods of 72 and 168 h. After that, the product was
centrifuged, triple washed with deionized water, and dried at 80°C under N2 flow.
4.2.5 Characterization

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 200B
diffractometer equipped with Cu Ko X-ray radiation and a curved crystal graphite
monochromator, operating at 45 kV and 100 mA. Powder samples were pressed
onto a glass X-ray sample holder.

N2 adsorption-desorption isotherms were recorded at -196°C on a
Micrometritcs Tristar 3000 sorptometer. Prior to analysis samples were
outgassed at 150°C and 10‘6 Torr for a minimum of 12 h. BET surface areas
were calculated from the linear part of the BET plot and BJH pore sizes were
obtained from adsorption isotherms.

Transmission electron microscopy (TEM) images were obtained on a
JEOL 2200FS field emission microscope with a ZrONV Schottky electron gun and
an accelerating voltage of 200 W. The powdered samples were sonified in
ethanol, and dripped onto 300 mesh copper grids.

29Si and 27Al MAS NMR spectra were obtained at 79 MHz on a Varian
VXR-4008 solid-state NMR spectrometer equipped with a magic angle-spinning
probe. Sample was spun at 4 kHz for each measurement. The pulse delay for
29Si MAS NMR was 400 s, and chemical shifts were referenced to talc. The
pulse delay for 27Al MAS NMR was 0.50 s, and a 0.10 M aqueous AI(NO3)3

solution was used as the chemical shift reference. FTIR spectra of samples

91

dispersed in KBr disks were recorded at ambient temperature on a Mattson

Galaxy 3000 FTIR spectrometer over the range 400-4000 cm".

4.3 Results and Discussion

The wide angle M! X-ray diffractions demonstrate the phase transition

sequence of saponite-garronite-cancrinite at‘90°C as a result of increasing the

NaOH concentration in the starting solutions (Figure 4.1). The characteristic 001

Si:AlegzNaOH

    

 

 

3.6 :0.4 :3.0:30 3 days CAN+SAP

3.6 :0.4 :3.0 :20 3 days GIS+CAN+SAP

Intensity (a.u.)

 
   

 

 

 

3.6:0.4:3.0:10 3days GIS+SAP

 

3.6:0.4:3.0:5 lday SAP

 

T

10 20 3O 4O 50

6O 7O
29 (degree)

Figure 4.1 XRD patterns showing different crystalline product formation as a

function of increasing alkalinity at 90°C. The ratios represent the stoichiometry of

Si:AlegzNaOH in the starting solutions.

92

 

Figure 4.2 TEM image of synthetic SAP formed at 90°C after a reaction time of 1

day.

reflection is absent for saponite clay, indicating the absence of regular platelet
stacking. SAP has small platelets with dimension of 20-30 nm with almost no
stacking, as shown by the TEM image in Figure 4.2. When the NaOH
concentration is doubled, a new phase garronite (denoted GIS) forms along with
SAP. The co-crystallization of saponite is verified by the presence of an in-plane
060 reflection characteristic of clay. Garronite is replaced completely by
cancrinite upon increasing the NaOH concentration six-fold. Increasing the NaOH
concentration four-fold results in the formation of three phases, namely, garronite,
saponite and cancrinite. Further increasing the alkalinity results in a reduction of
crystallinity and a higher amorphous background in the X-ray patterns for the
CAN/SAP phase mixture (not shown here). Therefore, in all subsequent studies

of this phase mixture, the molar ratios of SizNaOH are maintained at 3.6230.

93

 

 

211
101
300

110 .
Si:AlegzNaOH

 

 

002
400 401 302

210 3.620.420230 7 days l
w LJ :

 

 

 

 

 

  
   
 
   

 

 

 

 

 

 

 

i
A I
=2 l
~“> l
g; 3.6:0.4:0z30 3days s
m r.
.E':
3.6:0.4:0.75:30 6days
‘ :
1
3.6:0.4:0.75:30 3days *
J . 3.6:0.4:1.5:30 3days l
l ' 3. :0.4:3z3r 4 -a s i
A A , A A , ill-uh , _ ‘ ___- --3_'~6:Q'423:30' 3days
10 20 30 40 50 60 70
29 (degree)

Figure 4.3 XRD patterns of crystalline products formed upon decreasing the
Mg(N03)2 content of the reaction mixture. For each product, the molar ratios of
Si:Al:Mg:NaOH in the synthetic solutions are labeled, as well as the reaction time

at 90°C.

94

The effect of magnesium nitrate concentration on the phase formation was
also studied. The X-ray diffraction patterns of the crystalline products formed at
different Mg ion concentrations are shown in Figure 4.3. With decreasing the Mg
amount in the initial reaction liquor, the intensities of the reflection peaks from
cancrinite increase, indicating a significant improvement in the crystallinity of
cancrinite, as well as higher proportion of cancrinite in the CAN/SAP phase
mixture. The existence of saponite in the products together with cancrinite can be
easily seen from TEM images (Figure 4.4 and Figure 4.5). Figure 4.4 parts a and
b show that after 4 days of aging, rod-like crystals of cancrinite with several
hundred nm in length emerge together with platy saponite particles. These
crystals were formed from the starting Si:AlegzNaOH stoichiometry of
3.6:0.4:3z30. Evidence supporting the presence of amorphous silica in the
product mixture, as suggested by the high background in the X-ray powder
patterns shown in Figure 4.1, is provided by the TEM image of Figure 4.4 c. A
TEM-XEDS analysis of this selected area indicated a Si:O molar ratio of
3315:6685, in agreement with pure silica. When the magnesium nitrate
concentrations were reduced to half and one fourth, co—crystallization of saponite
platelets is still observed in TEM imagings (Figure 4.5 a, b, c and d). Until
completely removing magnesium nitrate from the starting reaction solutions, no
SAP nanolayers are observed in the reaction product (cf. Figure 4.6 a, b, and 0),
suggesting the formation of pure-phase cancrinite. Moreover, some cancrinite
crystallites have grown to a rather large size of several micrometers in length

after a 7 day reaction period.

95

100 nm

 

Figure 4.4 TEM images of CAN/SAP phase mixture formed from a reaction
stoichiometry of Si:Aleg:NaOH=3.6:0.4:3:30 at 90°C after a reaction time of 4
days, and the corresponding ratios of SilAl=9, Mg/(Si+Al)=0.75, and OH'
/(Si+Al+Mg)=4.3.

96

100 nm

100 nm

 

Figure 4.5 TEM images of CAN/SAP phase mixtures formed at 90°C after a
reaction time of 3 days. (a) and (b) are for the CAN/SAP mixture made from
starting materials with the composition Si:Al:Mg:NaOH=3.6:0.4215:30, and the
corresponding ratios of SilAl=9, Mg/(Si+Al)=0.375, and OH'/(Si+Al+Mg)=5.45; (c)
and (d) are for the CAN/SAP mixture made from Si:Al:Mg:NaOH=3.6:0.4:0.75:30,
and the corresponding ratios of SilAl=9, Mg/(Si+Al)=0.188, and OH'
/(Si+Al+Mg)=6.32.

97

 

Figure 4.6 TEM images of single-phase cancrinite formed from a reaction

stoichiometry of Si:Aleg:NaOH=3.6:0.4:0:30 at 90°C after a reaction time of 7
days, and the corresponding ratios of SilAl=9, Mgl(Si+Al)=0, and OH'
/(Si+A|+Mg)=7.5.

98

The full width at half maximum (FWHM) of the reflection peaks were
measured and compared for cancrinite materials crystallized at different Mg
levels (Figure 4.3). It’s found that the widths of the diffraction peaks for single-
phase cancrinite are the narrowest among all of the cancrinite products. This is in
agreement with the largest crystallite size, found by the TEM observation.

The lattice structure of saponite and cancrinite were further characterized
using solid state NMR. The 29Si MAS NMR spectrum of saponite (of. Figure 4.7)
exhibits the presence of two resonances. The -85.4 ppm line is assigned to Si
centers linked through bridging oxygen atoms to three tetrahedral Si ions and
one tetrahedral Al ion, and the -92.7 ppm resonance is ascribed to Si linked
through bridging oxygen to four Si neighbors. The 27Al MAS NMR spectrum of
SAP contains two signals at chemical shifts of 61.6 ppm and 10.5 ppm, indicating
that Al substitution occurred in both tetrahedral and octahedral sheets of the 2:1
layered silicate structure. In the case of a CAN/SAP phase mixture, 29Si and 27Al
solid state NMR may give qualitative information as to the relative proportions of
the phases present. The sharp signal at -87.3 ppm correlates well with the
alternating ordering of the framework Si and AI atoms in cancrinite crystals. The
shoulder line at -92.7 ppm might be from saponite, or the formation of Si-rich
cancrinite, wherein the Si/Al ratio is more than 1 and some of the silicon centers
are linked through bridging oxygen to three Al and one Si. The latter assignment
of -92.7 ppm resonance is verified by the quantitative analysis of cancrinte
particles in the mixed CAN-SAP phase carried out by XEDS-TEM (cf. Figure 4.8),

indicating a Si/Al ratio of 1.25. Moreover, according to X-ray analysis, small

99

Figure 4.7 298i and 27Al MAS NMR spectra of (a) SAP formed after 1 day
reaction time (3.6:0.4:3:5), (b) CAN/SAP phase mixture formed after a reaction
time of 4 days (3.6:0.4:0.75:30) and (0) pure CAN made after 3 days
(3.6:0.4:0:30). The ratio in the parentheses represents the molar ratio of

Si:AlegzNaOH in the starting solution. The asterisks indicate spinning sidebands.

100

27
29 - Al
3' Saponite

 

UV”

-40 -60 -80 -100 -120 -140 -160 250 200 150 100 50 O -50 -100 -150

 

 

 

 

 

 

 

 

PPm ppm
298i 27Al N
Cancrinite/Saponite
WMMJ Wavy
rfiV'I”"F""IV"'I13V'I""I WWW-I
'40 '60 ’30 '100 “120 '140 '160 250 200 150 100 50 0 -50 -100 -150
PPm PPm
298i 27A| N
Cancrinite

 

 

 

w

WWW-W

40 -60 -80 -100 -120 -140 460 250 200 150 100 50 0 -50 -100 -150
PM" PM

6

101

 

Element Wei ght% Atomic%
O K 44.61 57.09

Na K 11.17 9.95

Mg K 1.57 1.32

Al K 18.56 14.09

Si K 24.08 17.55
Totals 100.00

 

 

   

3:2

6

$52“st Cursor: OI'IXHOV
Figure 4.8 TEM image and X-ray energy dispersive spectrum (XEDS) of
cancrinite crystals in the CAN/SAP phase mixture formed from a reaction mixture
with the composition Si:Aleg:NaOH=3.6:0.4:0.75:30 after 4 days. Unlabeled
peaks are from copper grid and carbon film. Quantitative analysis result is

summarized in the inserted table and note the SilAl=1.25.

amounts of magnesium are also incorporated into the cancrinite framework
structure in addition to sodium cations. The 27Al MAS NMR spectrum only

exhibits a single line at 57.1 ppm for the aluminum in tetrahedral coordination.

102

The fact that a resonance for aluminum in octahedral coordination is absent
suggests a low saponite population in the CAN/SAP phase mixture. The 29Si
MAS NMR spectrum of pure cancrinite shows the same two resonances as in
CAN/SAP, and in this case the two resonances arise because a large excess of
silica in the reaction mixture with Si/Al=9 facilitates the formation of a Si-rich

cancfinhe.

 

 

 

 

Absorbance (a.u.)

 

 

 

 

, g

3900 3400 2900 2400 1900 1400 900 400
Wavenumber (cm")

Figure 4.9 FTIR spectra of (a) pure saponite after 1 day of aging from initial
solutions with a Si:AlegzNaOH stoichiometry at 3.6:0.4:3:5, and (b) CAN/SAP
phase mixture after 4 days of aging from initial solutions with a Si:AlegzNaOH

stoichiometry at 3.6:0.4:0.75:30.

103

The FTIR spectra of pure SAP and CAN/SAP phase mixture are
compared in Figure 4.9. The water bending mode around 1630 cm'1 and a broad
OH stretching band between 3100 and 3600 cm'1 are found in both samples. The
IR spectroscopic result of pure saponite shows absorption bands at 1020 and
653 cm", clearly indicating the typical asymmetric and symmetric T-O-T
vibrations. The vibrations of the cancrinite phase dominate the IR spectrum of the
CAN/SAP phase mixture (cf. Figure 4.9 b). For nitrate cancrinite the absorbance
for the asymmetric stretch VA“) for the Al-O-Si framework is split into four peaks
at 1120, 1039, 998 and 963 cm“. The region of 400-800 cm" is known as the
“fingerprint” region.17 The characteristic fingerprint peaks for cancrinite occur at
approximately 575, 622, 683, and 770 cm". In addition to the weak bending
mode near 820 cm", the presence of nitrate groups in the zeolite channels can
be better detected by the strong vas vibration at 1425 cm". The weak absorption
at 1383 cm’1 has been attributed to small amounts of precipitated sodium nitrate
the outer surface of the crystals.22 Another potential occluded anion for our
cancrinite stnicture is 0032'. This ion may form through the reaction of
atmospheric CO2 with the highly concentrated NaOH in the reaction mixture.
However, carbonate was not incorporated to a detectable extent because the two
C032' stretching bands9 at 1410 and 1455 crn'1 are not observed here, indicating
that NO»; is the major anion in the channels of the cancrinite stnicture. Cancrinite
vibrations almost completely mask the vibrations of the minor saponite phase.
This further confirms the low proportion of saponite in the phase mixture, which is

consistent with the 298i and 27Al MAS NMR results.

104

Table 4.1 Textural properties of saponite, cancrinite and CAN/SAP phase

 

 

mixtures
. . BET Total BJH
Si:AlegzNaOH Reaction .
Product ratio of reaction time surface p|ore adsorption
mixture (days) arzea vo u3me pore Size
(00 lg) (cm Ig) (nm)
SAP 3.6:0.4:3:5 1 875 1.54 1.8
CAN/SAP 3.6:0.4:3230 3 329 0.42 1 .4
CAN/SAP 3.6:0.4:1.5:30 3 227 0.39 1.8
CAN/SAP 3.6:0.4:0.75:30 3 134 0.31 2.0
CAN 3.6:0.4:0:30 7 10.2 0.068 -
CAN 3.6:0.4:0:30 3 8.9 0.060 -
CAN 2.4:0.4:0:30 3 21.8 0.14 -
CAN 1.6:0.4:0:30 3 40.0 0.17 -
CAN 0.8:0.4:0:30 3 36.6 0.17 -
CAN 0.4:0.4:0:30 3 21.3 0.095 -

 

BET surface areas and pore volumes of single-phase saponite and
cancrinite, and CAN/SAP phase mixtures are summarized in Table 4.1. The N2
adsorption—desorption isotherms (Figure 4.10) show a Type H3 hysteresis loop
for pure saponite. The BET surface area and pore volume of the pure saponite is
875 m2/g and 1.54 cm3/g, respectively, resulting from the irregular stacking and
small in-plane dimension of clay platelets.

The pore aperture in the cancrinite framework structure is 5.9><5.9 A,
which is big enough for the access of N2 molecules. But the presence of
intercalated nitrate anions in CAN severely blocks the big channels and prevents
N2 molecules from accessing the interior of the pores. This accounts for the low
uptake of N2, as reflected in the gas adsorption isotherm behavior of pure

cancrinite (Figure 4.10). The very low surface area of pure cancrinite zeolite

105

(10.2 mzlg) is due to adsorption on the external surfaces of large particles up to
several um in length (Figure 4.6 a, b and c). As expected, CAN/SAP phase
mixtures have the surface area between the pure saponite and cancrinite.
Moreover, lower surface area of CAN/SAP (134 m2/g) made from a reaction
mixture with the composition Si:Aleg:NaOH=3.6:0.4:0.75:30 illustrates the
higher population of CAN in this phase mixture in comparison to that of CAN/SAP

with higher surface area of 227 or 329 mzlg.

1000 -~

.g’
---«2-~ SAP 3.6:0.4:3:5 °
900 2' + CAN/SAP 3.6:0.4:3230 1:,
+ CAN/SAP 3.6:0.4:1.5:30 5
a: 800 - -— CAN/SAP 3.620.420.75230 .
'5 + CAN 3.6:0.4:0:30
7 700 1 3";
3 .
a, 600 -
3 500 2 fi‘
> .7 :‘r‘r’ '
8 400 '
g >0 *7 A
8 300 « .22”
(U ~
f

 

 

 

0 0.2 0.4 P/POO'G 0.8 1

Figure 4.10 N2 adsorption-desorption isotherms of pure saponite and cancrinite,
and their phase mixtures. The ratios represent the molar ratios of Si:AlegzNaOH
in the reaction mixtures. The reaction time for pure saponite is one day and all

the other products are made after a reaction period of 3 days.

106

 

 

 

 

 

 

 

 

 

 

211 T

101 300 1

110 I

210 l 1 ’ l l

SizAl=9 l

l

,. l

:3

s . 1
.Q SLA1=6

g l

l

,5 '1

l

Si-Al=4 1

l

1

SizAl=2 g

1

SizAl=l l

10 20 30 40 50 60 70

29 (degree)

Figure 4.11 XRD patterns of cancrinite products formed as a function of the Si/Al

ratio within the starting system after 3 days of reaction time under Mg—free

condition.

107

2931

 

 

W

-40 -60 -80 -100 -120 -140 -160
ppm

 

Figure 4.12 29Si MAS NMR spectrum of pure cancrinite synthesized from the

educts with the composition of Si:Al:NaOH=1.6:0.4:30 under Mg-free condition.

Additional experiments were performed at various Si/AI ratios of 6, 4, 2
and 1 within the starting reaction mixtures. Powder XRD patterns of crystals of
cancrinites formed at different Si/Al ratios after 3 days reaction time are shown in
Figure 4.11. For each reaction, the molar ratio of AlzNaOH was maintained at
0.4:30. It can be seen that relative peak intensities are dependent upon the SilAl
ratio. For example, the relative intensity of the 300 peak decreases in
comparison to the 101 and 211 reflections as the Si/Al ratio is increased from 1.0
to 9.0. This result most likely reflects a varying chemical composition of cancrinite
crystals. As discussed before, a large excess of silica facilitates the formation of

Si-rich cancrinites, which may be the case for Si/Al=9 and 6. XEDS-TEM analysis

108

reveals a Si/Al ratio of 1.02 for cancrinite formed from a reaction mixture with
Si/Al=4, which is very close to the ratio expected for the alternating occupation of
tetrahedral sites by Si and AI. As a result of the Si(OAl)4 framework building units,

a single 2""Si line at -87.3ppm is found (Figure 4.12) for cancrinite grown from a

reaction mixture with a Si/Al ratio of 4.

100 nm

   

Figure 4.13 TEM images of cancrinite crystals formed at Si:Al:NaOH=1.6:O.4:30

under Mg-free condition.

The Si/Al ratios in the reaction mixture also have an influence on the
textural properties of cancrinite products. As can be seen from Table 4.1, the
effect of the Si/Al ratio on the surface areas and pore volumes is significant.
Cancrinites prepared at a Si/Al ratio of 9 and 4 exhibit the largest differences in
textural characteristics. TEM investigations of products made at Si/Al=4 show a

large number of small crystals of cancrinite less than 100 nm in length, along with

109

large needle-like crystals up to several micrometers in length (cf. Figure 4.13).
The increased fraction of small crystallites explains the comparatively high

surface area of 40 m2/g.

4.4 Summary

The synthesis of saponites was carried out at 90°C using sodium
hydroxide as the base source in a mixed gel containing sodium silicate,
aluminum nitrate and magnesium nitrate. At low alkalinity corresponding to a CH
ISi ratio of 1.39, pure saponite with small in-plane dimension of 20-30 nm and
disordered stacking is achieved. The use of NaOH as a base represents a
significant cost-effective improvement over the previously reported synthesis of
saponite which used urea as the base sources. The present work demonstrates
that the slow release of OH' through the thermal hydrolysis of urea is not a
prerequisite for the nucleation and growth of saponite nanolayers. Equivalent 2:1
layered silicate nanolayers are obtained through the direct addition of hydroxide
ions to the reaction mixture.

In an effort to improve the layer aspect ratio and stacking order of the
saponite nanolayers, the NaOH concentration in the reaction mixture was
increased over the range OH'lSi=2.78 to 8.33. However, some new zeolite
phases appear along with saponite upon increasing alkalinity. Garronite first
emerged together with saponite when the NaOH concentration was doubled (OH'

lSi=2.78). Increasing the NaOH concentration by six-fold (OH'ISi=8.33) results in

110

the co-crystallization of cancrinite with saponite. This finding represents the first
example of the direct synthesis of a cancrinite-saponite phase mixture.

The appearance of garronite and cancrinite as co-products demonstrates
that these aluminosilicate phases compete kinetically with the nucleation of
saponite. The depletion of silicon and aluminum due to garronite or cancrinite
formation leaves the reaction mixture rich in magnesium. The enrichment in
magnesium allows the nucleation of saponite to compete with the nucleation of
the zeolite phases.

The concentration of magnesium ions also has an influence on the
formation of mixed CAN-SAP phase. Decreasing the Mg(NOa)2 amount leads to
an improvement of crystallinity of cancrintie and a higher proportion of cancrinite
in the CAN/SAP phase mixture, as evidenced by the increased intensities of
diffraction peaks from cancrinite in the X-ray powder patterns. Removing Mg2+
completely from the reaction mixture results in the formation of single-phase
cancrinite with rod-like morphology up to several pm in length. Moreover, both
2"Si MAS NMR and composition analysis indicate the formation of Si-rich
cancrinite, wherein the Si/Al ratio in the CAN framework structure is greater than
1.0. This can be explained by the fact that a large excess of silica within the
synthesis solution with Si/Al=9 facilitates the crystallization of a Si-rich cancrinite.
Additional experiments were performed at various Si/Al ratios of 6, 4, 2 and 1 to
study the effect of the Si/Al ratio on the chemical composition, as well as the
textural properties of cancrinite zeolite. XEDS-TEM analysis reveals a Si/Al ratio

of 1.02 for cancrinite crystals formed from a reaction mixture with SilAl=4.

111

Therefore, relatively low Si/Al ratio (e.g. 4, 2 and 1) within the synthesis gel leads
to the formation of cancrinite with an ideal Si/AI ratio of 1.0. Also, with decreasing
the Si/AI ratio, higher fraction of small cancrinite crystallites tends to crystallize
with dimensions within 100 nm in length, as shown by the TEM images. The
decrease in the average crystal size with decreasing Si/AI ratio results in an

increase in textural properties.

112

References

(1) Kloprogge, J. T.; Komameni, S.; Amonette, J. E. Clays Clay Miner. 1999, 47,
529.

(2) Fajnor, V. S.; Kuchta, L. Chem. Papers, 1988, 42, 339.

(3) Kloprogge, J. T.; Breukelaar, J.; Jansen, B. H.; Geus, J. W. Clays Clay Miner.
1993.41,103.

(4) Vogels, R. J. M. J.; Breukelaar, J.; Kloprogge, J. T. Clays Clay Miner. 1997,
45, 1.

(5) Prihod’ko, R.; Sychev, M.; Hensen, E. J. M.; Veen, J. A. R. V. Stud. Surf. Sci.
Catal. 2002, 142,271.

(6) Vogels, R. J. M. J.; Kerkhoffs, M. J. H. V.; Geus, J. W. Stud. Surf. Sci. Catal.
1995,91,1153.

(7) Xue, S.; Pinnavaia, T. J. Micropor. Mesopor. Mater. 2008, 107, 134.
(8) Strunz, H. Mineralogische Tabellen 1997, 7.

(9) Buhl, J.-Ch.; Hackbarth, K.; Gesing, Th.M.; Fechtelkord, M.; Stief, F. Micropor.
Mesopor. Mater. 1999, 30, 347.

(10) Weller, M. T.; Kenyon, N. J. Micropor. Mesopor. Mater.2003, 59, 185.

(11) Merlino, S.; Mellini, M.; Bonaccorsi, E.; Marco, P.; Pasero, M.; Leoni, L.;
Ortandi, P. Am. Mineral. 1991, 76, 2003.

(12) Bonaccorsi, E.; Ventura, G. D.; Bellatreccia, F.; Medino, S. Micropor.
Mesopor. Mater. 2007, 99,225.

(13) Buhl, J.-Ch.; Hermeler, G.; Hoffmann, W. Catal. Today 1991, 8, 415.

(14) Buhl, J.-Ch.; Fechtelkord, M.; Hackbarth, K. React. Kinet. Catal. Lett. 1998,
65, 33.

(15) Dann, S. E.; Armstrong, J. A. Micropor. Mesopor. Mater. 2000, 41,89.

(16) Gerson, A. R.; Barnes, M. C.; Addai-Mensah, J. Micropor. Mesopor. Mater.
1999, 31,287.

113

(17) Gerson, A. R.; Barnes, M. C.; Addai-Mensah, J. Micropor. Mesopor. Mater.
1999. 31,303.

(18) Zheng, K.; Gerson, A. R. J. Cryst. Growth 1997, 171, 209.

(19) Zheng, K.; Gerson, A. R.; Addai-Mensah, J.; Smart, R. S. J. Cryst. Growth
1997, 171, 197.

(20) Navrotsky, A.; Liu, 0.; Xu, H. Micropor. Mesopor. Mater. 2005, 87, 146.

(21) Navrotsky, A.; Liu, 0.; Jove-Colon, C. F.; Bonhomme, F. Micropor. Mesopor.
Mater. 2007, 98,227.

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J. Alloy. Compd. 2000, 305, 93.

114

CHAPTER 5

Influence of the Alkalinity and Temperature on the Synthesis of

Cancrinite-Sodalite System

5.1 Introduction

Cancrinite and sodalite are very interesting zeolites because they have
very low Si/Al ratio (Si/Al=1.0) and they're able to incorporate anions in their
structure. Cancrinite (CAN) is found in nature with the composition
NaeCa[AlSiO4]6003(H2O)2, while sodalite (SOD) with the composition
Nag[AlSiO4]r,Cl2.1 The cancrinite framework can be easily built from ABAB
stacking of hexagonal arrays of 6-membered rings of $104 and AIO4 tetrahedra
along the c-axis (see Figure 1.14).2'3 Small 11-hedral cages (e-cages) and a big
12-ring channel along the [001] direction are the results of this arrangement of
sheets. On the other hand, stacking these layers in an ABCABC sequence leads
to the cubic sodalite framework with its characteristic B-cage (Figure 1.14). In
cancrinite, isomorphous positioning of aluminum in tetrahedral coordination
results in negative charges on the framework counterbalanced by cations held
within the cavities and channels. The big open channels host cation-anion pairs,
while the e-cages contain chains of alternating charge-balancing cations and
water molecules.“-5

The structural and compositional similarities between the sodalite and

cancrinite often results in the co-crystallization of both aluminosilicates. Therefore,

115

the preparation of large amounts of pure-phase canrinite can be difficult due to
the formation of disordered cancrinite and an intermediate phase between
cancrinite and sodalite, and the co-crystallization of sodalite.2"°"“*23 Buhl et al.“7
investigate the best conditions for the synthesis of sodium carbonate cancrinite
as a single phase. Cancrinite was only shown to form as a pure phase at
temperatures not below 200°C. The concentration of reactants and the alkalinity
of the hydrothermal solution had a much lower influence on the phase formation.
An intermediate phase with one-dimensional stacking disorder of the (AlSiO4)n
layers was obtained at lower reaction temperature of 80°C. Weller reported the
effect of calcium on the hydrothermal synthesis of the cancrinite framework in the
presence of carbonate at 150 and 220°C.3 Ideal cancrinite was formed under
calcium-free conditions at 220°C, whereas a disordered cancrinite/sodalite
intergrown phase was observed at 150°C. The addition of calcium into the
reaction mixture leads to its incorporation into the cancrinite structure. At high
concentration of calcium a new phase (NaCaSianH) is formed. A phase
transition sequence of amorphous phase—zeolite A— sodalite—cancrinite was
reported by Barns and co-workers."'10 The pathway by which sodalite converts to
cancrinite is the result of sodalite dissolution and subsequent cancrinite
precipitation. Pure nitrate cancrinite was prepared by Navrotsky et al.2‘1”26 based
on a simple synthetic route at 90°C, without the formation of intermediate phase
or sodalite. This can be attributed to the difference in concentrations of the

starting solutions and the fact that no crystalline precursors (e.g. kaolinite, zeolite

116

A) were used in their synthesis process, which were used by most of the reported
synthesis.2'3'€*1°'151”922

In the previous chapter, we reported a new phase mixture of saponite and
cancrinite and investigated the influence of Mg(N03)2 and NaOH concentrations
on the synthesis of this mixed phase. Pure cancrinite can be prepared under Mg-
free conditions, because Mg cations are essential for saponite formation, but not
for cancrinite formation. Also, various Si/Al ratios in the reaction mixture affected
the chemical composition and textural property of the resulting cancrinite.

In the present chapter, we investigate the effect of alkalinity and reaction
temperature on the synthesis of nitrate cancrinite and nitrate sodalite. Ideal
nitrate cancrinite is formed in a wide range of reaction parameters, 9.9. NaOH
concentration, temperature, and duration. The concentration of NaOH influences
the product distribution at a reaction temperature of 90°C in the following way:
high OH' concentrations lead preferentially to the formation of single-phase CAN,
while low concentrations produce a CAN-SOD mixture. Also, pure cancrinite
could be synthesized if the aging temperature is not below 40°C. However, a
cancrinite-sodalite mixture with sodalite as, the major phase is formed at room

temperature.

5.2 Experimental

5.2.1 Materials

117

Sodium silicate containing 27 wt.% SiO2 and 14 wt.% NaOH and
AI(N03)-9H2O were purchased from Aldrich Chemical Co.. All other chemicals
were purchased from Aldrich Chemical Co. and used without further purification.
5.2.2 Synthesis of Nitrate Cancrinite

8.55 g of sodium silicate solution (27 wt.% silica, 14 wt.% NaOH) (0.03850
mol silica) and 14.44 9 (0.03850 mol) of Al(NOa)3-9H2O were first added to 87 ml
of deionized water in a sealed polypropylene bottle. A cloudy gel-like solution
was obtained after complete mixing. Then, NaOH as the source of base ranging
between 28.8 g (0.72 mol) and 115.2 g (2.88 mol) was introduced and the bottle
was vigorously shaken to ensure homogeneous mixing. The molar ratio for
sodium silicate, aluminum nitrate, sodium hydroxide (not including NaOH in the
sodium silicate solution) and water is 1:1:18.8—75:125. The resulting gels were
placed in oven heated at 40°C to 90°C in tightly stoppered polypropylene bottles.
Samples were taken from the oven at pre-detennined time periods ranging from
1 day to 7 days. All products were washed using deionized water via
centrifugation 3 times. All samples were dried overnight in oven at 80°C under
nitrogen gas flow.

5.2.3 Synthesis of Phase Mixtures of Nitrate Cancrinite and Nitrate Sodalite

Synthesis steps are similar to those used for the synthesis of nitrate
cancrinite except that the molar ratio for sodium silicate, aluminum nitrate,
sodium hydroxide and water in the initial solutions is 1:1:4.69-9.38:125.

5.2.4 Characterization

118

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 200B
diffractometer equipped with Cu Ka X-ray radiation and a curved crystal graphite
monochromator, operating at 45 kV and 100 mA. Powder samples were pressed
onto a glass X-ray sample holder.

N2 adsorption-desorption isotherms were recorded at —196°C on a
Micrometritcs Tristar 3000 sorptometer. Prior to analysis samples were
outgassed at 150°C and 10’5 Torr for a minimum of 12 h. BET surface areas
were calculated from the linear part of the BET plot and BJH pore sizes were
obtained from adsorption isotherms.

Transmission electron microscopy (TEM) images were obtained on a
JEOL 2200FS field emission microscope with a ZrONV Schottky electron gun and
an accelerating voltage of 200 W. The powdered samples were sonified in
ethanol, and dripped onto 300 mesh copper grids.

29Si MAS NMR spectra were obtained at 79 MHz on a Varian VXR-4OOS
solid-state NMR spectrometer equipped with a magic angle-spinning probe.
Sample was spun at 4 kHz for each measurement. The pulse delay for 29Si MAS
NMR was 400 s, and chemical shifts were referenced to talc. FTIR spectra of
samples dispersed in KBr disks were recorded at ambient temperature on a

Mattson Galaxy 3000 FTIR spectrometer over the range 400-4000 cm".
5.3 Results and Discussion

To investigate the effect of the alkalinity of the starting solutions on

cancrinite phase formation, we selected five different NaOH concentrations. The

119

 

S211

S110 Si : Al : NaOH :HzO

A C101 1| C211 CAN/SOD 1:1:4.69:125 3 days 1

 

CAN/SOD 1:1:9.38:125 3 days 1

l
i
l
1
l

 

Intensity (a.u.)
§
Lo.)
8
£3

 

 

 

 

 

 

I CAN 1:1:75:125 3 days i
7 LA.) 01’ J i 1
L_.L._l L225 1. J ”JAMAJW
.— «ems—T— ake—:22--——.——.----s——, —. — .—a——.——4-—. “we. ——r——. ~ .2 .—~-i
20 30 40 50 60 70

29 (degree)

Figure 5.1 XRD patterns of reaction products as a result of varying alkalinity in
the starting solutions after a reaction time of 3 days at 90°C. The ratios represent
the Si:Al:NaOHzH2O stoichiometry of the reaction system. (S=sodalite,

=cancnnne)

120

Figure 5.2 TEM images of (a), (b) and (c) ideal cancrinite crystals formed at 3
different alkalinities of x=75, 37.5 and 18.8, respectively; (d) cancrinite in
CAN/SOD phase mixture formed at x=9.38; (e) and (f) rod-like cancrinite and
cbtton-ball like sodalite in CAN/SOD phase mixture formed at x=4.69. All the

samples were prepared at 90°C after a reaction time of 3 days.

121

 

122

Table 5.1 Synthetic conditions together with the textural properties of reaction

 

 

products
Si:Al:NaOHon2O Reaction Temperature BET :3:
Products in the starting time (°C) surfacze volume
materials (days) area (m lg) (cm3 I 9)
CAN 1 :1 :75:125 3 90 7.9 0.056
CAN 1 :1 :37.5:125 1 90 68.9 0.55
CAN 1 :1 :37.5:125 3 90 41.9 0.41
CAN 1 :1 :37.5:125 7 90 48.6 0.38
CAN 1 :1 :18.8:125 1 90 97.6 0.66
CAN 1 :1 :18.8:125 3 90 74.8 0.49
CAN 1 :1 :18.8:125 7 90 59.6 0.35
CAN/SOD 1 :1 :9.38:125 1 90 24.0 0.13
CAN/SOD 1 :1 :9.38:125 3 90 22.6 0.15
CAN/SOD 1 :1 :4.692125 1 90 71.4 0.19
CAN/SOD 1 :1 :4.692125 3 90 41.0 0.20
CAN 1:1:18.8:125 3 60 109.5 0.70
CAN 1:1:18.8:125 3 40 125.7 0.61
SOD/CAN 1 :1 :18.8:125 3 25 208.6 0.29

 

X-ray powder patterns of the products after 3 days reaction at 90°C are
summarized in Figure 5.1. Pure-phase nitrate cancrinite was formed at relatively
high alkalinity for the reaction systems with the compositions Si:Al:NaOHzH2O
=1:1:x:125, x=75, 37.5 and 18.8. The diffraction patterns of these three
crystalline powders matched the data reported by Buhl15 for ideal nitrate
cancrinite. Crystallization of nitrate cancrinite was almost completed after 1 day
at 90°C, as verified by the X-ray diffraction pattern (not shown here). The
formation of nitrate cancrinite was rather fast considering the low reaction
temperature involved. TEM images in Figure 5.2 (parts a, b and 0) illustrate the

differences in morphology and crystallite size for cancrinites formed at three

123

different alkalinities corresponding to x=75, 37.5 and 18.8. At x=75, cancrinite
grows into rather large rod-like crystals with several hundred nm in length. This
product has a very low BET surface area of 7.9 m2/g (Table 5.1), which is
associated with the large crystallite size. When the alkalinity was reduced to half
and one fourth, mostly rod-like cancrinite crystals exhibit uniform size distribution
within 100 nm in length. Additionally, crystals formed at x=37.5 seem to grow
wider than those at x=18.8, and both of their BET surface areas increase a lot in
comparison to the large crystals prepared at x=75. With longer reaction time,
cancrinite crystals tends to nucleate and grow larger, as indicated by the
decrease in the surface area with increasing reaction period (see Table 5.1).
XRD patterns show that upon decreasing the NaOH concentration further,
some new phases appear. The precipitation of a phase mixture of cancrinite and
sodalite takes place at a Si:Al:NaOHzH2O stoichiometry of 1:1:9.38:125 and
1:1:4.69:125. The reflection peaks of the CAN/SOD phase mixtures indicate the
typical cancrinite reflections in the powder pattern, but with much weaker
intensities for the 101 and 211 peaks. Pure sodalite and cancrinite have very
similar XRD patterns, and the two can be distinguished by the presence of the
cancrinite 101 and 211 diffraction peaks. Barnes and co-workers developed a
simple methodology for determining the cancrinite proportion of a sodalite and
cancrinite phase mixture using powder X-ray diffraction.‘°They used overlapping
C110+S110 and C300+8211 (where C=cancrinite and S=sodalite) diffraction
peaks as the reference peaks and the C101 and 0211 peaks as variables

because they’re unique to cancrinite. The relative peak intensities for the C101

124

and 0211 diffraction peaks, when normalized against the C300+8211 peak, were
each found to increase linearly with increasing CAN concentration. Here, Table
5.2 shows an increase in the relative peak height of the characteristic cancrinite
101 and 211 diffraction peaks with increasing reaction time, which can be
attributed to an increased CAN population in the phase mixture. This result
suggests that sodalite-to-cancrinite phase transformation can occur in the

reaction system.

Table 5.2 Relative intensities of four diffraction peaks for sodalite and cancrinite

phase mixtures

 

 

Si: AlzNaOHzH2O Aging Relative diffraction peak intensity

 

Products _ , , , time C110+ C300+
-1.1.x.125 (days) S110 0101 $211 0211
_ 1 41.6 14.9 100 16.5
CAN/SOD ”938 3 42.4 20.2 100 25.5
_ 1 62.7 4.1 100 11.5
CAN’SOD ”4°69 3 57.5 16.1 100 20.0

 

A TEM image of aggregated cancrinite in the CAN/SOD phase mixture
(Figure 5.2 d) shows that it retains the rod-like morphology, but forms much
larger crystallites in comparison to cancrinite formed at x=37.5 and 18.8. In
addition, the hkl reflections for ideal cancrinite at x=37.5 and 18.8 are broader
than those for CAN/SOD at x=9.38, indicative of a much smaller crystalline
dimension and in agreement with the TEM observation. In the case of CAN/SOD
at x=4.69, sodalite with “cotton-ball” like morphology was also observed in

addition to the rod-like cancrinite morphology (Figure 5.2 e, f).

125

 

 

CAN/SOD 1-1-4.69-1 25 30
W

 

 

CAN/SOD 1-1-9.38-125 30

 

Absorbance (a.u.)

 

 

CAN 1-1-18.8-125 SD

A
i
l -

l
l
l
l
|

    

-_-__.l

4000 3600 3200 2800 2400 2000 1600 1200 800 400
Wavenumber (cm'1)

 

 

 

7 _7

Figure 5.3 FTIR spectra of selected products formed at varying alkalinity in the
starting solutions after 3 days reaction at 90°C. The ratios represent the

Si:Al:NaOHzH2O stoichiometry in the reaction system.

A further indication of a cancrinite-sodalite mixture formation is provided
by IR spectroscopic results (Figure 5.3). The broad band at 3100-3600 cm'1 and
the bending mode at 1630 cm'1 can be assigned to water molecules. For the
ideal nitrate cancrinite prepared at x=18.8 and a reaction time of 3 days, the
absorbance for the asymmetric stretch who for the Si-O-Al framework is split into

four peaks at 1120, 1039, 998 and 963 cm", in accordance with the reported IR

126

assignments for cancrinite.1°"5'20 In the phase mixture of cancrinite and sodalite
formed at x=9.38 and 4.69, absorptions at 1120, 1040 (shoulder) and 996 cm‘1
can be resolved. Pure sodalite should exhibit a single peak at 996 cm‘1 for the
asymmetric framework stretch. The region of 400-800 cm‘1 is known as
“fingerprint” region. The characteristic fingerprint peaks for pure cancrinite occur
at 575, 622, 683, and 770 cm“1 as shown in Figure 5.3. The absorptions at 683
and 770 crn'1 correspond to the symmetric Al-O stretch for the Si-O-Al framework
and the bands at 575 and 622 cm”1 arise from the vibration of the parallel 4- or 6-
membered double rings of cancrinite. For sodalite, fingerprint peaks occur at 665,
704 and 731 cm". For CAN/SOD phase mixtures made at x=9.38 and 4.69,
characteristic fingerprint peaks for both cancrinite and sodalite can be resolved in
the region of 500-800 cm". In addition to the weak bending mode near 820 cm",
the presence of nitrate groups inside the cancrinite channels can be better
detected by the vas vibration at 1425 cm". The absorption band at 1383 cm'1
could be attributed to small amounts of precipitated sodium nitrate on the outer
surface of the crystals. lntercalated nitrate within the sodalite cages also exhibits
a band at 1383 cm".19'27 Thus, the intensity of this peak increases relative to the
peak at 1425 cm'1 in the phase mixture.

The 29Si MAS NMR spectrum of pure cancrinite (Figure 5.4) formed from a
reaction composition of Si:Al:NaOHzH2O=1:1:18.8:125 exhibits a single sharp
signal at -87.3 ppm, in agreement with the alternating ordering of tetrahedrally
coordinated Si, Al atoms in the CAN framework for a Si/AI ratio of 1.0. Some

studies”19 have demonstrated that 29Si MAS NMR is not suitable for determining

127

 

whether an aluminosilicate sample is sodalite or cancrinite because both phases
have the same or very similar chemical shifts. Therefore, 29Si MAS NMR is not

helpful for distinguishing cancrinite from a cancrinite-sodalite phase mixture.

298i

 

 

 

-40 -60 -80 -100 -120 -140 -160
PP")

Figure 5.4 29Si MAS NMR spectrum of pure cancrinite precipitated from the

educts with composition of Si:Al:NaOH2H2O=1 :1 :18.8:125 at 90°C after 3 days.

N2 adsorption-desorption isotherms for products synthesized at different
alkalinities are presented in Figure 5.5. A fairly narrow hysteresis loop for pure
cancrinite formed at x=37.5 and 18.8 implies a mesoporous structure for these
products. However, it should be noted that the pore sizes shown in Figure 5.5 (a)
are the result of secondary pores between aggregated cancrinite crystals rather
than the pores in the cancrinite framework structure. The micropores (5.9><5.9 A)

of cancrinite are blocked by nitrate anions, thus preventing the N2 molecules from

128

Figure 5.5 N2 adsorption-desorption isotherms of (a) pure cancrinite products
and (b) CAN/SOD phase mixtures, formed at different alkalinities after 1 day or 3
days of reaction period. The insert provides BJH pore size distributions
determined from the adsorption branches of the N2 isotherms of selected

products.

129

(a)

(b)

350 —

300

250 ~

200 2

150 71

100 —

N2 adsorbed volume (cc/g. STP)

50~

 

350

300 2

N
0|
0

150

N2 adsorbed volume (cc/g, STP)

0'!
O

200 -

100 ~

0.035
0.03
D0.025
2 0.02
%0015
0.01
0.005
0 . . . .

010 20 30 40 50 60 70 80
Pore Diameter (nm)

 

 

  
         

 

 

T 1 T

+ CAN 1-1-18.8-125 30
+ CAN 1-1-37.5—125 3D
+ CAN 1-1-75—125 3D

 

-cn -o_-----—----——----—l

0.2 0.4 0.6 0.8 1

TI

 

l

    

 

l T

0 102030405060 7080
Pore Diameter (nm)

+ CAN/SOD 1-1-4.69—125 1D
-*- CAN/SOD 1-1-4.69-125 3D
~9- CAN/SOD 1-1-9.38-125 3D

I

  
  
 
 

 

_ -
..... _ _ --
- - - . _ _.
. — - ‘ - ' - - — t- " " - n - ‘-
if‘iz'é'é'i'31'z':—:':': : : : : - _ .. - ~ - -
O'-

 

—.
i Tf 1”’T~T‘T‘j’mgr_’r__r—T_ T V T I ‘1 l T I T I T Tfi

0 0.2 0.4 0.6 0.8 1
P/Po

130

 

accessing the interior of the pores. This explains the very limited N2 uptake in the
adsorption branch at low relative pressure (P/Po < 0.1) for cancrinite. With regard
to the SOD/CAN phase mixture prepared at x=4.69 after 1 day, the nitrogen
shows a broad hysteresis loop and the pore filling step begins at about 0.45 P/Po.
Moreover, the hysteresis loop shape of this cancrinite-sodalite mixture transforms
to that of the pure cancrinite when the reaction period was increased to 3 days,

suggesting that sodalite-to-cancrinite phase conversion may occur.

‘~ C101 C211

1
l (3110

  
 
 
 
 
 
  
 
 

CAN 1-1-18.8-125@90°C 30

 

 

 

CAN 1 -1-18.8-125@60°C 3D

Intensity (a.u.)

CAN 1-1-18.8-125@40°C 3D

 

SOD/CAN 1-1-18.8-125@25°C SD

 

 

1 0 20 30 50 60 70

40
29 (degree)

Figure 5.6 XRD patterns of products formed at different reaction temperatures.

(S=sodalite, C=cancrinite)

131

 

We also studied the influence of reaction temperature on the synthesis of
cancrinite and sodalite. The wide angle hkl reflections of reaction products
formed at 90, 60, 40 and 25°C are provided in Figure 5.6. All of the products
were formed from the same initial reaction composition of Si:Al:NaOH2H2O at
1:1:18.8:125 after 3-day period. As illustrated by XRD patterns, pure cancrinite
can be synthesized at temperatures as low as 40°C. In addition, the lower
temperatures may be in favor of the formation of smaller cancrinite crystallites, as
indicated by the broader reflection peaks for CAN@60°C and 40°C in comparison
to CAN@90°C. Sodalite crystals emerge together with cancrinite at room
temperature, but the cancrinite is a minor co-product, as indicated by the very
weak intensities of the CAN 101 and 211 reflections. TEM images in Figure 5.7
clearly show the cotton-ball like aggregation of nanometric sodalite crystals with
dimension less than 500 nm. Actually, no cancrinite crystals with a characteristic
rod-like morphology are observed during the TEM investigation due to the minor

amount.

   

Figure 5.7 TEM images of sodalite crystals prepared at 25°C after 3 days at high

and low magnifications.

132

 

 

 

CAN °C SD

 

CAN@60°C so
___—_J—x

Absorbance (a.u.)

CAN@40°C 3D
A

 

SOD/CAN@25°C so

 

 

 

 

F _l l w_—'T——.—_ T T r

3900 3400 2900 2400 1900 1400 900 400
Wavenumber (cm")

Figure 5.8 FTIR spectra of products as a result of changing reaction temperature.

Pure cancrinite framework structures formed at 40°C or above also can be
verified by FT-IR results. A combination of fingerprint peaks in the region of 500-
800 cm'1 from both sodalite and cancrinite can be resolved for SOD/CAN@25°C
(Figure 5.8). Also, the strong absorption band at ~1000 cm"1 attributed to vas
vibration for sodalite framework together with the shoulder peak at 1120 cm‘1
from v... vibration for cancrinite framework further suggests that the major phase
should be sodalite in this mixture. According to N2 adsorption-desorption
measurements, this SOD/CAN phase mixture exhibits a large BET surface area

of 209 m2/g, a uniform BJH adsorption pore size distribution centered at 3.0 nm

133

 

and a total pore volume of 0.29 cm3/g. The nitrogen isotherm of SOD/CAN is

easily distinguished from those of pure CAN, as shown in Figure 5.9.

0.09
500 7~ 0.08
0.07
0.06
20.05
%0.04 ~
0.03
0.02
0.01

0

4:- r:-
o 01
o o

D

OJ
01
O

   

 

0.)
O
O

0

10 20
Pore Diameter (nm)
250 7

+CAN 1-1-18.8-125@900 so
200 ‘ +CAN1-1-18.8-125@6003D
150 j +CAN1-1-18.8-125@4003D

N2 Volume Adsorbed (cm3/g, STP)
8
O

 

     

 

I T

0 0.2 0.4 P/Po 0.6 0.8 1

"1-7 7' " ' .' —T-"'.

Figure 5.9 N2 adsorption-desorption isotherms of samples formed as a function
of reaction temperature after 3 days reaction time. The insert provides the BJH
pore-size distribution determined from the adsorption branch of the N2 isotherm

of SOD/CAN@25°C phase mixture.

5.4 Summary
Nitrate cancrinite and sodalite have been synthesized based on a simple

batch reaction method using sodium silicate and aluminum nitrate as silica and

134

alumina sources. The reaction conditions, 9.9. amount of base, reaction
temperature and reaction time, were varied to determine their influence on the
zeolite phase formation. Single-phase nitrate cancrinite can be formed in a wide
range of synthesis parameters. At high concentrations of NaOH corresponding to
OH'lSi ratios of 18.8 to 75, cancrinite emerged at 90°C after 24 h. The
morphology and crystallite size of the resulting products are very sensitive to the
alkalinity, according to the BET surface area and TEM measurements. At low
concentrations of NaOH corresponding to OH'lSi ratios of 4.69 to 9.38, a
cancrinite-sodalite mixture was obtained, as confirmed by the powder X-ray
diffraction and FTIR spectroscopy characterizations. Moreover, a phase
transformation from sodalite to cancrinite occurred in this mixture with increasing
reaction time. This observation suggests that sodalite is the kinetically favored
phase under reaction conditions, but cancrinite predominates as the
thermodynamically stable phase. With respect to the dependence on the
temperature, pure cancrinite was found over the range 40-90°C, whereas a
reaction temperature below 40°C results in the co-crystallization of sodalite as

the major phase.

135

 

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137

 

CHAPTER 6

Epoxy-Zeolite Composites with Enhanced Tensile Preperties

6.1 Introduction

The chemical stability, excellent adhesion properties, and optical
transparency of thermoset epoxy polymers make them suitable for a variety of
applications, such as coatings, adhesives and electronic insulators. So far, most
of the epoxy-based composites reported are reinforced by layered silicate clay
nanoparticles”, mesostructured silicas‘”13 and carbon nanotubes (CNTs)“*18 in
order to improve mechanical, thermal, and barrier properties. Only a few studies

‘9'” wherein the focus was on

have been reported on zeolite-epoxy composites,
the effect of natural zeolite particles on the curing kinetics and mechanical
properties of the epoxy system. With increasing zeolite content, tensile strength
decreased dramatically, while Young's modulus increased, which is the result
typical of conventional composites.19

In Chapter 5, we described the low-temperature synthesis of nitrate
cancrinite (denoted as CAN) and nitrate sodalite (denoted as SOD) zeolites. In
the present study, we investigate the use of CAN and SOD zeolites as
reinforcement agents for thermoset epoxy polymer, particularly for a glassy

epoxy system. The large one-dimensional 12-ring channel (5.9 AXS.9 A) and the

small particle size (<100 nm) of rod-like cancrinite, and the high surface area

138

(~200 mzlg) and a large average BJH pore size (~3.0 nm) of cotton-ball like
sodalite are shown to substantially improve the tensile properties of the polymer.
Though a few studies have been reported concerning applications of

cancrinite or cancrinite-type zeolites?"28

including as antacid?"26 heavy-metal
trapping agents,” and cholesterol-controlling drug,28 no efforts have been made
on the polymer reinforcement application. The composites formed from synthetic
nitrate cancrinite and sodalite represent the first example wherein the modulus,

strength and toughness of epoxy-zeolite composites all are improved.

6.2 Experimental
6.2.1 Materials

Sodium silicate containing 27 wt.% SiO2 and 14 wt.% NaOH and
AI(NO3)'9H20 were purchased from Aldrich Chemical Co.. Mg(NOa)2-6H2O was
purchased from Columbus Chemical Industries, Inc.. The epoxide resin used to
form glassy epoxy matrices was a diglycidyl ether of bisphenol A (DGEBA), more
specifically, EPON 826 with an epoxy equivalent weight of ~182 glequivalent.
The curing agent was a polyoxypropylene diamine of the type
NH2CH(CH3)CH2[OCH2CH(CH3)]XNH2 (x = 2.6), which was provided by
Huntsman Chemicals under the tradename Jeffamine D230. All other chemicals
were purchased from Aldrich Chemical Co. and used without further purification.
6.2.2 Cancrinite and Sodalite Synthesis

In a typical synthesis, 8.55 g of sodium silicate solution (27 wt.% silica, 14
wt.% NaOH) (0.03850 mol silica) and 14.44 9 (0.03850 mol) of Al(NOa)3-9H2O

139

were first added to 87 ml of deionized water in a sealed polypropylene bottle. A
cloudy gel-like mixture was obtained after complete mixing. Then, 28.8 g (0.72
mol) or 57.6 g (1.44 mol) of NaOH was introduced and the bottle was vigorously
shaken to ensure homogeneous mixing. The molar ratio for sodium silicate,
aluminum nitrate, sodium hydroxide (not including NaOH in sodium silicate
solution) and water is 1:1:18.8/37.5:125. The resulting gels were heated in oven
at 90°C or were allowed to age at room temperature in tightly stoppered
polypropylene bottles for 3 days. All products were washed using deionized
water via centrifugation 3 times. All samples were dried overnight in oven at 80°C
under nitrogen gas flow.
6.2.3 Preparation of Epoxy-Zeolite Composites

A predetermined amount of cancrinite or sodalite was added to the epoxy
resin and mixed at 50°C for 10 min. The amount of Jeffamine D230 curing agent
needed to achieve an overall NH: epoxide stoichiometry of 1:1 was then added
and the mixture was stirred with a stir bar at room temperature for 15 min,
outgassed under vacuum for 20 min at room temperature, and transferred to a
silicone mold. The composites were partially cured at 75°C for 3 h and then fully
cured at 125°C for 3 h under N2 flow.
6.2.4 Characterization

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 2008
diffractometer equipped with Cu Ko X-ray radiation and a curved crystal graphite
monochromator, operating at 45 kV and 100 mA. Rectangular flat specimens of

cross-linked composites were mounted onto an aluminum holder for XRD

140

examination. N2 adsorption-desorption isotherms were recorded at —196°C on a
Micrometritcs Tristar 3000 sorptometer. Prior to analysis samples were
outgassed at 150°C and 1045 Torr for a minimum of 12 h. BET surface areas
were calculated from the linear part of the BET plot and BJH pore size
distribution were obtained from adsorption isotherms.

Tensile properties of dog-bone shaped specimens were tested according
to ASTM procedure D3039 using an SFM-20 United Testing System equipped
with a laser extensometer. The measurements were conducted at room
temperature with a crosshead speed of 0.5 mm/min for glassy specimens. The
reported tensile values are the averages for 4-5 specimens with dog-bone

shapes.

6.3 Results and Discussion

Detailed characterization of nitrate cancrinite and nitrate sodalite is
described in Chapter 5. The synthetic conditions together with the textural
properties of the selected zeolites used for reinforcements of glassy epoxy
polymers are summarized in Table 6.1. Note that SOD-18.8-25°C is actually a
mixed SOD-CAN phase with minor amount of nitrate cancrinite. X-ray powder
patterns and N2 adsorption-desorption isotherms of selected zeolites are

provided in Figure 6.1 and Figure 6.2, respectively.

141

Table 6.1 Textural properties of synthetic zeolites

 

 

Si:Al:NaOHzH2O BET BJH
Sodium = 1:1:xz125 in surface Totall pore adsorption
aluminosilicatea the starting area V0 "3219 pore size
solutions (mzlg) (cm 9) (nm)
CAN-37.5-90°C =37.5 41 .9 0.41 -
CAN-18.8-90°C x=18.8 74.8 0.49 -
SOD-18.8-25°C x=18.8 208.6 0.29 3.0

 

a All zeolite products are aged for 3 days.

 

 

 

 

 

 

 

    

 

 

1 C300
1 6101 c211
1 C110
C210 CAN-37.5-90°C
E7
8
E
U CAN-18.8-90°C
S211
S110
Aw SOD-18.8-25°C
10 20 30 50 60 70

40
29 (degree)

Figure 6.1 XRD patterns of synthetic zeolites after a reaction time of 3 days.

142

350- ---------------- ...... . ..... .............

_._ CAN-18.8-90C
300 - -.— CAN-37.5900
+ SOD-18.8—25C

      

.........................

N2 adsorbed volume (cc/g. STP)

 

 

 

P/Po

Figure 6.2 N2 adsorption-desorption isotherms of synthetic zeolites after a

reaction time of 3 days.

All of the selected zeolites exhibited visually very good dispersion in epoxy
prepolymers. The distribution of sodium aluminosilicates was qualitatively judged
by examining the relative intensities of zeolite diffraction peaks obtained from the
lower and upper surfaces of the cured composite specimens. Little or no settling
of aluminosilicate occurred during the polymer curing process because the
intensities of the Bragg reflection peaks of zeolite particles from the top and

bottom surfaces of the composite were equivalent.

143

Table 6.2 Tensile properties of glassy epoxy composites reinforced by nitrate

cancrinite and sodalite

 

 

. . Tensile Tensile .
3.3333123... 1.313239 strength modu'us 3332115}; 1‘"
Wm? (Gpa)

Pristine Epoxy 0 82.1 3.12 5.14

CAN-37.5-90°C 2 85.9 3.19 5.16

5 85.9 3.31 4.63

10 94.3 3.65 5.00

CAN-18.8—90°C 2 90.4 3.29 6.39

5 89.1 3.95 5.68

10 94.0 3.85 5.25

SOD-18.8-25°C 5 94.9 4.19 6.95

10 90.3 4.52 5.35

 

a The standard deviations for all values are <6%, except for the values of elongation-at-break in
which case the standard deviation is <20%.

The tensile properties of glassy epoxy composites prepared from CAN
and SOD zeolites and a comparison of representative stress-strain curves are
provided in Table 6.2 and Figure 6.1, respectively. When compared to the
pristine epoxy, all of the epoxy composites reinforced by cancrinite and sodalite
demonstrate significant improvement in Young's modulus and tensile strength at
all loadings. Moreover, elongation-at-break is also enhanced in most of the cases
except for the cases of CAN-37.5-90°C at 5 wt% and 10 wt% loadings. For these
composites, reduced elongation-at-break may be caused by the local defects
(e.g. air bubbles) in the composites. The increase in the elongation-at-break
additionally results in enhanced toughness.

The dramatic influence of cancrinite and sodalite on the tensile properties
of the glassy polymer is obviously evidenced by the stress-strain curves (Figure

6.3). The pristine epoxy showed a failure behavior typical of a brittle material.

144

100

Stress (Mpa)
O) ‘1 co to
O O O O O

#01
O

302

20
10

100
90
80
70

Stress (Mpa)
_s N 0) «F U" C)
O O O O O O

O

 

 

 

 

(a)

- 10%
i.-

P 5% 2%
¢ / 0%
'L ,
!
L
ir—

L.__ L l 1 1 _.
0 2 4 6 8 1 0
Strain (%)

(b)
1‘ oo o
, . 0 / 10 A. 5%
f
._ /‘
.
0 4. 8
Strain (%)

Figure 6.3 Representative stress-strain curves for glassy epoxy nancomposites

reinforced by (a) CAN-18.8-90°C and (b) SOD-18.8—25°C at the wt% zeolite

loadings indicated.

145

The stress increases with increasing load to a maximum value, after which the
specimen ruptured with little yield. Similar behavior is observed for the
composites, except that the composites exhibit substantial yield and the initial
slopes are larger and the maximum stress is higher due to increased modulus
and strength, respectively. Elastic modulus of the nanocomposites increases with
increasing zeolite concentration, while stress and strain-at-break suffer from a
loss in some cases. For instance, a 15.6% increase in stress and 35.2% increase
in elongtation-at-break were observed at 5 wt% of SOD-18.75-25°C loading,
while at 10 wt%, only a 10.0% increase in stress and 4.1% increase in
elongation-at-break were achieved. It’s been noticed that with increasing fillers
content, especially for those fillers having good adhesion with epoxy monomers,
the viscosity of epoxy precursors goes up rapidly, which may lead to the difficulty
of degassing the prepolymers and associated air-bubble defects in the dog-bone
samples. This may explain the loss of strength and elongation-at-break at higher
loadings.

The substantial improvements in the tensile strength, stiffness and
toughness achieved for the epoxy-zeolite composites in comparison to the
pristine polymer most likely are a consequence of strong interfacial interaction
between the epoxy polymer chains and microporous zeolites and the good
dispersion of zeolite particles throughout the matrix.

It had been previously reported that poly(ethylene oxide) (PEO),
poly(acrylic acid) (PA), dextran and proteins were found to be adsorbed on a

dealuminated FAU-type zeolite.29 The high specific interaction suggested the

146

penetration of zeolite pores by macromolecules to some extent. Recently, Macko
et al.”33 reported for the first time that the highly hydrophobic polyethylene (PE)
and isotactic polypropylene (PP) may be adsorbed from solutions onto certain
zeolites with pore/channel diameters ranging from ~5 to 6 A. These authors
conclude that both hydrophobic polyolefins likely penetrate and are at least
partially included in the interiors of the narrow zeolite nanopores. Moreover,
polymers may be synthesized within the zeolite pores, but this occurs after pre-
intercalating the pores with the monomers, and then the polymer is fixed as an
integral part of the zeolite lattice.34 The cancrinite structure consists of a main
one-dimensional channel system along the [001] axis and is formed from elliptical
12-rings of tetrahedra giving a pore dimension of 5.QX5.9 A, which is big enough
for the diffusion of epoxy monomers. However, there are invariably nitrate anions
situated inside the channels which might impede efficient penetration of epoxy
resin and curing agents from one side to another. Also, the entrance to the 8-
cages of cancrinite is too small to allow entry of epoxy prepolymers. Hence, the
cross-linked epoxy chains after curing process might not be included in the
interiors of the cancrinite pores. Thus, the polymer most likely is adsorbed onto
external surfaces of the cancrinite crystal. The increased Young’s modulus is due
to the cancrinite crystallites being stiffer than the epoxy matrix and the improved
tensile strength can be ascribed to the strong chemical interaction between the
hydroxyl groups generated during the cure reaction of epoxy monomers and the

sodium cations at cancrinite external surface sites.

147

The magnitude of tensile property improvement achieved by CAN-18.75-
90°C is better than that of CAN-37.5-90°C and this can be rationalized by the fact
that CAN-18.8-90°C has a smaller particle size and associated higher BET
surface area (Table 6.1) compared to CAN-37.5—90°C, thereby resulting in better
dispersion and interfacial interaction with the epoxy matrix.

The pore opening In the sodalite framework structure is only 2.4xz.4 A
and is too small to expect intercalation of the epoxy polymer within the
nanopores. However, the high surface area (~200 mzlg) and a large average
BJH pore size (~3.0 nm) of sodalite with a cotton-ball like morphology allows the
polymer to fill the secondary pores between sodalite nanoparticles and thus
eliminate the tendency toward microcrack formation between the two phases,
leading to improved tensile properties. Furthermore, the reinforcing benefits
provided by sodalite are superior to those provided by cancrinite at equivalent

loadings.

 

Figure 6.4 Optical image comparing the transparency of (a) pristine epoxy
polymer and epoxy composites prepared from (b) CAN-18.8-90°C, (c) CAN-37.5-

90°C and (d) SOD-18.8-25°C. All of the composites are at 5 wt% zeolite loading.

148

Another significant advantage of synthetic cancrinite as a reinforcing agent
is the optical transparency of the resulting composites (cf. Figure 6.4). The
transparency of the epoxy-cancrinite nanocomposite is comparable to that of the
pure polymer. On the other hand, the composite made from sodalite is opaque.
Thus, in comparison to sodalite, the refraction index of the cancrinite zeolite more
nearly matches that of the epoxy matrix.

In order to evaluate the advantages of synthetic nitrate cancrinite and
nitrate sodalite as polymer reinforcing agents, the observed mechanical
properties of epoxy-zeolite nanocomposites are compared with those of epoxy
nanocomposites made from organoclay, mesoporous silicas, and carbon
nanotubes. Organoclays have been successfully used for reinforcing rubbery
epoxy matrices”, which are relatively soft and easy to reinforce. However, the
reinforcement provided by organoclays became less significant for rigid glassy
epoxy matrices. The strength of glassy epoxy nanocomposites was not improved
by the addition of organoclays.35 Even when uniform dispersion of exfoliated or
intercalated clay nanolayers is achieved in epoxy, the tensile strength is
decreased by 25% at 5 wt% of organoclay loading, although the modulus is
improved by 40%.35 The plasticizing effect of the clay organic modifiers may
compromise the interfacial interactions needed for enhanced strength.

Mesoporous silicas (denoted MSU-J-TEOS and MSU-J-SS) with
wormhole framework structure have been proven to be effective reinforcement
agents for glassy epoxy system.11 The larger pore mesophases (MSU-J-TEOS)

provide better reinforcement at 5 wt% loading, than the corresponding smaller

149

pore analogs (MSU-J-SS), most likely, because of more efficient polymer
impregnation of the larger mesopores. A 18.2% increase in strength and a 31.3%
increase in modulus were observed at 5 wt% loading of MSU-J-TEOS, in
comparison to the pristine glassy epoxy. However, elongation-at-break is
decreased by 15.6%. The improvements in the tensile strength and modulus
provided by MSU-J-TEOS silica is comparable with those achieved by synthetic
zeolites of this investigation at the same loading. Epoxy-zeolite composites made
from nitrate sodalite exhibit a 15.6% and 34.3% increase in strength and modulus,
respectively, at 5 wt% loading. Additionally, a 35.2% increase in elongation-at-
break is realized, resulting in enhanced toughness. Therefore, microporous
zeolites are better toughening agents for glassy epoxy polymer.

When carbon nanotubes are considered for reinforcement applications, it
is important to address bundle formation arising from intertubular van der Waals
forces.36 Normally, surface functionalization of the nanotubes is necessary in
order to facilitate the dispersion of CNTs in the polymer matrix.” 18 In one study,
a fully integrated nanotube-glassy epoxy polymer composite with covalent
chemical bonding between the matrix and functionalized SWNTs is reported.16 At
a loading of only 1 wt% functionalized SWNTs, the resulting glassy epoxy
composites demonstrated significant 25.3%, 30.8% and 30.8% increases in
tensile strength, Young’s modulus and elongation-at-break, respectively. This
also suggests higher toughness with respect to the neat epoxy. For the
composites with higher loading of 4 wt%, up to 70% improvement in Young's

modulus was found, although no further increase of tensile strength was

150

observed. This might indicate an upper limit to the content of nanotubes that can

be added directly into epoxy resin because of a significant increase in viscosity.

6.4 Summary

Microporous nitrate cancrinite and nitrate sodalite zeolites have been used
to form glassy epoxy nanocomposites containing 2-10 wt% zeolite. In
comparison to the pristine epoxy polymer, epoxy composites reinforced by
cancrinite and sodalite demonstrate substantial improvement in the tensile
strength, modulus, elongation-at-break and toughness. For example, a 15.6%
increase in strength, a 34.3% increase in modulus and a 35.2% increase in strain
were observed at 5 wt% loading of SOD-18.8-25°C. The reinforcement benefits
are most likely a consequence of strong interfacial interactions between the
epoxy matrix and microporous zeolite. Moreover, the increases in tensile
properties occur with little or no sacrifice in optical transparency.

When compared with organoclay, mesoporous silica and carbon
nanotubes, microporous zeolites still exhibit advantages as reinforcing agents for
glassy epoxy system. In order to achieve homogeneous dispersion in an epoxy
matrix, both clay and carbon nanotubes need to be surface modified. Organically
modified clays can only improve the Young’s modulus, not the strength of rigid
glassy epoxy. Thus, nitrate CAN and SOD are better reinforcing agents for
glassy epoxy system than organoclays. Regarding functionalized CNTs, only a 1
wt% loading can reach the reinforcement benefits provided by synthetic zeolites

at a 5 wt% loading. However, increasing the content of nanotubes will decrease

151

rather than further increase the tensile strength owing to a significant increase in
viscosity. The reinforcements provided by mesoporous silicas with large average
pore size (~ 20 nm) are comparable with those achieved by synthetic zeolites at
equivalent loading. But epoxy-silica mesocomposites suffer from a loss in the
elongtation—at-break. Therefore, synthetic zeolites are better toughening agents
for glassy epoxy polymer. Also, it should be noted that the improvement in tensile

properties are realized without the need for organic surface modification.

152

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