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ORGANIC—MODIFIER-FREE PATHWAYS FOR THE PREPARATION OF
POLYMER - METAL OXIDE NANOCOMPOSITES

By

Siqi Xue

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

2007

ABSTRACT

ORGANIC-MODIFIER-FREE PATHWAYS FOR THE PREPARATION OF
POLYMER — METAL OXIDE NANOCOMPOSITES

By
Siqi Xue

Novel preparation strategies for the formation of polymer - metal oxide
nanocomposites have been developed to eliminate the need for surface organic-
modifiers and thereby to avoid the disadvantages of the modifiers, such as their
plasticizer effects and low thermal stabilities. The success of these strategies
relies on the smart design of the synthetic or naturally occurring metal oxides and
the composite preparation methodology.

Synthetic clay materials exhibit high purity, uniform surface properties and
composition, and are promising substitutes for natural clay minerals with more
variable properties. We synthesized three types of saponite-Iike silicates with
different octahedral sheet compositions and different layer stacking orders
depending on synthesis methodology. When bis(triethoxylsilyl) methane was
used as the Si source, an inorganic-organic hybrid clay material, which had -CH2-
tetrahedra bridging groups in the basal plane, was achieved.

Among the synthetic saponites, an irregularly stacked derivative (denoted
SAP) was examined as an epoxy polymer reinforcing agent. SAP exhibited a
large surface area (920 mzlg), small lateral dimension (20 - 30 nm), irregularly

stacked layered, morphology with porous surfaces, which allowed polymer

intercalation without any surface modifications. A Uniform dispersion of SAP
nano-aggregates was achieved for the epoxy - SAP nanocomposites, providing
improved strength, modulus and toughness of the polymer. For example, for a
glassy epoxy system, a 10% by weight loading of SAP provided a 10% increase
in strength, 30% increase in modulus and 45% increase in toughness, compared
with the pristine polymer. In addition, the nanocomposites had the same glass
transition temperature as the pristine epoxy, because the plasticizer effects of the
organic modifiers were avoided.

Palygorskite, a silicate clay with a pleated 2:1 layered structure, has a
lath-like particle morphology, and low surface charges that make it an attractive
candidate for the formation of polymer nanocomposites. The pristine clay mineral
provided reinforcement to epoxy matrix without organic modification. The
silylated derivatives of palygorskite provided better dispersions in rubbery epoxy
matricx than the pristine mineral, affording further improvements in mechanical
properties, especially at low loadings levels of 2 and 5 % (w/w).

Mesostructured silicas with pore sizes larger than 20 nm allowed
polyethylene intercalation in the porous regions without silica surface
modifications. Even with an inhomogeneous dispersion of mesoporous silica, the
polymer mesocomposites exhibited improvement in tensile strength and modulus
that were analogous with those achieved in polyethylene — clay systems where

extensive modifications were required.

Copyright by
SIQI XUE
2007

ACKNOWLEDGEMENT

I would first like to express my appreciation to Prof. Thomas J. Pinnavaia
for his thoughtful guidance during the five years of my Ph.D. study. Prof.
Pinnavaia h as been my role model of not only a n e rudite scientist, b ut also a
great person. I feel extremely lucky to be his student.

I would also like to thank former and current group members. Discussions
with Dr. Kim, Dr. Liu, Dr. Reinholdt and Dr. Zhang have always been very
inspiring a nd helpful whenever I met problems during research. D r. Park had
great patience in teaching me various techniques to get started in research when
I just joined the group. My friendship with Emily, Hua, Jainisha and Xin extended
outside the lab, and we went for movies, dinners and shopping together.
Christian, D.K., Hui, Joel are all great person to work with. My close friend
Shujuan has been like my sister. I am also very grateful for the staff members in
the Composite Materials and Structure Center and Microcopy Center at Michigan
State University for teaching me how to use their facilities.

I would like to dedicate this humble thesis to my parents, Mr. Ligen Xue
and Mrs. Xue, Xiuhua Shi. I could not have completed my Ph.D. degree without

their love and support.

TABLE OF CONTENTS

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

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

CHAPTER 1

AN OVERVIEW OF POLYMER — CLAY NANOCOMPOSITES
1.1 Introduction .......................................................................................... 1
1.2 Clay Surface Modification ........................................................... 7
1.3 PCN Preparation Technology .................................................... 13
1.4 Characterization of Clay Nanolayer Dispersion ............................. 20
1.5 Nanolayer Orientation .............................................................. 28
1.6 Synthetic Materials for Polymer Reinforcement ............................. 30
1.7 Conclusion ............................................................................ 41
References ................................................................................. 43

CHAPTER 2

SYNTHESIS OF SAPONITE-LIKE PLATEY SILICATES AND THEIR
INORGANIC-ORGANIC HYBRID DERIVATIVES

2.1 Introduction ............................................................................ 51
2.2 Experimental

2.2.1, M aterials ................................................................... 55

2.2.2 Saponite synthesis ...................................................... 55

2.2.3 Ion exchange for Mg-SAP ............................................. 56

2.2.4 Hybrid saponite synthesis ............................................. 57

2.2.5 Characterization ......................................................... 58
2.3 Results

2.3.1 Inorganic saponite derivatives ........................................ 59

2.3.2 Hybrid organo-saponite derivatives ................................. 71
2.4 Discussion

2.4.1 Inorganic saponite derivatives ........................................ 78

2.4.2 Hybrid organo-saponite derivatives ................................. 79
2.5 Conclusion ............................................................................ 81
References ................................................................................. 82

CHAPTER 3

Porous Synthetic Smectic Clay for the Reinforcement of Epoxy Polymers

vi

3.1 Introduction ........................................................................... 84
3.2 Experimental

CHAPTER 4

3.2.1 M aterials ................................................................... 86
3.2.2 Saponite synthesis ....................................................... 86
3.2.3 Composite preparation .................................................. 87
3.2.4 C haracterization .......................................................... 88
3.3 Results
3.3.1 Epoxy — saponite nanocomposites ....................................... 89
3.3.2 Composite Properties ................................................... 91
3.4 Discussion ............................................................................. 96
3.5 Conclusion ............................................................................. 98
References ................................................................................. 99
PALYGORSKITE AS AN EPOXY POLYMER REINFORCING AGENT
4.1 Introduction ......................................................................... 101
4.2 Experimental
4.2.1 M aterials ................................................................. 102
4.2.2 Palygorskite purification .............................................. 103
4.2.3 Palygorskite silylation ................................................. 104
4.2.4 Composites Preparation .............................................. 104
4.2.5 Characterization Methods ............................................ 105

4.3 Results
4.3.1 Characterization and Purification of Palygorskite..............106

4.3.2 Silylated Palygorskite .................................................. 110
4.3.3 Epoxy Nanocomposites ............................................... 111
4.4 Discussion ........................................................................... 114
4.5 Conclusion ........................................................................... 120
References ................................................................................ 122
CHAPTER 5
FOAM-STRCTURED MESOPOROUS SILICAS FOR POLYETHYLENE
REINFORCEMENT
5.1 Introduction .......................................................................... 124
5.2 Experimental
5.2.1 Materials .................................................................. 127
5.2.2 Mesoporous silica synthesis ......................................... 127
5.2.3 Mesocomposites preparation ........................................ 128
5.2.4 C haracterization ........................................................ 128
5.3 Results and Discussion ........................................................... 130
5.4 Conclusion ........................................................................... 139

vii

References ................................................................................ 140

CHAPTER 6
CONCLUSIONS AND FUTURE DIRECTIONS
6.1 Comparison of Various Polymer Reinforcing Agents

6.1.1 Nanoparticles for Epoxy Reinforcement ......................... 142
6.1.2 Nanoparticles for Polyethylene Reinforcement ................ 145
6.2 Future Directions .................................................................. 147
References .............................................................................. 1 51

viii

LIST OF TABLES

Table 1.1 Thermal stability date for imidazolium- and alkylammonium-MMT.......9

Table 2.1 The BET surface areas and pore volumes of Mg-SAP, Zn-SAP, and
SAP-200 and the oazt, Cs“-, and K"- exchanged forms of Mg-SAP ................ 61

Table 2 .2 Elemental atomic ratios for Mg-SAP a nd its exchanged derivatives
obtained from quantitative SEM-EDS analysis ............................................ 71

Table 3.1 Tensile properties of glassy epoxy composites reinforced by SAP and
NaMMT clay nanoparticles ..................................................................... 92

Table 3 .2 D ynamic mechanical analysis (40 °C) of glassy epoxy composites
filled with SAP and ClG-SAP clay ............................................................. 94

Table 3.3 Tensile properties of rubbery epoxy composites reinforced by SAP clay
nanoparticles ........................................................................................ 95

Table 4.1 The tensile properties of rubbery palygorskite-epoxy composites.
Numbers in the parentheses are relative standard deviations ....................... 115

Table 4.2 Tensile properties of glassy epoxy reinforced by silylated palygorskite
HMSZ-PLG. Numbers in the parentheses are relative standard deviations......117

Table 5.1 The extrusion and injection molding conditions for pure polymers and

their m esocomposites .......................................................................... 128
Table 5.2 The surface properties of MSU-F and MSU-F2 ............................ 133
Table 5.3 The crystallinity of HDPE and LDPE mesocomposites .................. 136
Table 5.4 The mechanical properties and oxygen perrneabilities of LDPE - silica
mesocomposites ................................................................................. 1 38
Table 5.5 The mechanical properties of HDPE - silica mesocomposites. .......139

Table 6.1 Comparison of various inorganic particles for PE reinforcement......146

ix

LIST OF FIGURES

Figure 1.1 The number of publications on “polymer - clay/layered silicate
composites” from the year 1990 to 2006 ..................................................... 2

Figure 1.2 Structure of 2:1 phyllosilicates .................................................... 3

Figure 1.3 Nanoscale arrangements of layered silicates in polymer
nanocomposites ..................................................................................... 6

Figure 1.4 Cryo-TEM image of styrene-butyl acrylate particles coated by
monofunctional silane-treated Laponite clay nanolayers, as obtained by emulsion
polymerization in the presence of 10 wt% Laponite clay platelets ................... 13

Figure 1.5 Schematic illustration of the intercalated state the exfoliation process
showing the forces acting on a pair of clay layers: (a) organically modified clay; (b)
epoxy-intercalated state; (c) forces acting on a two-particle clay tactoids ......... 17

Figure 1.6 TEM images of an epoxy nanocomposite (2.5 wt% clay) prepared by
slurry compounding ............................................................................... 19

Figure 1.7 (a) WAXD patterns and (b) TEM images of three different types of
PCNs ................................................................................................. 21

Figure 1.8 A phase contrast AFM image of an epoxy/synthetic fluoromica
nanocomposite .................................................................................... 24

Figure 1.9 (a) FTIR spectrum of hectorite powder and delaminated hectorite in
water; (b) a calibration curve of clay in-plane Si-O bandwidth vs. percent activator
added and percent dispersion ................................................................. 27

Figure 1 .10 z-axis polarized FTIR spectra of a LDPE film containing oriented
BentoneTM clay layers, with different tilt angle. Inset: schematic representation of
clay in-plane and out-of-pane Si—O bonds .................................................. 28

Figure 1.11 Hermans orientation parameter, Sd (squares), and Iayerspacing
(circles) determined from in-situ scattering (3 °C/min). Note that the intensity of
the basal reflection (d001=36A) rapidly decreases above 120 °C, effectively

disappearing by 180 °C. Orientation parameter above 120 °C was estimated from
the featureless scattering around the beam stop30

Figure 1.12 XRD patterns of (a) natural Na-hectorite, herein the peaks labeled
by * are due to (Mg, Fe)Si03 and CaC03 impurities; (b) synthetic Li-hectorite, and
(c) synthetic PVP-hectorite. Typical clay hkl reflections are identified ............... 31

Figure 1.13 Proposed schematic structures of the functionalized clays. Two
possible structures for C 16-LMS include a single layer structure a nd a d ouble
layer structure which is shown in (A). The proposed structure of C16-LMAS (b)
with both tetrahedral and octahedral AI present in the octahedral layer. In C16-
SiOz-LMAS (C) additional silicate groups without alkyl chains are present and
tetrahedral Al is likely to occupy some of the sites in tetrahedral layer ............. 35

Figure 1.14 Elastic moduli of PS/clay nanocomposites using (a) C16-SiOz-LMAS,
(b) C16-LMAS, (c) C16-LMS, and (d) Cloisite 20A from melt rheology data
obtained at T=150 °C, and 6.3 wt% inorganic material .................................. 36

Figure 1.15 DSC traces of the organosilicate. Inset shows the phase transition
from solid to liquid upon heating ............................................................... 37

Figure 1.16 (a) Tensile modulus, (b) tensile strength, (c)strain-at-break, (d)
toughness, and (e) oxygen permeabilities epoxy composites containing different
loadings of as-made MSU-J (solid circles), and calcined MSU-J (open circles)
mesostructured silica ............................................................................. 40

Figure 2.1 Schemes illustrating the difference in tetrahedral sheet structure for (a)
an inorganic clay and (b) an organo hybrid clay made from BTESM ................ 54

Figure 2.2 Wide angle hkl X-ray diffractions of (a) pressed Mg-SAP, (b) as-made
Mg-SAP, (c) Zn-SAP and (d) SAP-200 ...................................................... 60

Figure 2.3 TEM images of the synthetic saponites: (a) Mg-SAP; (b) Zn-SAP and
(c) SAP-200 ......................................................................................... 62

Flgure 2.4 (a) Nitrogen adsorption-desorption curves and (b) BJH adsorption
pore size distributions (off set by 0.1 for clarification) of as-made Mg-SAP and its
Ca2*-, Cs*-, and K“- exchanged derivatives ................................................ 64

Figure 2.5 The FTIR spectrum of Mg-SAP ................................................. 66

xi

Figure 2.6 (a) 29Si MAS NMR and (b) 27Al MAS NMR spectra for Mg-SAP.......68

Figure 2.7 Wide angle hkl X-ray diffractions of the Ca2*-, Cs*-, and K’-
exchanged Mg-SAPs, compared with that of as-made Mg-SAP ...................... 70

Figure 2.8 XRD patterns of Mg-SAP, C1-SAP and C2-SAP reaction products..72
Figure 2.9 The nitrogen adsorption-desorption isotherms for C1-SAP ............. 73

Figure 2.10 The (a) 29s: MAS NMR, (b) 13c MAS NMR and (c) ”Al MAS NMR
spectra of C1-SAP ................................................................................ 76

Figure 2.11 FTIR spectrum of C1-SAP. The circled regions in the spectrum of
C1-SAP indicate the incorporation of methylene groups in C1-SAP structure....77

Figure 2.12 TEM images of C1-SAP showing its platy morphology ................. 77

Figure 3.1 The XRD patterns of the topside and bottom of an epoxy — SAP 5.4%
composite, both showing 060 clay diffraction peak of similar intensity .............. 90

Figure 3.2 A TEM image of thin-sections of a glassy epoxy — SAP 5.4 wt%
composite ........................................................................................... 90

Figure 3 .3 R epresentative stress - strain c urves for glassy e poxy - SAP clay
nanocomposites at the wt % clay loadings indicated .................................... 93

Figure 3.4 Pictures of (a) a pristine glassy epoxy polymer and (b) the
corresponding epoxy - SAP nanocomposite containing 5.4 wt% clay .............. 95

Figure 4.1 A schematic [100] projection of palygorskite structure.................107
Figure 4.2 The hkl X-ray reflections of as-received palygorskite and of the
edimented fractions obtained upon aging a 10% (w/w) aqueous slurry of the
mineral. Q and C mark diffraction peaks indicative of quartz and carbonate
mineral impurities, respectively .............................................................. 108
Figure 4.3 TEM images of as-received palygorskite ................................... 110
Figure 4.4 FTIR spectra (KBr) of as-received palygorskite and the HMSZ

silylation product. The inset shows an expansion of the C-H stretching region
between 2800-3000 cm’1 for HMSZ-palygorskite ....................................... 111

xii

Figure 4.5 XRD patterns of the bottom sides of epoxy nanocomposite specimens
prepared from as-received and silylated palygorskite .................................. 112

Figure 4.6 TEM image of a thin section of a glassy epoxy nanocomposite
containing 5% (w/w) of HMSZ-PLG ......................................................... 112

Figure 4.7 Representative strain-stress curves for rubbery epoxy composites
containing different loadings of silylated palygorskite HMSZ-PLG ................. 114

Figure 5.1 (a) The N2 adsorption-desorption isotherms of MSU-F and MSU-F2.
The curves were offset by 200 for clarity. (b) Cell size (solid lines) and window
size (dash lines) distributions of MSU-F and MSU-F2 obtained from the
adsorption and desorption isotherm curves, respectively. The distributions were
determined by applying BJH model ........................................................ 132

Figure 5.2 The TEM images of (a) MSU-F and (b) MSU-F2 ......................... 133

Figure 5.3 TEM images of thin-sections of a HDPE - MSU-F 4.2 wt%
mesocomposites .............................................................................. 1 34

Figure 5.4 PLM images of (a) LDPE and (b) a LDPE — MSU-F 5.5 wt% loaded
mesocomposite .................................................................................. 137

xiii

CHAPTER 1
AN OVERVIEW OF POLYMER — CLAY NANOCOMPOSITES

1.1 Introduction

When a polymer materials is reinforced by a particle that is nanometric in
at least one dimension, the resulting polymer composite usually exhibits
remarkable improvements in material properties relative to the pristine polymers
or conventional composites. Nanocomposites can be divided into three
categories depending on whether the filler has one, two, or three dimensions in
the nanometer range. Smectic clay minerals in exfoliated form are
representative of the first type of nanoparticle filler, since the silicate sheets are
1-nm thickness, but up to several microns in diameter.

Polymer - clay nanocomposites (PCNs) have been extensively studied
ever since Toyota researchers first reported the nylon - exfoliated montmorillonite
(MMT) nanocomposites with improved mechanical properties (Figure 1.1) [1-14].
What makes clay minerals promising fillers to complement the drawbacks of
conventional polymer materials is the unique combination of the properties of
clay, such as the high stiffness and modulus, thermal stabilities. and low gas
permeabilities. Additionally, clays are inexpensive and abundant natural
minerals. The unique properties of clay can be transferred into polymer

1

composites only when clay sheets are uniformly dispersed in polymer materials
on a nanometer length scale. The resulting nanocomposites usually exhibit
improved mechanical properties, better thermal stabilities, reduced gas
permeability and flammability at clay loadings below 10 wt%, and even below 5

wt%.

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Figure 1.1 The number of publications on “polymer - clay/layered silicate

composites” from the year 1990 to 2006.

The commonly used clays for the preparation of PCNs belong to the
smectite family, which possesses a layered structure where silicate sheets stack

due to the electrostatic interactions between negatively charged silicate sheets

and exchangeable gallery alkaline or alkaline earth cations (Figure 1.2). Simply
adding clays to polymers results in the formation of conventional composites (or
microcomposites) that have micrometer-sized clay tactoids dispersed in the
polymer phase. However, when clay surfaces are modified and made
compatible with polymers, the polymer chains can wet the clay basal surfaces,
penetrate into clay galleries and exfoliate the tactoids into nanolayers, forming

nanocomposites.

       

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Two types of dispersed clay structures have been recognized in PCN
formation. One is the intercalated state, where the clay gallery heights are
expanded to a substantial extend (usually 10 A or more) by the polymer, but the
parallel orientation of the silicate layer, which is characteristic of the initial
tactoids, remains in registry upon composites. Such composites generally are
referred to as “intercalated nanocomposites”, particularly if the clay phase in the
nanocomposite exhibits a 001 X-ray reflection indicative of layer stacking. Even
when Bragg scattering is not observed due to irregular or very large spacings
between nanolayers, transmission electron microscopy can reveal the parallel
orientation of the expanded clay galleries, in which case the term “intercalated
nanocomposite” may still be appropriate. The other state of clay particle
dispersion is the exfoliated state, in which the average separation between
nanolayers is very large, usually up to 100 A or more. An exfoliated clay
nanocomposite does not exhibit any X-ray evidence for layer stacking and,
ideally, the parallel registry of the nanolayers is lost upon nanocomposites
formation, as evidenced by transmission electron microscopy. That is, there are
no remnants of the initial tactoid structure of the clay. The distinction between
intercalated and exfoliated states is difficult to establish, much less quantify,
particularly if the composite provides no X-ray diffraction evidence for layer
stacking. In many PCN systems that do not exhibit Bragg scattering for clay, a
combination of intercalated and exfoliated clay structures usually are present.

4

Vaia has suggested an expanded classification to describe PCN
morphologies on the basis of (1) relative changes in layer spacing and correlation
(d and 6d, respectively); (2) the relative volume fraction of single layers and layer
stacks; (3) the dependence of single-layer separation on silicate volume fraction
(rp) above a critical volume fraction for ordering (<p*) (Figure 1.3) [19]. However,
quantifying these states of clay particle dispersion for a particular polymer-clay
generally is experimentally difficult, though some progress is being made in this
direction.

Extensive experimental and theoretic studies have been carried out on
PCNs for engineering applications. The most successful examples of exfoliated
PCN are for nylon-6 and epoxy clay nanocomposites prepared via in-situ
polymerization [2, 3]. Nylon-clay nanocomposites have been used in under the
hood applications in Toyota automobiles. Although great improvements have
been made in PCN performance properties, realizing the full potential of these
remarkable materials and achieving their commercialization are still challenging
tasks. For some non-polar polymers, such as polyolefins, achieving facile
exfoliation of clay nanoparticles in the polymer matrix on a commercial scale still
is a challenging problem [15-18]. The complexity of PCN systems and the
limitations of quantitative structural characterization methods make it difficult to
establish morphology — processing — sthcture - property relationships [1].
Nevertheless, researchers have been developing new strategies in PCN

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chemistry to reach a better understanding in this field. Several of these emerging

advances, along with additional technical problems facing the field, are outlined

below.

1.2 Clay Surface Modificatlon

Smectites are hydrophilic materials, immiscible with most of engineering
polymers that are hydrophobic. Additionally, clay layers are closely stacked
through electrostatic interactions, with galleries occupied by hydrated alkaline or
alkaline earth cations, making it difficult for macromolecule diffusion. Therefore,
clays usually require organic modifications prior to the preparation of PCNs. The
most commonly used modification method is alkyl-ammonium-ion-exchange.
When the inorganic gallery cations are replaced by long-chain alkyl-ammonium
cations, not only the clay surfaces become more hydrophobic and, consequently,
more compatible with polymers, but also the gallery spaces expand, readily for
the intercalation of polymer chains. Depending on the cation exchange
capacities (CEC) of clays, the alkyl chain length, alkyl ammonium cations in clay
galleries may adopt structures that vary from lateral monolayer-, lateral bilayer-,
paraffin-, or lipid-type orientations, in the order of increasing gallery height [20].

While the onium-ion-exchange is widely used as a convenient and
effective clay surface modification method, it has several disadvantages. First,
the plasticizing effect of the alkylammonium cations may compromise the

7

benefits of clay nanolayer reinforcement [21, 22]. Second, ammonium cations
are susceptible to thermal decompose via Hoffmann degradation mechanism at
temperatures above 200 °C, which ares the processing temperature of some
engineering polymers, such as polyamide. The degradation of the organic cation
will cause the re-stacking of clay nanolayers and a loss of potential benefits [23,
24]. Third, during the preparation of PCN, the gallery alkyl ammonium cations
interact with the polymer matrices and clay layers, adding complexicity to the
nanocomposite systems [25]. Okamoto and coworkers argued that, in a
polylactide (PLA)- clay system that involves organo-clay with a large platelet
diameter and CEC value, the alkylammonium surfactant ions on the basal
surfaces of the clay form a restricted conformation due to steric limitations and,
consequently, hinder polymer intercalation [26].

Alkylimidazolium, a lkylphosphonium a nd a lkylpyridinum ions have been
used as clay modifiers for the reinforcement of high-melt-temperature
engineering polymers, such as polyamide-6 (PA-6), PA-6,6, polyethylene
terephthalate (PET) and polycarbonate (PC), as well as polymers bearing
medium-melt temperatures, such as polystyrene (PS) and polypropylene (PP)
[27-36]. Gilman and coworkers have used different dimethyl alkylimidazolium
salts to modify MMT, proving a higher thermal stability of the alkylimidazolium-
MMTs compared with alkylammonium-MMTs (Table 1) [34, 36]. The 1,2-
dimethyl-3-hexadecyl-imidazolium-MMT was exfoliated in PA-6 via melt-blending

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processes at 300 °C. Dong and coworkers compared the stability of imidazolium-
MMT (l-MMT) and ammonium-MMT (A-MMT) as reinforcement agents for PP
[29]. Both l-MMT and A-MMT exfoliated in PP upon in-situ polymerization.
However, l-MMT exhibited higher thermal stability than A-MMT, whether
analyzed as the pristine clay or as the composite form. The phase-structure
stability of the nanocomposite was investigated by melt-extrusion at 200 °C for
durations varying from 2-10 min. Re-stacking of clay layers was observed for PP
- A-MMT nanocomposites, due to the decomposition of alkylammonium modifier,
leading to inferior incompatibility between PP and MMT. The structure integrity
of PP — l-MMT nanocomposites was maintained during the extrusion.
Pyridinium- and imidazolium-based ionic liquids also have been used to modify
clay [37, 38].

Silylation of surface silanol groups is a common method to modify glasses
and silicates for use in polymer composite formulations. Silylation has also been
used for the polymer compatiblilization of smectite clays, layered silicic acids,
and silica materials [39, 40]. But this method is seldom used as the sole surface
modification technique in PCN studies, because only a small fraction of the
hydroxyl groups of a smectite clay are found at the surfaces of the platelet edges.
The remaining hydroxyl groups are buried in the second and third atomic planes
of the 2:1 layered structure and are unavailable for coupling reactions with silane
modifiers. On the other hand, one of the advantages of silylation is that the

10

treated clay contains less organic components (less than 5 wt%) than organicly
exchanged clay (25-45 wt%). Silylation has been used to graft functional groups
to organo-clay to improve clay dispersion and material properties of the
nanocomposites [40-45]. Bourgeat-Lami and coworkers synthesized organic-
inorganic nanocomposite colloids containing Laponite, which was modified by
either a monofunctional (y-methacryloxy propyl dimethyl ethoxylsilane) or a
trifunctional (y-methacryloxy propyl trimethoxysilane) silane coupling agent, via
emulsion polymerization [42]. Their results showed that trifunctional silane
coupling agents chemically locked the clay platelets into irreversibly stacked
nanolayers in tactoid form. However, the Laponite modified by monofunctional
silane derivatives can be satisfactorily dispersed in water and was found to cover
the styrene-butyl acrylate copolymer latex surface during the preparation of
colloids, as shown in Figure 1.4.

It has been realized that unmodified, purely inorganic clays are able to
disperse homogeneously in polymers and provide reinforcement, once the
surface properties of clays are compatible with polymers. Synthetic clays are
commonly used in this area since they usually possess lower aspect ratio than
natural clays. Haraguchi and coworkers reported nanocomposite hydrogels
containing rubbery poly(2-methoxyethyl acrylate) (PMEA), which is a
biocompatible polymer material, and inorganic clay (Laponite) [46]. The
composites were made via in-situ polymerization of MEA monomer and Laponite

11

in aqueous solution. The resulting composites were completely transparent and
had unique mechanical properties that they can undergo reversible necking
deformation of more than 1000%. The authors claimed that, during the
polymerization, PMEA monomers aggregated and excluded the hydrophilic clay
platelets so that the nanocomposite products were composed of PMEA
nanospheres that were covered by clay from the outside. When the composites
were under stress, the PMEA core provided the high elongation while clay
provided stiffness. More recently, McKinley and coworkers reported high-
performance elastomeric nanocomposites of polyurethane-polyether copolymer
and Laponite via solvent exchange processing technique [47]. Laponite
preferentially dispersed in and interact with the hard microdomains of
polyurethane, and, consequently, improved stiff, toughness and heat distortion
temperature of the polymer, especially when it reaches the weight loading to form
a percolative network. Pinnavaia and coworkers developed a strategy to prepare
a porous saponite-like silicate material and to use it to reinforce epoxy polymers
[48]. Due to its porous surfaces and irregularly stacked morphology, the
synthetic saponite dispersed uniformly in epoxy in the form of nanometric
tactoids. The nanocomposites exhibited improved strength, modulus, elongation

and toughness.

12

 

Figure 1.4 Cryo-TEM image of styrene-butyl acrylate particles coated by
monofunctional silane-treated Laponite clay nanolayers, as obtained by emulsion

polymerization in the presence of 10 wt% Laponite clay platelets [42].

1.3 PCN Preparation Technology

Three main methods have been used for the preparation of PCNs, namely,
solvent intercalation, in-situ polymerization, and melt blending. The solvent
intercalation method requires the polymer or pre-polymer to be soluble in a
solvent that also is capable of swelling the clay galleries. Polymer chains are
allowed to access the clay galleries, thereby, promoting the formation of either an
intercalated or exfoliated nanocomposites upon the removal of the solvent. This
method is commonly used for water-soluble polymers, and for composite film

13

fabrication [49-55]. For hydrophobic polymers, solvent intercalation methods
require the use of organic solvents, such as chloroform [56, 57]. From practical
point of view, organic solvent processing is environmentally challenging.

PCNs also can be prepared via in-situ polymerization of monomers in the
presence of the modified clay. The in-situ polymerization methodology is
particularly useful for the preparation of therrnoset PCNs and for thennoplastics
that can be readily prepared under near-ambient conditions. Nylon-6 - exfoliated
MMT nanocomposites, the first example of an exfoliated PCN, were prepared by
polymerization of a mixture of s-caprolactam and MMT modified by the onium ion
of a o, w-amino acid [8, 9]. This method has also been applied to epoxy,
polyolefin, PS, PA, and polyimides (Pl) systems [22, 58-63]. The advantage of
the in-situ polymerization approach is the ease of clay — pre-polymer processing.
It is usually much easier to pre-intercalate or to pre-exfoliate a clay with a
polymer precursor than the polymer itself.

The melt-blending method blends a mixture of polymer and clay at a
temperature above the softening point of the polymer [34, 64-68]. This method
has been identified as the most desired approach to the preparation of
thermoplastic PCNs because it does not require the use of solvents or the re-
tooling of processing equipment and conforms to standard industrial methods for
polymer processing. However, thermoplastic polymers. especially non-polar
polymers are not good wetting agents for organoclays. Moreover, many

14

organoclays are not stable under melt processing conditions due to the low
thermal stability of the organic modifiers. Thus, it generally is difficult to achieve
organoclay nanolayer exfoliation under common melt processing conditions. To
overcome these limitations, extra modification of the polymer needs to be carried
out to achieve compatibility with the clay and to facilitate intercalation.
Commonly used technique in the preparation of polyolefin — clay nanocomposites
is to incorporate maleic anhydride grafted polyolefin as a compatibilizer, although
the addition of the lower molecular weight component lowers the mechanical
properties of the polymer matrix [13-18, 69].

The mechanism of polymer intercalation in clay has been studied
extensively. At thermodynamic equilibrium, the dispersion of the platelets in
polymer matrix is favored only if the enthalpy changes overcome the entropy loss
of the confined polymer chains within the clay galleries [70]. Vaia and Giannelis
studied the polymer melt intercalation in organo-clay using a mean-field, lattice-
based model, and pointed out that the entropic penalty of polymer confinement
may be compensated by the increased conformational freedom of surfactant
chains a s the layers 8 eparate. The dispersion of clay was e nergy favored by
maximizing the polymer-surface interactions and minimizing the unfavorable
apolar interactions between polymer and surfactants [25].

Epoxy-based PCNs prepared via in-situ polymerization method provide
examples for clay exfoliation mechanism research. Pinnavaia and coworkers

15

pointed out that the exfoliation of clay is possible when comparable intra- and
extragallery polymerization rates were achieved by incorporating acidic exchange
cation as both a clay surface modifier and a catalyst for intergallery epoxide
polymerization [71]. Chen et al. divided the interlayer expansion process into
three stages: initial interlayer expansion due to the monomer and curing agent
diffusion, steady-state interlayer expansion due to the intergallery polymerization,
and the cessation of interlayer expansion [72]. More recently, Jana and
coworkers argued that the elastic force exerted by cross-linked epoxy molecules
inside the clay galleries was the primary force for exfoliation [73]. The elastic
force pushed the outermost clay layers out from the tactoids against the
opposing forces arising from electrostatic and van der Waals attraction and shear
until gelation of the system occurs (Figure 1.5). Complete exfoliation can occur if
the ratio of shear modulus to complex viscosity was maintained at or above 2-4
1/s, so that the elastic forces inside the galleries outweighed the viscous forces
of the extragallery epoxy.

Organic solvents have been used to assist in the preswelling of an
organo-clay in the presence of a prepolymer [22, 74]. Hasegawa reported the
preparation of a PCN by compounding nylon-6 with a Na*-MMT aqueous slurry
[75]. After the evaporation of water, Na"-MMT layers were found to be exfoliated
and dispersed homogeneously in the nanocomposites. The properties of these
novel nanocomposites were nearly equal to those of conventional nylon-organo-

16

clay nanocomposites. Also, Mai and coworkers synthesized exfoliated silicone
rubber - clay nanocomposites based on a siloxane surfactant modified clay

aqueous suspension [76].

 

clay layer
+

+
. l clay layer

', :l 1‘
_. epoxy monomer
- / /
l + l +
l .. - - - - I

l l l l l l l extra-gallery viscous force
(c)
l _ l
1 1 + :I
van der Waals force

++++r fl]

1 l l 1 II I

Figure 1.5 Schematic illustration of the intercalated state of the exfoliation

 

(a)

1.92 nm quaternary ammonium ion

 

 

 

 

 

elastic force , _
electrostatic attractive

force

 

process showing the forces acting on a pair of clay layers: (a) organically
modified clay; (b) epoxy-intercalated state; (c) forces acting on a two-particle clay

tactoids [73].

17

More recently, He and co-workers developed a so-called “slurry
compounding’ process to prepare exfoliated epoxy-clay nanocomposites [44,45].
The exfoliated morphology of the clay in aqueous suspension can be transferred
to the epoxy matrix. The mechanism of exfoliation was believed to involve the
silylation of hydroxyl groups on the basal planes of the clay structure. However,
since the hydroxyl groups are buried in the second and third planes of oxygen
atoms and are inaccessible for reaction, the proposed mechanism is unlikely.
Instead, silylation most likely occurred only at the edges of the clay platelets.
This would be consistent with the observation that the method required only
small amounts of silane as a surface modifier. Figure 1.6 shows TEM images of
the exfoliated clay nanolayers in the epoxy matrix. The same group also
reported preparing PP — clay nanocomposites by reactive compounding PP with
an epoxy - clay master batch made via “slurry compounding” [77]. Because
epoxy and PP are immiscible, maleic anhydride grafted PP was incorporated as
both a compatiblizer and an epoxy curing agent. Most clay particles were
exfoliated or dispersed into small stacks clay layers in the resulting
nanocomposites. The sluny compounding approach for nanocomposite
formation is deserving of verification and extension to other polymer - clay

systems.

18

 

Figure 1.6 TEM images of an epoxy nanocomposite (2.5 wt% clay) prepared by

slurry compounding [45].

Supercritical carbon dioxide fluid, as an environmentally friendly,
inexpensive, low-viscosity and nonflammable liquid, has been used to replace
conventional solvents in the preparation of PCN compositions based PS,
Poly(ethylene oxide) and poly(methyl methacrylate) matrices [78-80]. Manke and
coworkers patented a unique technique using supercritical fluid to delaminate
clay [81]. A catastrophic depressurization of a supercritical - clay system and
the associated catastrophic volume change of the fluid caused the clay layers to
exfoliate, as judged by XRD patterns of the solids. The inventors claimed that
the treated silicate particles were ready for polymer reinforcement purposes.
More recently, Gulari and coworkers patented a technique to exfoliate contacted
particles and reduce reaggregation of the structures by supercritical fluid diffused

19

with coating agents, such as polydimethylsiloxane [82]. These techniques
eliminate the use of any clay surface modifiers. However, they require high-

pressure environments.

1.4 Characterization of Clay Nanolayer Dispersion

PCN studies rely on accurate characterization of the nanometer-scaled
clay morphology in polymer matrices. The most commonly used methods are
wide-angel X-ray diffraction (WAXD) and transmission electron microscopy
(TEM). WAXD has been used to measure the position, shape and intensity of
the basal reflection of both initial clay particles and clay layers dispersed in
polymers (Figure 1.7a). An increase in interlayer spacing due to the intercalation
of polymer results in a lower-angle shift of the 001 peak, compared with the basal
reflection of the initial clay particles. The peak also broadens when the layer
orientation is disordered. The disappearance of the 001 reflection indicates the
formation of some form of an exfoliated morphology. Small angle X-ray
scattering (SAXS) technique can be used to detect interlayer spacings of 200 A
and beyond. SAXS has also been used to determine three-dimensional clay
layer orientation in PCNs [83]. Although WAXD has been a convenient method
to determine the d-spacing, the X-ray line broadening is directed relative to the
crystal size and number of diffracting layers. Therefore, the d-spacing of only a
few stacked layers or a relatively disordered intercalated structure may not

20

be detected using WAXD. Additionally, WAXD does not provide information on

structure homogeneities throughout the composites.

in
U'

 

Original OMLS

      

AlllllllllllAllllL

 

1 Ir; run I.
'lrvl'II|IYI

 

Intensity /A.U.

 

allllnlknlllnlll

   

 

Exfoliated

O
l v .-
nnlninilinn 'alll A

 

 

:3
,—
-
u
.

 

u
A
OK
no
5

 

2 theta (degrees)

Exfoliated 200 nm

Figure 1.7 (a) WAXD patterns and (b) TEM images of three different types of

PCNs [4].

21

High magnification and high resolution TEM has been a powerful
technique for PCN characterization. The edge of individual clay layer can be
directly imaged (Figure 1.7b). However, the sample preparation of TEM requires
ultrathin-sectioning of the composite into 70-nm thick slices, which is a delicate
and time-consuming process. To determine the overall homogeneity, a large
number of TEM images must be taken from different regions of the composite
sample. Quantitative analysis of a large number of images can be carried out by
enumerating the interlayer spacing and layer aspect ratios of the platelets [84,
85]. The combination of WAXD and TEM usually provides an reasonable
representation of PCN morphologies.

Optical microscopy and scanning electron microscopy (SEM) have been
used to investigate clay dispersion on a micrometer scale, and can be
complementary methods to WAXD and TEM. SEM is also a useful tool to
investigate the fracture surfaces of the composites [44].

Vanderhart and coworkers first used ‘30 solid state NMR to determine
clay dispersion in polymers [86]. The spin-exchange interaction between the
unpaired electrons of paramagnetic Fe atoms in clay particles shortens the
longitudinal relaxation time (T1”) of protons within 1 nm of clay surfaces, and,
further, contributes to the T1” of bulk composite by spin diffusion. Better clay
dispersion leads to greater average paramagnetic contribution to T1“. The
success of this method depends on two factors: the paramagnetic components of

22

the clay, such as Fe3+ in MMT, and an efficient spin diffusion processes. NMR
has been an effective quantitative method to characterize polymer - MMT
nanocomposites based on PS [87], nylon [86, 88, 89], styrene-acrylonitrile
copolymer [90] and poly(£-caprolactone) [91]. However, in rubbery polymers,
where the dipolar interactions are too weak, NMR cannot be applied [90].

Atomic Force Microscopy (AFM) has also been used to investigate clay
dispersion in PCNs. After chemical etching, the silicate particles were detectable
on the composite surfaces. AFM can also be applied directly to the composite
thin-sections [92-94]. Figure 1.8 shows a phase contrast AFM image of an epoxy
— fluromica nanocomposite, where the intercalated morphology can be clearly
seen. AFM has been found to be especially useful in mapping the heterogeneity
of polymer blends [95-97]. Karger-Kocsis and coworkers studied polyamide-6 —
PP — organo-clay nanocomposites by AFM and found that organoclay was
located only in the PA-6 phase of the uncompatibilized blends [97]. When maleic
anhydride grafted PP (MAH-g-PP) was added as a compatiblizer, the clay
became embedded in the PA6-g-PP phase only.

Lo’pez-Manchado and coworkers introduced a novel PCN
characterization method by measuring the freezing-point depression of a solvent
encapsulated the swollen composite [99]. It has been found that vulcanizates
containing highly dispersed clay particles exhibited a larger freezing point
depression for the encapsulated solvent than the pristine rubber matrix. The

23

difference in freezing point depressions was attributed to a tighter network and
smaller solvent cages for the composite. However, the change in the freezing
point depression was comparatively small, and this method cannot reliably

provide quantitative information on the degree of clay particle dispersion.

 

Figure 1.8 A phase contrast AFM image of an epoxy - synthetic fluoromica

nanocomposite [92].

Researchers at Elementis Specialties, Inc. recently developed a technique
to use FTIR to characterize the exfoliation and alignment of clay layers in

24

polymer matrices (W. ljdo, personal communication). The bandwidth of the
silicate vibrations of an organo-hectorite was found to narrow significantly upon
layer delamination (Figure 1.9 a). It was possible to construct a calibration curve
relating the bandwidth for the in-plane Si-O vibration vs. the fraction of nanolayer
dispersion. A solvent containing different amounts of a chemical activator, such
as propylene carbonate, was used to control the level of clay delamination
(Figure 1.9 b). When other clays were used to reinforce polymers, the dispersion
could be determined by measuring the in-plane Si-O bandwidth in comparison to
the hectorite calibration curve. The same researchers also have used polarized
FTIR to evaluate clay alignment in PNC films. Since the polarized FT IR
absorbance changes with the electron density in the polarized direction, tilting a
PNC film sample containing oriented clay particles in the path of a z-axis
polarized IR beam will cause a decrease in the intensity of the in-plane Si-O
absorbance and an increase in the out-plane Si-O absorbance (Figure 1.10).
This phenomena was not observed in composite films with randomly orientated

clay layers.

25

Figure 1.9 (a) FTIR spectrum of hectorite powder and delaminated hectorite in
water; (b) a calibration curve of clay in-plane Si-O bandwidth vs. percent activator
added and percent dispersion. (Figures provided by W. ljdo, Elementis

Specialties, Inc.)

26

(a)

 

FT-IR Hectorite clay
powder

    
 
    

delaminated

 

 

 

1 150 1 100 1050 1000 950 900

Wavenumber [cm"]

Percent nanodispersion

0 20 40 60 80 100
.......ll....f..l

A
0'
v

 

 

ln-plane Si-O bandwidth (FWHM) [cm"]

 

 

 

u ' u I
..4 ' : ' , I :
r n . r ' ‘
r . I I ' .
25 TTTTT i1 T l T I {1 T I r I] IIIII [I I I T I i Y I TTU I T1 Tfil

Percent activator added

27

Polarization along z-axes
(BENTONE 108 in LDPE)

 

 

0.7
‘ —0 deg tilt ‘ '

j — 15 deg tilt I .\Out-of-plane f 0'6

——— 30 deg tilt l - Si-O band i 0 5
l - 45degtilt l I §
L —60 deg tilt ' -. 0.4 g
\l a
.n
In-plane " 0-3 ‘1

Si-O bonds

-» 0.2

7~ 0.1

l ‘ ‘ ‘ . , , ._.¢g§§;E? on

 

 

1150 1100 1050 1000 950 900 850
Wavenumber [cm'1]

Figure 1.10 z-axis polarized FTIR spectra of a LDPE film containing oriented
BentoneTM clay layers, with different tilt angle. Inset: schematic representation of
clay in-plane and out-of-pane Si-O bonds. (Figure provided by W. ljdo, Elementis

Specialties, Inc.)

1.5 Nanolayer Orientation

To achieve control on the orientation of anisotropic fillers has been a
challenge in nanocomposite studies. Conventional techniques, such as fiber
weaving, obviously are not feasible for clay nanoparticles. Theoretical modeling
studies of PCN composition have shown that substantial improvements in
mechanical and barrier properties are possible for a composite reinforced by

28

aligned clay layers, compared with those containing random clay dispersions
[100-103]. It has also been predicted that the enhanced mechanical properties
observed for PCNs based on elastomers are due to the ability of realignment of
clay layers under stress [61]. Relatively fewer experimental studies have been
carried out on achieving good alignment of clay layers in PCNs. Vaia and
coworkers applied AC electric field during the curing of epoxy-organo-clay
nanocomposites [104]. The in-situ SAXS showed that the clay layers reoriented
parallel to the electric field, and that, upon thermal curing, the complete
development of alignment occurred before the onset of chemical crosslinking
(~80 °C). At the same time, the clay layers swelled toward an exfoliated
morphology (Figure 1.11). The driving force for clay particle alignment was
possibly the induced dipole of the ‘mobile’ exchangeable organic cations on clay
surfaces. Compared with epoxy — clay nanocomposite with random clay
morphology, those with aligned morphology exhibited lower coefficient of thermal
expansion, higher modulus and better optical clarity. More recently, Vaia and
coworkers reported using a magnetic field > 1 T to align clay in an epoxy matrix
before curing [105]. Clays, particularly MMT that contains antiferro- or
ferrimagnetic impurities, were able to align parallel or perpendicular to the
magnetic field respectively, indicating that the composition of the clay has great
impact on the alignment, presaging the design of triaxial nanoparticle
reinforcement on the basis of magnetic properties.

29

3"
O

 

 

 

 

~ 44
0.8 -
En 0.5-
to ~40 -<
°- O.4 - :.
g 8
..., as o
9, 0.2-
5 as
'5 0.0«
+ D-Spacing _ 3‘
.02 . I Orientation Parameter

0 40 80 120 160 200
Figure 1.11 Hermans orientation parameter, Sd (squares), and layer spacing
(circles) determined from in-situ scattering (3 °C/min). Note that the intensity of
the basal reflection (d001=36A) rapidly decreases above 120 °C, effectively
disappearing by 180 °C. Orientation parameter above 120 °C was estimated from

the featureless scattering around the beam stop [104].

1.6 Synthetic Materials for Polymer Reinforcement
Although naturally occurring clays are inexpensive and readily available,
they have certain disadvantages for PCN applications, including the difficulty of

purification and the need for organic modification. Researchers have been

30

looking for alternative synthetic clay-like materials for polymer reinforcing
purposes. Clay materials can be synthesized under various laboratory conditions,
ranging from ambient pressure and temperature to extreme hydrothermal
conditions [106]. The products usually exhibit high purity and uniform particle
morphology and charge distribution. More importantly, a well designed synthesis
strategy raises the possibilities of tailoring clay properties to meet certain

research requirements.

 

 

Relataive Intensity

 

 

 

 

0 20 25 30 35 40
Degreesze

Figure 1.12 XRD patterns of (a) natural Na-hectorite, herein the peaks labeled
by * are due to (Mg, Fe)SiOa and CaCOa impurities; (b) synthetic Li-hectorite, and

(c) synthetic PVP-hectorite. Typical clay hkl reflections are identified [107].

31

Carrado and coworkers have developed a general technique to
synthesize hectorite-like clay materials by refluxing a freshly precipitated brucite
in solutions containing LiF and silica precursors [107]. A series of water-soluble,
cationic or electrically neutral organics, varying from small molecules like
tetramethyl ammonium to large polymer molecules like polyvinylpyrrolidone (PVP)
can be present in the course of hectorite synthesis. When cationic organics were
present, they ended up as the interlayer cations. The presence of neutral
organic molecules during synthesis required larger amounts of LiF for successful
hectorite synthesis, and both the neutral organic and Li+ were incorporated into
the interlayer 5 paces [107]. Figure 1 .12 s hows the X RD p attems of a n atural
Na+-hectorite, synthetic Li*-hectorite and a PVP-hectorite, showing the
characteristic XRD diffraction peaks and higher level of purity of the synthetic
hectorites. PVP-hectorite composites exhibited an increase in basal spacing with
polymer uptake from 16.5 to 23.5 A. The authors proposed two potential
applications for synthetic-hectorite [107]. One was for the preparation of porous
materials for potential use as catalysts or catalyst supports upon the removal of
the organic polymer phase by combustion. A correlation was observed between
catalyst mesoporous size and the size and concentration of a polymer used in
the synthesis. The other potential application was for the preparation of organic-
inorganic composites through a one-step process by allowing more polymer
intercalation during the hectorite synthesis. The highest organic loadings, up to

32

86%, were observed for the polyaniline-hectorite composites. These composites
had a so-called semi-exfoliated structure, in which small clay crystallites were
intercalated by polymer and well-dispersed within a continuous polymer matrix.
Carrado and coworkers also reported a one-step synthesis strategy of organo-
smectite clays with organics grafted on the interlayer surfaces by using
organotrialkylsilane as the silica sources [108]. These synthetic clay may be an
interesting substitute for conventional organo-clay prepared from ion-exchange of
natural clay.

Chastek, Stein and coworkers synthesized hexadecyl-functionalized
lamellar silicates C16-LMS, aluminosilicates C16-LMAS and C16-SiOz-LMAS by
sol-gel synthesis, whose structures are shown in Figure 1.13 [109]. C16-LMS
consisted of single or double layers of tetrahedral silicates. C16-LMAS had a
pyrophyllite-type structure with an octahedral aluminum layer sandwiched
between two tetrahedral silicate layers. C16-SiOz-LMAS had additional silicate
groups in the structure, compared with C16-LMAS and, thus, had a larger
inorganic layer thickness and greater structural disorder. Hexadecyl groups were
covalently attached to the inorganic layers in all three synthetic silicates, as
verified via 29Si-NMR. In a subsequent report, Stein and his coworkers
investigated the dispersion of C16-LMS, C16-LMAS and C16-SiOz-LMAS in
different organic solvents and in polystyrene [110]. The solvent study showed
that the synthetic clays formed stronger gels with aromatic solvents, such as

33

toluene, than with linear, branched or cyclic alkyl solvents, indicating that the
synthetic clays were more suited for dispersion in polystyrene, which has a
similar structure as the aromatic solvents. The clay — toluene gel strength
increased with increasing clay concentration, although SAXS showed no
intercalation of toluene in clay galleries. These three synthetic clays and Cloisite
20A (an organo-MMT) were compared in polystyrene - clay nanocomposites
prepared from melt-blending process. The rheology study showed that, at the
same inorganic content, composites made from C16-LMAS and C16-SiOz-LMAS
had higher elastic moduli (6’) than from C16-LMS, due to the higher aspect ratio
and better dispersion of the former two clays in polystyrene. However, SAXS
showed that no polymer intercalation occurred in the composites. The
flocculation of clay caused by the interactions of hydroxyl groups on clay
surfaces, and the strong interactions between the interlayer covalently bonded
alkyl groups made the polymer intercalation more difficult. Polystyrene-Cloisite
20A exhibited highest 6’, attributed to the better dispersion and possibly higher

layer stiffness of Cloisite 20A (Figure 1.14).

34

 

 

 

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A n r A v

Figure 1.13 Proposed schematic structures of the functionalized clays. Two
possible structures for C16—LMS include a single layer structure and a double
layer structure which is shown in (A). The proposed structure of C16-LMAS (b)
' with both tetrahedral and octahedral Al present in the octahedral layer. In C16-
SiOz-LMAS (C) additional silicate groups without alkyl chains are present and
tetrahedral Al is likely to occupy some of the sites in tetrahedral layer [109].
35

reactant concentrations.

 

105':me r recur:

 

 

. *.*
Id tittt*****’**.
4 ..:
I
10 1 ... ‘
Ibleualllllll"'.. 5“
s “A‘ AC.
13 AAAAAAAAAA“‘ -c’
13‘”: O
r‘r'ho '3'
a .. O 300 00 0 v 0
4 . 4 r ".) ‘
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21 0
i o
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‘I
10 1 d
i c
: PS 0
4 00
O
10 .Enmq ..“mq ..nflw -4n".
0.01 0.1 1 10 100

‘9 (rad/s)

Figure 1.14 Elastic moduli of PS - clay nanocomposites using (a) C16-Si02-
LMAS, (b) C16-LMAS, (c) C16-LMS, and (d) Cloisite 20A from melt rheology

data obtained at T=150 °C, and 6.3 wt% inorganic material [110].

Clearfield and coworkers have used synthetic d-zirconium phosphate (d-
ZrP) to make epoxy nanocomposites [111, 112]. The a-ZrP had a higher ion
exchange capacity and narrower particle size distribution than MMT. The size

and aspect ratio of o-ZrP can be controlled by varying the reaction time and

readily exfoliated in epoxy, while o-ZrP remained unswelled. A 1.9 vol% M-o-ZrP

36

 

F l

A monoamine modified o-ZrP (M-a-ZrP) has been

loaded epoxy nanocomposite showed 50% increase in tensile modulus and 10%
increase in yield strength. However, the ductility was dramatically reduced. or-
ZrP has also been used to study the fracture behavior of epoxy nanocomposites

[113].

 

 

-2.

Heat flow/ng
O

.4.

 

.5-

53°C

 

 

 

320' 6 'io'ab'sb'sb'rbo
17°C

Figure 1.15 DSC traces of the organosilicate. Inset shows the phase transition

from solid to liquid upon heating [114].

Giannelis and coworkers have reported surface-functionalized
nanoparticles with liquid-like behavior [1 15]. For instance,

37

[(CH30)3SI(CH2)3N+(CH3)(C1oH21)2]Cl' functionalized silica nanoparticle with
sulfonate anion R(OCH20H2)70(CH2)3303' as counterion appeared as a clear
liquid at room temperature. This concept has been extended to various
nanoparticles, including magnetic iron oxide nanoparticles, layer-like
organosilicate, polyoxometalate cluster, anatase nanoparticles, and DNA
oligonucleotide [116]. Layered organosilicate nanoparticles have been
synthesized by refluxing octadecyltrichlorosilane in toluene with the presence of
a small amount of water [114]. The solid product can be readily dispersed in
various organo solvent, and exhibit reversible solid-liquid transition upon heating
(Figure 1.15). These solvent-free, zero-vapor pressure, liquid-like nanostructures
circumvent toxic solvents, possess unique conducting, magnetic or
electrorheological properties, and exhibit low dimension restrictions, bringing new
scientific and technological opportunities.

Meso-structured silica with large mesopore size has been a promising
reinforcing agent for polymer materials. Improvement in mechanical properties
have been reported based on nylon-6,6 [117], poly(vinyl acetate) [118], poly((3-
trimethoxysilyl)propyl methacrylate) [119] etc. Very recently, Pinnavaia and
coworkers reported epoxy — silica mesocomposites with enhanced tensile
properties and oxygen permeability [120]. The mesoporous silica, denoted MSU-
J, was synthesized using a d,w-diamine polypropylene oxide (Jeffamine D2000).
It had a wormhole framework structure with a pore diameter of 5.3 nm, a pore

38

volume 1.41 cm3/g and a surface area 974 mzlg. Both as-made MSU-J and
calcined MSU-J, where the Jeffamine D2000 was removed in the latter case,
dispersed uniformly in a rubbery epoxy matrix. Compared with epoxy - organo-
clay nanocomposites, epoxy — silica mesocomposites exhibited comparable
improvement in mechanical properties and thermal stability (Figure 1.16). The
tensile modulus, strength, toughness, and elongation-at—break for the
mesocomposites were systematically reinforced by up to 4.8, 5.7, 1.6, and 8.5
times, respectively, in comparison to the pure epoxy polymer. Additionally, no
organic modification was needed for MSU-J, and the composites were prepared
with a simple mix and cure method. Another remarkable property of the epoxy -
as—made MSU-J mesocomposites is that the oxygen permeability increased
dramatically at loadings 25 wt% due to the partitioning of curing agent between
the as—made mesostructure and the liquid prepolymer, and, consequently,
reduced chain cross-linking in the vicinity of the silica particles. Such
phenomenon was not observed in epoxy - calcined MSU-J mesocomposites
since the curing a gent partitioning was not allowed (Figure 1 .16). Balkus a nd
coworkers have also reported improved gas permeability in polysulfone - MCM-
41 membranes prepared by a solvent-blending method. A 30 wt% MSM-41
loaded polysulfone membrane had the oxygen, nitrogen, carbon dioxide,

methane permeability increased by 2.5, 2.9, 2.7, and 2.6 times respectively [121].

39

 

 

 

 

15 16
(a)
A “'8
ii" 3
2 3
v10- g;
3 -e I”
3 i
‘3 e
2
2 " 3
... 5_ A
Q
3
5 (b) 1:
|— in
P2 v
9 r V r ' r - r r C
12

2 4 Is ' é ' 1'o
Silica Loading (wM v.)

 

 

'- 1 e
30- (c) -roo : ,(I
.. ’ ~eoo F
325- §‘..'° ‘
V ' 10—
a; mg 5 .
2"“ = E
~400§ o
.115- I‘ll .. i
E #300? E 5‘
0,10- 400‘" '8
100 a.
,, .
c c o o .

 

 

 

 

 

 

' ' ' ' I I l I l

2 4 6 8 10 1'2
Silica Loadings (wlw %)

°-

2 4 i ' I ' 13 '1':
Silica Loading (wlw %)
Figure 1.16 (a) Tensile modulus, (b) tensile strength, (c)strain-at-break, (d)
toughness, and (e) oxygen permeabilities epoxy composites containing different

loadings of as-made MSU-J (solid circles), and calcined MSU-J (open circles)

mesostructured silica [120].

40

1.7 Conclusion

This chapter has reviewed certain aspects of PCN chemistry, particularly
clay nanolayer modification, PCN preparation and characterization, nanolayer
alignment and alternatives to natural clays as polymer reinforcement agents.
Conventional techniques, such as onium-ion exchange, WAXD, TEM, have been
and will continue to be effective for PNC study. New techniques and concepts
are rising everyday, and holding additional promise for routine use in future PCN
studies.

Despite all the efforts, it is still a great challenge to establish the
morphology — processing - structure - property relationships of polymer - clay
nanocomposites. To reach this goal, it is important to understand and evaluate
the necessity of clay organic modification, considering the advantages of the
modifiers as surface compatibilizers and their disadvantages in causing
plasticizer effects and low thermal stability. It is also important to further develop
polyolefin nanocomposites, since polyolefins have a wide range of commercial
applications. Current techniques in forming polyolefin — clay nanocomposites
requires extensive organic modifications and has not yet given satisfactory
results on commercial scales. This thesis aims at developing organic-modifier-
free preparation pathways for the formation of polymer - metal oxide

nanocomposites to prove the feasibility of the proposed study, to discover novel

41

inorganic metal oxides for polymer reinforcement purposes and to attract more
research interest in this area.

In the following chapters, both thermoset polymer epoxy and thermoplastic
polyolefins are chosen as polymer matrices. Three different metal oxides with
unique morphology and surface properties are studied: synthetic saponite,
palygorskite, and mesoporous silicas. Specific research interests include:

0 Synthesis of clay-like silicate materials in order to control the
morphology and composition of the products by choice of synthesis
conditions.

. Develop strategies for using inorganic metal oxides as polymer
reinforcing agents. The inorganic metal oxides include synthetic
saponite, palygorskite, and mesoporous silicas. The polymer
matrices include epoxy and polyethylene.

- Compare the properties of nanocomposites reinforced by inorganic

metal oxides with those of nanocomposites formed from organo-clay.

42

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50

CHAPTER 2

SYNTHESIS OF SAPONITE-LIKE PLATEY SILICATES AND THEIR ORGANO

HYBRID DERIVATIVES

2.1 Introduction

The term “clay” is usually used to describe particles of mineral sediments
that are smaller than 0.2 pm in size [1]. Clay minerals are some of the most
important industrial minerals, and find applications in process industries,
agriculture, environmental remediation, construction. pigments, coatings,
ceramics, catalysts and polymer materials [2]. The major structural building
blocks of clay minerals are tetrahedral SiO4 units and octahedral M06 units,
where is M = Al, Mg, Fe, Zn and the like. Saponite, the clay of interest in the
present study, belongs to the smectite clay family. Natural saponite has a
stacked layered morphology wherein the layers are made of a 2:1 ratio of
tetrahedral silica sheets and magnesia sheets, and a unit cell formula of (M2,.
ynH20)[Mgg.,,AI,,](Si3_,.Al,.)Ozo(OH)4, where M“ is the exchangeable gallery cations.
lsomorphous substitution by Al takes place in both tetrahedrals and octahedrals
and generates negative layer charges that are balanced by exchangeable alkali

or alkaline earth cations located inside the galleries.

51

Natural clay minerals are abundant and inexpensive resources. However,
they usually contain impurities and their purification is a time-consuming process.
The nanolayer compositions and corresponding surface charges are usually
heterogeneous, even in the same clay sample, making it difficult to obtain reliable
results in systems where the clay surface properties play important roles.
Natural layered clays usually exhibit a well-stacked layered morphology, and
their low accessible surface area limits some of their potential applications, such
as catalysts. On the other hand, synthetic clays are promising substitutes to
natural clays. Smectic clays can be synthesized through sol-gel methods under
conditions varying from mild room-temperature conditions to extreme
hydrothermal conditions [3-8]. Synthetic clays exhibit reproducible high purity,
uniform composition and surface properties. These properties are controllable
through mediation of synthesis conditions.

Natural smectic clay minerals are hydrophilic inorganic materials. Their
organic modification can be achieved by either organo-cation exchange or
silylation of the edge hydroxyl groups. Previously, attempts have been made to
synthesize organo-clay derivatives via one-pot sol-gel synthesis using
organosiloxanes of the type RSi(OR')3 as the silicon source for clay formation [9-
12]. The products contain gallery alkyl groups, which are covalently bonded to
the basal silicon centers. Consequently, the products exhibit a stacked
morphology and hydrophobic surface properties. However, since the alkyl

52

groups of the organosiloxane cannot serve as bridging groups for neighboring
silica tetrahedra, the products have significant structural distortion, in comparison
to a well-ordered clay structure [10]. It is also possible that inversion of silica
tetrahedral may occur, so that the apical alkyl group points to the clay gallery,
while the oxygen atoms are bridged to adjacent silicon center of the tetrahedral
sheet

Bissiloxanes of the type (R’O)3Si-R-Si(OR’)3 have been used for the
synthesis of hybrid zeolites and periodic mesoporous organosilicates (PMOs) [13,
14], which have 3-D structures made exclusively of tetrahedral SiO4 and NO.
centers. When bis(triethoxylsilyl)methane (BTESM) is used as the silicon source
for clay synthesis, the bridging methylene group might substitute in part for the
bridging oxygen atoms in the basal plane of the clay structure, as illustrated in
Figure 2.1. The organically modified clay layers differ from organo-cation-
exchanged clays or silylated clays in that the organic groups are incorporated
directly into the clay lattice structure. These novel hybrid clay materials may
have promising application as absorbances, polymer reinforcing agents, etc.

In the present study, inorganic saponite-like layered silicates also were
synthesized through modification of literature methods [3, 4]. Depending on the
choice of reagents and synthesis temperature, the products exhibited different
layer stacking order that varied between irregularly stacked layers to a well-
stacked layered morphology. The structure and surface properties of the

53

synthetic saponites were characterized using XRD, NMR, TEM and N2
adsorption-desorption isotherms. The organic hybrid derivatives were prepared
using bissiloxanes as the Si sources, specifically bis(triethoxylsi|yl)methane
(BTESM) and bis(triethoxylsilyl)ethane (BTESE). The former siloxane afforded a
hybrid clay-like structure with methylene groups incorporated in clay lattice
stnicture in place of bridging oxygen atoms, while the latter could not direct the

formation of clay structure due to the bulky size of the ethylene groups.

 

Figure 2.1 Schemes illustrating the difference in tetrahedral sheet structure for (a)

an inorganic clay and (b) an organo hybrid clay made from BTESM.

2.2 Experimental
2.2.1 Materials

BTESM and BTESE were purchased from Gelest and used without
purification. All other chemicals were purchased from Aldrich and used without
punficafion.
2.2.2 Inorganic saponite synthesis

A synthetic saponite, denoted Mg-SAP, was prepared at 90 °C by

modification of previously described methods [3] using water glass solution (27
wt% silica, 14 wt% NaOH), AI(N03)3-9H20, Mg(NOa)2-6H20, NaOH and urea in
the molar ratio 3.6 : 0.40 : 3.0 : 3.0 : 10 per 400 moles of water. In a typical
synthesis, 38.48 g of the sodium silicate solution (0.173 mol silica) and 7.218 g
(0.019 mol) Al(N03)3°9H20 were mixed in 346 g deionized water at room
temperature for 1.5 h. Then 28.8 g (0.48 mol) urea 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 consitions for 30 min,
followed by the introduction of 37.0 g (0.144 mol) of Mg(NOa)2-6H2O as a solid.
The mixture was refluxed under stirring conditions at 90 °C for 24 h in a 1 L flask.
Then 70 mL of hot NaOH aqueous solution containing 5.64 g (0.141 mol) NaOH
was added to the synthesis mixture. The mixture was allowed to age for another

15 h. Finally, the product (denoted Mg-SAP) was centrifuged, triple washed with

55

deionized water, and dried at 50 °C under N2 flow. Around 20 g Mg-SAP can be
obtained, indicating the yield to be close to 100%.

A synthetic saponite, denoted Zn-SAP, with Zn as the dominant
octahedral species was synthesized using the same method as the Mg-SAP
synthesis [3], except that Zn(NO3)2-6H20 was used instead of Mg(N03)2-6H20.

A well-crystallized saponite, denoted SAP-200, was obtained under
hydrothermal conditions [4]. In a typical synthesis, 0.35 g of 28wt% NH4OH
solution (2.8 mmol NH4OH), 0.61 g (3.0 mmol) Al(OC3H7)3, 1.6 g (27 mmol)
fumed silica, and 4.77 g Mg(CH3002)2 (22 mmol) were introduced into 40 g
deionized water one by one, and the mixture was stirred for 10 minutes after the
addition of each reagent. The mixture was then stirred at room temperature for
24 h before being introduced into an autoclave vessel. The crystallization was
carried out at 200 °C for 6 days. The product (denoted SAP-200) was
centrifuged, washed with water, and dried at 50 °C under N2 flow. A 2.2 9
quantity of SAP-200 can be obtained, indicating the yield to be 80%.

2.2.3 Ion exchange of Mg-SAP

Prior to the exchange reaction, 120 mL 5 wt% aqueous suspension of Mg-
SAP was stirred under ambient conditions for 24 h. The Cs+ exchange reaction
was carried out by adding 3.37 g (20 mmol) CsCl to 40 mL of the 5 wt% Mg-SAP
aqueous suspension, followed by stirring for 8 h and centrifugation once to
separate the solid from the liquid. The supernatant liquid was discarded, and the

56

solid was again suspended in 40 mL water. The reaction was repeated for 3
times to ensure a complete ion exchange. The product (denoted Cs-[Mg-SAP])
was washed with water four times before drying at 50 °C. K*- and Caz"- forms of
exchanged Mg-SAP were prepared in the same way, using KCI and CaCl2-2H20
as the exchange salts, respectively. T he latter products were denoted K-[Mg-
SAP] and Ca-[Mg-SAP].
2.2.4 Hybrid organo-saponite synthesis

The hybrid organo-saponite was synthesized using a method analogous to
the method for the preparation of Mg-SAP, except that hydrolyzed bissiloxane
were used exclusively as the Si source. Two types of bissiloxane were used:
bis(triethoxylsilyl)methane (BTESM) and bis(triethoxylsilyl)ethane (BTESE), and
the products were denoted C1-SAP and C2-SAP respectively. In a typical
synthesis, 1.236 g (3.6 mmol) of BTESM was prehydrolyzed in 20 mL of 0.3 M
NaOH aqueous solution at 90 °C overnight. After cooling the solution to room
temperature, 0.300 g (0.8 mmol) AI(N03)3-9H20 was mixed with the hydrolyzed
BTESM solution and stirred for 3 h, followed by the introduction of 1.200 g (20
mmol) urea and another 1.5 h of stirring. The suspension was then brought to 90
°C and heated for 1 hour before the addition-of 1.538 g (6 mmol) Mg(NO3)2-6H20.
The reaction mixture was allowed to age at 90 °C for 24 h with stirring. The

product (denoted C1-SAP) was filtered, washed with water and dried at 50 °C.

57

The CZ-SAP was prepared in the same way, expect that 1.276 g (3.6 mmol)
BTESE was used in place of BTESM.
2.2.5 Characterization

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotaflex 2008
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 ASAP 2010 sorptometer. Prior to analysis samples were
outgassed at 150 °C and 106 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 (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 W. The powdered samples were sonified in
ethanol, and the resulting suspension was dripped onto 300 mesh copper grids.

Scanning electron microscope energy dispersive X-ray microanalysis
(SEM EDS) was performed on a JEOL 6400V microscope with LaBe emitter, and
the EDS detector was an INCA x-sight. The powder samples were mounted on

aluminum stubs and coated with carbon before the analysis. The quantitative

58

analysis was performed using an INCA suite version 4.07 (Oxford Instruments)
analysis program.

2S’Si 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. The samples were spun at 4 kHz for each measurement. The pulse
delay for 29Si MAS NMR was 400 s. Talc was used as a chemical shift reference
(-98.1 ppm). The pulse delay for 27Al MAS NMR was 0.50 s. A 0.10 M aqueous
AI(N03)3 solution was used as the chemical shift reference (0.0 ppm).

FTIR spectra of samples dispersed in KBr disks were recorded at ambient
temperature on a Mattson Galaxy 3000 FTIR spectrometer over the range 400 to

4000 cm".

2.3 Results
2.3.1 Inorganic saponite derivatives

The XRD patterns of the synthetic saponite clays Mg-SAP, Zn-SAP and
SAP-200 are compared in Figure 2.2. Characteristic smectite clay in-plane
diffraction peaks are observed in all three cases. Mg-SAP did not possess a
well-stacked morphology, as verified by the absence of a 001 peak in the XRD
pattern (Figure 2.2b). However, once the powder was pressed in a hydraulic
presser at 10,000 psi for 1 min to form a pellet and subject to XRD analysis, a
weak 001 peak appeared (Figure 2.2a), indicating that the as-made Mg-SAP had

59

an irregularly stacked platy morphology. In contrast, Zn-SAP and SAP-200
exhibited a sharp 001 peak, indicating a well-stacked morphology. Zn-SAP and
SAP-200 had better crystallinity compared with Mg-SAP, as revealed by the

narrower line widths of the clay in-pIane diffraction peaks.

 

 

 

i
I 001
i
3: ' 110, 020 130.200
3‘ 150 06°
3‘ i
.7) l d
C
.03 i
C
_ i c
‘ A. b
2__'_./\ a
0 10 20 30 40 50 60 70
29(0)

Figure 2.2 Wide angle hkl X-ray diffractions of (a) pressed Mg-SAP, (b) as—made

Mg-SAP, (c) Zn-SAP and (d) SAP-200.

The platy morphologies of the synthetic saponites are clearly shown in
TEM images (Figure 2.3). The dark thin lines in the images represent clay

60

platelets with edge-on orientations to the electron beam. The length of such
edges increases dramatically in the order of Mg-SAP < Zn-SAP < SAP-200,
indicating an increase in lateral dimension and layer stacking. Also, the number
of the edges appearing in the TEM images decreases in the reverse order. The
in-plane dimensions of the Mg-SAP platelets were of the order of 50 nm,
whereas Zn-SAP had in-plane platelet dimensions of several hundred
nanometers, and SAP-200 had an even larger lateral dimensions. Layer
stacking was evident for both Zn-SAP and SAP-200. Although the layer stacking
was absent in Mg-SAP, the clay platelets were highly aggregated and interacted

with each other through edge-face interactions.

Table 2.1 The BET surface areas and pore volumes of Mg-SAP, Zn-SAP, and
SAP-200 and the Ca2*-, Cs“-, and K*- exchanged forms of Mg-SAP.

 

BET surface Total pore BJH adsorption pore

 

area (mzlg) volume (cm3/g) size (nm)
Mg-SAP 920 1.98 3
Zn-SAP 310 0.70 -
SAP-200 270 0.10 -
K-[Mg-SAP] 820 1.16 -
Cs-[Mg-SAP] 700 0.98 -
Ca-[MLSAFfl 690 1 .02

 

The relative standard deviations of the BET surface areas are less than 0.5%.

61

 

Figure 2.3 TEM images of the synthetic saponites: (a) Mg-SAP; (b) Zn-SAP and

(c) SAP-200.

62

Figure 2.4 (a) Nitrogen adsorption-desorption curves and (b) BJH adsorption
pore size distributions (off set by 0.1 for clarification) of as-made Mg-SAP and its

Ca2+-, Cs“-, and K+- exchanged derivatives.

63

(a) 1400 _

   

 

 

1200
1000
EC
*—
U) 800
23
t")
.5,
7 600
8
>
400
200 -—+—-K-[Mg-SAP]
' +Cs-[Mg-SAP]
—-o-—Ca-[Mg-SAP]
0 — I I l I I I
O 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 ‘I
P/Po
(b)0.6 -
—B—Mg-SAP
0.5 - —I—K-[Mg-SAP]

+ Cs-[Mg-SAP]
+ Ca-[Mg-SAP]

.0
h
I

 

dV/dD (cm3/g.nm)
_o o
N 00

 

 

0 5 10 15 20 25 30
Pore diameter (nm)

64

The textural properties of Mg-SAP, Zn-SAP and SAP-200 were
determined from the nitrogen adsorption—desorption isotherms. The results are
summarized in Table 2.1. Mg-SAP exhibited a large BET surface area (920
mzlg), a uniform BJH adsorption pore size (3.0 nm) and a large pore volume
(1.98 cng). The abrupt nitrogen uptake at low pressure indicates the presence
of micropores. (Figure 2.4) The well-stacked Zn-SAP and SAP-200 derivatives
had lower surface areas and pore volumes. The sodium exchange form of
naturally occurring sodium montmorillonite exhibited a surface area < 10 mzlg
and virtually no pore volume. Clearly, the exceptionally high surface area and
porosity of SAP is a manifestation ofthe relatively small platelet size and the
aggregation of the platelets in poorly stacked form. We have shown that this
unique porous morphology makes Mg-SAP a good reinforcing agent for epoxy
polymers [15].

Vibration modes typical of a 2:1 layered silicate structure were observed in
the FTIR spectrum of Mg-SAP (Figure 2.5). The broad adsorption band 3000-
3600 cm‘1 is a mixture of the stretching vibrations of OH from structural hydroxyl
groups, symmetric and asymmetric OH stretching frequencies from surface water
molecules, and gallery NH4“. The adsorption bands at 1646 and 1400 crn'1 are
assigned to the bending vibrations of OH from surface water and gallery NH4+,
respectively. The strong adsorption band at 1050 cm'1 is assigned to Si-O
stretching. The band at 476 cm'1 is assigned to Si-O-Si bending [16].

65

Transmittance (a. u.)

 

l T T l l 1
3900 3400 2900 2400 1900 1400 900 400
Wavenumver (cm’1)

Figure 2.5 The FTIR spectrum of Mg-SAP.

The 2:1 lattice structure of Mg-SAP was further characterized using solid
state NMR. The 29Si MAS NMR spectrum (Figure 2.6a) showed the presence of
two Q3 signals at chemical shifts of -87.3 and —95.2 ppm. In accord with previous
assignments [17], the -87.3 ppm line is assigned to Si centers linked through
bridging oxygen atoms to one tetrahedral Al site and two tetrahedral Si sites.
The -95.2 ppm resonance is attributed to Si linked to three adjacent tetrahedral
sites that are exclusively occupied by Si. In addition, a weak resonance near -

110 ppm is present which may be indicative of the presence of a small amount of

amorphous silica.

66

The 27Al MAS NMR spectrum (Figure 2.6b) of Mg-SAP contains two
signals at chemical shifts of 62.8 ppm and 6.5 ppm, indicating that Al substitution
occurred in both the tetrahedral and octahedral sheets of the 2:1 layered silicate
structure. Deconvolution analysis of the overlapping 27AI NMR lines at 62.8 ppm
and 6.5 ppm revealed the ratio of tetrahedral to octahedral aluminum to be 0.83 :
0.17. Due to potential differences in the electric field gradient at the quadrupolar
27AI nucleus in tetrahedral and octahedral environments, the relative intensities of
the resonances may not be proportional to the aluminum populations in these
sites. Nevertheless, the presence of octahedral and tetrahedral was clearly
indicated by the ”N NMR spectrum. Similar 298i NMR and 27AI NMR spectra
were observed for Zn-SAP and SAP-200, indicating both had a saponite-like
structure.

An SEM EDS analysis of Mg-SAP showed that the molar ratio of Si : Mg :
Al equals to 17 : 13 : 2, corresponding to a saponite unit cell formula of (Na*,
NH4*)o,5(Si7,3Alo,7)[Mgs,aAlo_2]Ozo(OH)4. The atomic ratio of Si : Zn : Al was 19 :
15 : 2 in Zn-SAP, as revealed by SEM-EDS, corresponding to a unit cell formula
of Zn-SAP was (Na*, NH4*)o,4(Si7,4Alo,6)[Zn5,aAlo,2]02o(OH)4. The atomic ratio of
Si : Mg : Al was 18 : 14 : 3 for SAP-200, and corresponding unit cell formula was

(N H4+)0.4(Si723i0.8)[M95.6A|0.4]O20(OH)4-

67

(a)

 

 

-4O -60 -80 -100 -120 -140 -160
ppm

(b)

 

g L I L

225 150 75 o -75 -150
ppm

Figure 2.6 (a) 299i MAS NMR and (b) 27AI MAS NMR spectra for Mg-SAP.

68

As-made Mg-SAP contains mixed exchange cations of Na+ and NH4“.
These gallery cations can be exchanged by various other cations. The Ca”, Cs+
and K“ exchanged derivatives of Mg-SAP were obtained by mixing Mg-SAP and
excessive corresponding metal chloride salts in water. The exchanged products
were denoted Ca-[Mg-SAP], Cs—[Mg-SAP] and K-[Mg-SAP] respectively. The
irregularly stacked morphology of Mg-SAP was maintained after the
exchange reaction, as the exchanged saponites showed XRD patterns
equivalent to Mg-SAP. (Figure 2.7) However, the 110 diffraction peak of Cs-[Mg-
SAP] was weakened. These suggest that the Cs+ cation may be bound tightly to
the hexagonal cavities in the clay basal plane, leading to an X-ray phase shift
which weakens the 110 diffraction peak. The other exchange cations could be
more heavily hydrated, and not tightly bound to hexagonal cavities. Therefore,
these cations do not affect the intensity of the 110 peak.

Nitrogen adsorption-desorption isotherms and BJH adsorption pore size
distribution curves of the saponites are shown in Figure 2.4. The exchanged Mg-
SAP clays had lower BET surface areas and total pore volumes compared with

Mg-SAP (Table 2.1).

69

   
  

130, 200
060

K-[Mg-SAP]

    
 

Cs-[Mg-SAP]

Intensity (a. u.)

Ca-[Mg-SAP]

Mg-SAP

0 10 20 30 40 50 60 70
2 theta (°)

Figure 2.7 Wide angle hkl X-ray diffractions of the Ca2*-, Cs‘-, and K‘-

exchanged Mg-SAPs, compared with that of as-made Mg-SAP.

The ion exchange of Ca2+, Cs", or K+ was verified using SEM-EDS
analysis. The atomic ratios of Si : Mg : AI of the exchanged saponites were
consistent within experimental error with the unexchanged Mg-SAP, while the
gallery Na+ cations of Mg-SAP were replaced by Ca”, Cs", and K", respectively
in the exchanged forms (Table 2.2). A lower Si ratio was observed in Cs-[Mg-
SAP], due perhaps to leeching, consistent with the XRD results. The
concentrations of the metal cations were the same in Ca-[Mg-SAP], Cs-[Mg-SAP]

70

and K-[Mg-SAP]. Mg—SAP exhibited a lower exchange cation concentration
since NH4“, as the other part of the Mg-SAP surface cations, cannot be detected

accurately using SEM-EDS.

Table 2.2 Elemental atomic ratios for Mg-SAP and its exchanged derivatives

obtained from quantitative SEM-EDS analysis.

 

 

Si ' Mg ' AF M -
Mg-SAP 19.0 14.7 2.0 0.5 Na‘
Ca-[Mg-SAP] 19.6 14.5 2.0 1.1 Ca[(OH)]*
Cs-[Mg-SAP] 18.1 15.0 2.0 1.0 Cs”
K-[Mg-SAP] 19.0 15.5 2.0 1.1 K“

 

. Each value is the average of three measurements performed on different regions of the sample.

2.3.2 Hybrid organo-saponite derivatives

We used a method analogous to that for Mg-SAP synthesis to prepare an
inorganic-organic hybrid saponite. Prehydrolyzed BTESM and BTESE
containing two triethoxylsilyl groups bridged by a methane and ethylene group,
respectively, were used as the Si sources for hybrid clay synthesis. The XRD
patterns of C1-SAP (made from BTESM) and CZ—SAP (made from BTESE) are
shown in Figure 2.8, compared with Mg-SAP. The organo saponite synthesis
was successful with BTESM, as verified by the characteristic clay in-plane
diffraction peaks in the XRD pattern of C1-SAP. The 110 peak was not well

resolved, indicating the possibility of structural distortion in the basal plane.

71

Although a weak 060 diffraction peak is present in the XRD pattern of CZ-SAP,
the principal product is a mixture of clay and hydroxides of magnesium and
aluminum, indicating that BTESE was not completely converted to a clay
structure due to the unsuitable bulky size of the ethylene bridging group.

C1-SAP had a BET surface area of 656 m2/g. The presence of
micropores on C1-SAP surfaces was verified by the uptake of nitrogen at low
pressure in the nitrogen adsorption-desorption isotherm curve (Figure 2.9),
similar as in the isotherm curve of the inorganic Mg-SAP. However, the
mesoporous region (0.1 ~ 1 P/Po) of C1-SAP was dramatically reduced, resulting

in a lower total pore volume of 0.54 cm3/g, compared with 1.98 cm3/g for Mg-SAP.

\
I

3: 1
g 1
a? r
g I CZ-SAP
2 1
E

, 110. 020 C1-SAP

130,200
060
Mg-SAP
0 10 20 30 40 50 60 70
2theta(°)

Figure 2.8 XRD patterns of Mg-SAP, C1-SAP and CZ-SAP reaction products.

72

400 c

350 -

Veda (cmalg, STP)
N
O
O

 

 

 

150
100
50
0 I I I I I J J I I I
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1
P/Po

Figure 2.9 The nitrogen adsorption-desorption isotherms for C1 -SAP.

The incorporation of methylene group in the structure of C1-SAP was
verified using 29Si, ‘30, and 27Al MAS NMR (Figure 2.10) and FTIR. In the 2"Si
NMR spectrum, two chemical shifts at -55 ppm and -90 ppm were observed, and
were assigned to T2 and ()2 silicon species respectively. The presence of
methylene group and C-Si bonds in C1-SAP was verified by the T bands in the
”Si NMR spectrum, as well as by the chemical shift at -3 ppm in the 130 NMR
spectrum. The presence of 02 silicon species in the C1-SAP structure was
possibly due to the cleavage of C-Si bonds during BTESM hydrolysis. Details

will be discussed in the Discussion part. lsomorphous substitution by Al in both

73

the tetrahedral and octahedral centers of the C1-SAP structure was verified by
the presence of tetrahedral and octahedral Al resonances at 55 and 4 ppm,
respectively, in the 27AI NMR spectrum (Figure 2.10c).

The FTIR spectrum of C1-SAP exhibited characteristic clay in-plane lattice
vibration modes that were comparable with those of Mg-SAP (Figure 2.11). The
new bands at 2980 and 1275 cm'1 in the spectrum of C1-SAP are assigned to C-
H and Si-C stretching bands, respectively. C1-SAP has a platy morphology with
lateral dimensions in the range of 20 ~ 50 nm, as shown in its TEM image (Figure
2.12). Unlike Mg-SAP, C1-SAP did not exhibit an aggregated nanolayer
morphology. CHN elemental analysis revealed that C1-SAP contained around

4.7 wt% carbon.

74

Figure 2.10 The (a) 293i MAS NMR, (b) ‘30 MAS NMR and (c) 27Al MAS NMR

spectra of C1 -SAP.

75

 

 

 

 

40 0 -40 -80 -120 -160 -200
ppm
(b)
100 50 0 -50 -100
ppm
(C)
250 1 50 50 -50 -1 50
ppm

76

Transmittance (a. u.)

 

 

-Vi_ - 21‘,_l___. ...--7 4

3900 3400 2900 2400 1900 1400 900 400
Wavenumber (cm'1)

Figure 2.11 FTIR spectrum of C1-SAP. The circled regions in the spectrum of

C1-SAP indicate the incorporation of methylene groups in C1-SAP structure.

 

Figure 2.12 TEM images of C1-SAP showing its platy morphology.

77

2.4 Discussion
2.4.1 Inorganic saponite derivatives

The physical properties of the synthetic inorganic clay products are
directly determined by synthesis conditions, particularly the synthesis
temperature and the composition of the octahedral sheet. Comparing Mg-SAP
made at 90 °C and SAP-200 made at 200 °C, the higher synthesis temperature
and a longer crystallization time for SAP-200 resulted in ordered layer stacking
and larger in-plane platelet dimensions. Consequently, SAP-200 exhibited a well
stacked layered morphology, while the stacking was not ordered for Mg—SAP.
The octahedral species also plays important role in determining clay particle
morphology determination. With Zn as the primary octahedral species, Zn-SAP
had a faster platelets octahedral s heet g rowth rate, which afforded larger clay
platelets dimensions and better layer stacking. This fact also indicates that the
clay stacking formation is dependent on the growth rate of the octahedral sheet.

Synthetic saponites can be subject to ion-exchange reaction. Although K“,
05", Ca”, Na+ and NH4+ have different atomic diameters and could result in
different gallery spacings, the irregularly stacked morphology of Mg-SAP was not
affected by the ion exchange reactions. The exchanged Mg-SAP exhibited
surface porosities, mainly micropores, as verified by the abrupt nitrogen uptake in
the low partial pressure region of the N2 adsorption-desorption isotherms. (Figure

2.4) The mesoporous volumes forthe exchanged Mg-SAPs were lower than the

78

pore volume of as-made Mg-SAP, as judged by the lower nitrogen uptake in the
range of 0.1 ~ 1.0 P/Po. The BJH pore size distribution curves also revealed that
exchanged forms of Mg—SAP had smaller pore sizes than Mg-SAP. The
decrease in surface areas and pore volumes of the exchanged saponites might
be due to the collapse of the aggregated clay layers upon replacement of
mixtures of Na+ and NH4+ cations by Ca”, Cs+ or K“.
2.4.2 Hybrid organo-saponite derivatives

The clay-like structure has been verified for C1-SAP based on the
characteristic smectic clay diffraction peaks in the XRD pattern. The
incorporation of methylene group as groups between silica tetrahedra in the
basal plane was verified by the 29Si and ‘30 MAS NMR and FTIR. The size of
the organo bridging group in bissiloxane, as the starting reagent for clay
synthesis, has to be suitable for the replacement of oxygen in the Kagome
network. Bulky groups, such as the ethylene group in BTESE, can not substitute
for oxygen and cannot lead to the formation of the hybrid clay structure.
However, the methylene group is isoelectronic with oxygen and well suited as a
replacement for bridging oxygen.

The surface porosity was much reduced in C1-SAP, compared with

inorganic Mg-SAP. The surface porosity of Mg-SAP arises from the formation of
an irregularly aggregated layer morphology. However, the possible presence of

organo groups on the edges of C1-SAP platelets weakened the edge-face

79

interactions. Therefore, the platelets correlated with each other without leaving
inter-platelet pores. Some layer stacking was observed in the TEM image
(Figure 2.11). However, a stacking peak was not resolved in the XRD pattern of
C1-SAP due to the low stacking order and small domain sizes of the platelets.
Despite the incorporation of methylene groups in C1-SAP basal plane,
some structural distortions occurred, compared with an ideal smectic clay
structure, as evidenced by the 02 silicon species in the 2"Si NMR and the
absence of a 110 diffraction peak in the XRD pattern. There are two main
causes for the structural distortions. First, since BTESM was used exclusively as
the Si source to prepare C1-SAP, there was a high possibility that one of the
siloxanes is in the clay basal plane while the other is inverted and oriented out of
the basal plane. Using a mixture of an inorganic Si source, such as a sodium
Silicate solution, and BTESM for clay synthesis may solve this problem, although
the products would have a lower methylene loading. Second, BTESM was
susceptible to nucleophilic substitution under basic conditions, and would yield
inorganic Si—(OH)4 species and intermediate Si-CH2' species [13]. The hydrolysis
reaction introduces 02 Si species and methyl terminal groups into the Kagome
network. Although the prehydrolysis of bissiloxane has been used in the
preparation of hybrid zeolite materials, the hydrolysis conditions, such as pH and
reaction temperature, have to be adjusted for hybrid clay synthesis to avoid the
C-Si bond cleavage and to minimize clay structural distortion. It is expected that,

80

by using a mixture of sodium silicate solution and BTESM as the Si source, the
prehydrolysis step of BTESM can be removed, and the hydrolysis of BTESM will

take place in the basic sodium silicate solution even at room temperature.

2.5 Conclusion

Synthetic saponites have been prepared at 90 °C and under hydrothermal
conditions. The products shared a similar smectic clay structure, while their layer
stacking order, surface porosity varied according to the synthesis conditions and
octahedral composition. A higher synthesis temperature and rapid octahedral
Sheet formation leads to better crystallinity and layer stacking in the products. An
analogous synthesis method was used to make organo clay hybrids using
exclusively BTESM as the Si source so that the hybrid product would have
methylene group bridging silica tetrahedra in the clay basal plane. Although
some structural distortions still occurred, this strategy provides a novel approach
to achieve organic decoration of a clay structure from a one-pot sol-gel synthesis
while preserving the basic clay lattice structure. Potential applications of the

synthetic hybrid clay include polymer reinforcement, catalysts, absorbance etc.

81

References

(1) Brindley, G. W.; Brown, G. Crystal Structures of Clay Minerals and Their X-
ray Identification; Mineralogical Society: London, 1980.

(2) Murray, H. H. Clay Miner. 1999, 34, 39.

(3) Vogels, R. J. M. J.; Kerkhoffs, M. J. H. V.; Geus, J. W. Stud. Surf. Sci. Catal.
1995.91.1153.

(4) Kloprogge, J.T.; Breukelaar, J.; Jansen, J. B. H.; Geus, J. W. Clays Clay
Miner. 1993, 41, 103.

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

(6) Reinholdt, M.; Miehe—Brendle, J.; Delmotte, L.; Tuilier, M.-H.; le Dred, R.;
Cortes, R.; Flank, A.-M. Eur. J. Inorg. Chem. 2001, 11,2831.

(7) Carrado, K. A. Appl. Clay Sci. 2000, 17, 1.

(8) Prihod'ko, R.; Sychev, M.; Hensen, E. J. M.; van Veen, J. A. R.; van Santen,
R. A. Stud. Surf. Sci. Cafe]. 2002, 142, 271.

(9) Chastek, T. T.; Que, E. L.; Shore, J. S.; Lowy, R. J.; Macosko, C.; Stein, A.
Polymer 2005, 46, 4421 .

(10) Chastek, T. T.; Stein, A.; Macosko, C. Polymer 2005, 46,4431.
(11)Whilton, N. T.; Burkett, S. L.; Mann, S. J. Mater. Chem. 1998, 8, 1927.
(12) Burkett, S. L.; Press, A.; Mann, 8. Chem. Mater. 1997, 9, 1071.

(13) Yamamoto, K.; Sakata, Y.; Nohara, Y.; Takahashi, Y.; Tatsumi, T. Science
2003, 300, 470.

(14) Hatton, B.; Landskron, K.; Whitnall, W.; Perovic, D.; Ozin, G. A. Acc. Chem.
Res. 2005, 38, 305.

82

(15) Xue, S.; Pinnavaia, T. J. Microporous Mesoporous Mater. 2007, doi:
10.1016/j.micromeso. 2007.02.042.

(16) Madejova, J.; Komadel, P. Clays Clay Miner. 2001 , 49, 410.

(17) Lipsicas, M.; Raythatha, R. H.; Pinnavaia, T. J.; Johnson, I. D.; Giese Jr, R.
F.; Costanzo, P. M.; Robert, J.-L. Nature, 1984, 309, 604.

83

CHAPTER 3

Porous Synthetic Smectic Clay for the Reinforcement of Epoxy Polymers

3.1 Introductlon

Polymer — smectic clay nanocomposites, first demonstrated for a Nylon-6
polymer [1], have attracted much research interest over the past decade. When
the clay nanolayers are fully dispersed in the polymer matrix, significant
improvements in mechanical strength, thermal stability, barrier properties, and
other properties can be realized [2-5] at low clay loadings (< 10 wt%). Naturally
occurring smectite clays have poor wetting properties when combined with
polymer or polymer precursor due to incompatible surface polarity. Thus, when
unmodified clays are combined with a polymer, the nanolayers retain a stacked
structure, resulting in the formation of a conventional composite containing clay
tactoids in micrometer size. In order to achieve exfoliation of the clay nanolayers
in the polymer matrix, it is almost always necessary to replace the inorganic
exchange cations on the clay basal surfaces with alkylammonium or other
organic cations. The organocations enlarged the galley space between stacked
nanolayers, lower the polarity of the surface and allow for the intercalation of

polymer between nanolayers. Under appropriate though often stringent [6]

84

processing conditions, complete exfoliation of the nanolayers into the polymer
matrix can be achieved.

Although natural smectic clay minerals are abundant and inexpensive,
they require extensive purification for nanocomposite applications. On the other
hand, they also can be synthesized through sol—gel methods under hydrothermal
conditions [7-12]. Organosilicates with lamellar structures analogous to smectic
clays also have been prepared using one-pot sol-gel methods [11, 13, 14]. In
addition, the composition and surface polarity of synthetic clays can be easily
modified through a judicious choice of cations in the tetrahedral and octahedral
interstacies of the layers.

In the present work we investigate the properties of a synthetic smectic
clay (saponite) for the reinforcement epoxy polymers. The obtained clay forms a
2D micrometric network of irregularly stacked nanolayers approximately 50 nm in
diameter. Sonification reduces the domain Size of the aggregated platelet
network and facilitates the dispersion of clay particles in the polymer. Moreover,
the dispersion of clay aggregates is achieved without the need for organo-
mod'rfication of the clay surfaces through ion exchange with alkylammonium ions.
Thus, it is possible to achieve polymer reinforcement while avoiding the
plasticizing effects of the alkylammonium ions [15, 16] and the complications
caused by Hoffman degradation of such ions at temperatures above 200 °C [17,
18]. Although epoxy — clay nanocomposites have been extensively studied [6,

85

15, 19-24], improving the mechanical strength of glassy epoxy derivatives
remains a challenge. We show that the synthetic clay of the present study
substantially improves the tensile properties of both rubbery and glassy epoxy

matrices.

3.2 Experimental
3.2.1 Materials

The diglycidyl ethers of bisphenol A (DGEBA), more specifically, Shell
EPON resin 826 and 828, and polyoxypropylene diamines of the type
NH2CH(CH3)CH2[OCH20H(CH3)]XNH2 (Huntsman Chemicals Jeffamine D230
with x=2.6 and D2000 with x=33) were used to form glassy and rubbery epoxy
matrices. PGV Na*-montmorillonite (NaMMT) was provided by Nanocor Inc. All
other chemicals were purchased from Aldrich Chemical Co. and used without
further purification.
3.2.2 Saponite synthesls

Synthetic saponite was prepared at 90 °C according to previously

described methods [9] using water glass solution (27 wt. % silica, 14 wt. %
NaOH), AI(N03)3-9H20, Mg(N03)2-6H20 and urea as the source of base in the
molar ratio 3.6 : 0.40 : 3.0 : 10 per 400 moles of water. After a crystallization
period of 24 h the as-made NH“, Na+ mixed ion form of the clay was converted

to the sodium ion form by ion exchange and denoted SAP. An organic ion

86

derivative, denoted C16-SAP, was prepared by ion exchange reaction of SAP
with cetyltrimethylammonium bromide (CTAB).
3.2.3 Composite preparation

SAP was used to reinforce both glassy and rubbery epoxies, prepared
from stoichiometric amounts of EPON 826 - Jeffamine 0230 (weight ratio 76 : 24)
and EPON 828 - Jeffamine D2000 (weight ratio 27.5 : 72.5), respectively. For
comparison purposes, sodium montmorillonite (NaMMT) and C16-SAP clays
were also used to make glassy epoxy composites.

Glassy epoxy composites were obtained by first suspending the clay in
ethylacetate and subjecting the suspension to ultrasonification for 10 min using a
Branson 1020 digital laboratory sonifier. The EPON resin was added to the
sonified clay suspension and the mixture was stirred in a fume hood overnight
and then heated at 50 °C for 4 h under vacuum to complete the evaporation of
solvent. T hen the curing agent was added to the epoxy-clay mixture and the
mixture was stirred at 75 °C for 10 min, outgassed at room temperature for 20
min, and then poured into silicone molds to obtain tensile testing specimens. The
composites were partially cured at 75 °C for 3 h and then fully cured at 125 °C for
3 h.

Rubbery composites were prepared in the same manner as the glassy
composites except that the clay, epoxy resin, and curing agent were directly
mixed without the aid of ethylacetate or ultra-sonication.

87

3.2.4 Characterization

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 s ample holder. R ectangular f lat s pecimens of c ross-Iinked
composites were mounted onto an aluminum holder for XRD examination.

Transmission electron microscopy (T EM) 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. For composites, thin sections
(~100 nm thick) of the composites were put on naked copper grids.

Tensile mechanical 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,
and 12.5 mm/min for rubbery specimens. The reported tensile values are the
averages for 4-5 specimens with dog-bone shapes. The standard deviation for
the modulus and strength was +/- 3%. The standard deviation for the toughness
values, obtained by integrating the area under the stress - strain curves, was +/-

35%.

88

Dynamic mechanical analysis (DMA) was performed on a DMA 2980
dynamic mechanic analyzer (TA Instruments) operated in three-point bending
mode at a frequency of 1Hz and amplitude of40 pm, from 25 to 1 40 °C at a
heating rate of 4.0 °C/min. The specimens were rectangular bars with

dimensions 60 X 13 X 3 mm.

3.3 Results
3.3.1 Epoxy — saponite nanocomposites

The synthetic SAP has an irregularly stacked platy morphology.
Consequently, it exhibited a large BET surface area (920 m2/g) and a average
BJH pore size of 5.0 nm and a pore volume of 1.98 cm3/g. Detailed
characterization of SAP is described in Chapter 2.

Both SAP and NaMMT were tested for the reinforcement of epoxy
polymers, particularly in the glassy state, without the use of organic surface
modifier. Rapid settling-out of clay particles was observed for suspensions of
NaMMT in the epoxy resin, curing agent, and stoichiometric mixtures of the resin
and curing agent, making it impossible to form uniform composites. However,
little or no particle settling occurred for SAP clay suspensions, as judged by the
depth profiling of the XRD intensity of the clay 060 reflection for cured

composites (Figure 3.1).

89

 

\
'1

topside

bottom

Intensity (a. u

 

l

65 70

 

 

50 55 60
2 theta (°)

Figure 3.1 The XRD patterns of the topside and bottom of an epoxy - SAP 5.4%

composite, both showing 060 clay diffraction peak of similar intensity.

 

Figure 3.2 A TEM image of thin—sections of a glassy epoxy — SAP 5.4 wt%

composite.

90

TEM images of thin-sections of a 5.4 wt% epoxy - SAP composite
showed a uniform dispersion ofaggregated SAP particles (Figure 3.2). The
edges of randomly oriented clay platelets with lateral dimensions ~ 50 nm are
cleariy observed in the aggregates. Clay platelets lying flat in the sections are
visible as lighter regions of the image due to the low contrast. The particle
aggregates are comprised of face-face stacks of nanolayers and edge-face
interactions between the stacks. There is little evidence for the presence of
single, delaminated platelets within the aggregates. Thus, the absence of a well-
resolved 001 stacking reflection in the XRD pattern of SAP and SAP-epoxy
composites is the result of irregular stacking of nanolayers and not the result of
nanolayer exfoliation. Although bent clay nanolayers are sometimes observed
for polymer — clay nanocomposites [25, 26], little or no layer bending is observed
for epoxy — SAP nanocomposites due primarily to the small lateral dimensions of
the nanolayers.

3.3.2 Composite Properties

The tensile properties of glassy epoxy - SAP and - NaMMT composites
are provided in Table 3.1. The improvement in the tensile properties of the SAP
composites is correlated with the exceptional dispersion of these particles in the
polymer matrix. On the other hand, NaMMT exhibited poor dispersion in the
epoxy matrix, and the corresponding composites exhibit mechanical properties
typical of conventional composites. Although the modulus is improved with the

91

addition of NaMMT, the tensile strength and elongation-at—break decreased
dramatically with increasing clay loading. In contrast, the SAP composites
showed improved tensile strength and elongation-at—break, in addition to
improved modulus. The increase in the elongation-at—break substantially

improves toughness. For example, a 45% increase in toughness was achieved

at a 9.4 wt% SAP loading.

Table 3.1 Tensile properties of glassy epoxy composites reinforced by SAP and

 

 

NaMMT clay nanoparticles.
Clay Clay Tensile Tensile Elongation- Toughness
Loadinga strength modulus at-break (MJ/m3)
(wt%) (MPa) (GPa) (%)

Pristine 0.0 66.1 2.9 4.3 2.1

SAP 2.1 67.8 3.1 5.7 3.0

5.4 70.2 3.1 8.9 4.9

9.4 72.5 3.8 5.4 3.1

NaMMT 7.2 34.8 5.1 0.8 0.2

12.4 29.4 5.2 0.6 0.1

 

a The clay loading is on a silicate basis; The standard deviations for all values is ~ 3%, except for
the values of elongation-at-break and toughness for the SAP composites in which case the
standard deviation is ~30%.

Representative strain-stress curves for epoxy - SAP nanocomposites are
provided in Figure 3.3. The pristine epoxy showed a typical failure behavior for a
brittle material. The stress increased with increasing loading to a maximum
value, after which the specimen ruptured with little yield. Similar behavior was

observed for the nanocomposites, except that the initial slopes were larger due to

92

the increased modulus and the stress maximum was increased due to increased
strength. Also, the composites exhibited substantial yield and an increase in
elongation-at—break, resulting in increased toughness. Although the standard
deviations for the modulus and strength was less than 3%, the magnitude of the
break at elongation was much larger (> 30%) and most likely dependent on local

defects in the composites.

80-

Stress (MPa)

 

 

 

Strain (%)

Figure 3.3 Representative stress - strain curves for glassy epoxy - SAP clay

nanocomposites at the wt % clay loadings indicated.

DMA also showed improved storage modulus for the glassy epoxy - SAP

nanocomposites, compared with pristine epoxy (Table 3.2). Additionally, the

93

glass transition temperatures of the composites are comparable to the pristine
epoxy. However, the onium ion modified clay formed by ion exchange with
cetyltrimethylammonium cations, denoted C16—SAP, which contains about 20
wt% organic components as determined from elemental analysis, causes the
glass transition temperature of the polymer to decrease due to the plasticizing

effect of the modifiers.

Table 3.2 Dynamic mechanical analysis (40 °C) of glassy epoxy composites filled

with SAP and C16-SAP clay.

 

Clay Loadinga (wt%) T9b (°C) Storage Modulus (GPa)

 

Pristine 0.0 87.2 2.7
SAP 5.2 86.9 3.0
10.3 87.6 3.3

C16-SAP 4.1 84.2 2.9
8.6 84.0 3.2

 

3 Clay loading is on a silicate basis. b Estimated from tan fimax.

AS expected for the dispersion of rigid inorganic particles in a soft polymer
matrix, far greater improvements in tensile strength, modulus and elongation-at-
break were observed for rubbery epoxy - SAP nanocomposites (Table 3.3). A
730% increase in strength, 360% increase in modulus, and 86% increase in

elongation-at—break was observed at a loading of 15.0 wt% SAP.

94

Table 3.3 Tensile properties of rubbery epoxy composites reinforced by SAP clay

nanoparticles.

 

Clay Loadinga Tensile strength Tensile modulus Elongation-at-

 

(wt%) (MPa) (MPa) break (%)
0 0.6 2.8 24.9
2.1 1.2 3.8 38.7
5.5 1.8 5.6 38.0
10.9 2.7 7.7 39.2
15.0 5.0 12.9 46.4

 

3 Clay loading is on a silicate basis. The standard deviations for the values of strength. modulus,
and elongation-at-break are <5 %.

 

 

'..‘ CD” “7‘ "
b: 1:19.24: at"
(fir—91390914 ”O
(13")ctli>~.moigi.
QT" ”0 Cia’bfi"
t- “O 0 ‘.
LiniUHEp-I‘ttg
“clclo‘ itdlmi

Figure 3.4 Pictures of (a) a pristine glassy epoxy polymer and (b) the

corresponding epoxy - SAP nanocomposite containing 5.4 wt% clay.

95

Another advantage of synthetic SAP clay as a reinforcement agent is the
optical transparency of the resulting composites (Figure 3.4). In contrast, the iron
impurities in naturally occurring Na-MMT clay imparts an off-white color to the

composites, along with reduced transparency [27].

3.4 Discussion

The irregular platelet stacking and associated high surface area of SAP
allows for the reinforcement of both rubbery and glassy epoxy matrices without
the need for organic modification of the basal surfaces (Tables 3.1 and 3.3). On
the other hand, NaMMT provide improvements in modulus, but suffer from a
severe loss of strength due to the non-uniform dispersion of aggregates in the
polymer matrix. Although organic modification of the SAP basal surfaces by ion
exchange with cetylammonium cations facilitates platelet dispersion and provides
reinforcement, the glass transition temperature of the polymer is compromised
due to the plastisizing effect of the modifier (Table 3.2). This is only one of
several reasons why eliminating the use of an organic modifier in the formation of
polymer nanocomposites can be advantageous.

The dispersion of SAP nanoparticles into glassy and rubbery epoxy
matrices occurs with the retention of nanolayer aggregated in the form of tactoids
(Figure 3.2). That is, the presence of single nanolayers in the polymer matrix is
virtually non-existent. However, the tactoids are uniformly dispersed throughout

96

the polymer matrix. It has been previously observed that aggregated clay
tactoids in glassy epoxy matrices typically improve modulus, but sacrifice
strength and toughness, especially at high clay loadings [21, 24, 26]. The
fracture mechanism in this case was attributed to the weak interaction between
the stacked nanolayers and the development of microcracks between the
stacked nanolayers. However, for the SAP-epoxy composites obtained in the
present work, the clay nanolayers are not regularly stacked and the disordering
of the aggregated nanolayers allows the polymer to fill at least some of the pores
between nanolayers, thereby reducing the tendency toward microcrack formation.
Thus, epoxy - SAP nanocomposites can approach the tensile properties of
composites containing fully delaminated organoclay nanolayers [22].

The ability of exfoliated clay nanolayers to improve the strength and
toughness of a rubbery epoxy matrix has been attributed in part to the ability of
the single nanolayers to align under applied strain. It also has been suggested
[28] that for the rubbery domain above To, the clay layers have better mobility
than below T9, leading to improvements in strength, as well as modulus. On the
basis of the present work it can now be said that disordered clay tactoids provide
strength and toughness analogous to exfoliated organoclay nanolayers.
Moreover, the reinforcement is extended for the first time to glassy, as well as
rubbery epoxy matrices. In fact, the present glassy system is the only reported
example for which the modulus, strength and toughness all are improved without

97

sacrificing the glass transition temperature. Thus, disordered SAP clay tactoids
are made even more appealing as reinforcing agents for epoxy matrices, not only
because they avoid the undesirable plasticizing effects of an organic modifier, but
also because they benefit both rubbery and glassy matrices and minimize costs
by eliminating the need for an organic modifier. Still further, these benefits can

be realized with relatively little loss in optical transparency (Figure 3.4).

3.5 Conclusion

A porous synthetic saponite-like material (denoted SAP) was used to
reinforce rubbery and glassy epoxy polymers. Remarkably, reinforcement
properties superior to those of organo-montmorillonite can be achieved without
the need for a surface organic modifier, as evidenced by the improvements in
tensile strength, modulus and toughness at temperature above and below the
glass transition temperature of the polymer. The disordered stacking of clay
platelets ~50 nm in lateral dimension gives rise to aggregated tactoids with a
BET surface area of 920 mzlg, an average BJH pore size of 5.0 nm, and a pore
volume of 1.98 cm3/g. These unique textural properties facilitate the dispersion of
the tactoids into the polymer matrix and allow reinforcement of the polymer

matrix in the absence of an organic modifier.

98

References

(1) Usuki, A.; Kojima, Y.; Kawasumi, M.; Okada, A.; Fukushima, Y.; Kurauchi, T.;
Kamigaito, O. J. Mater. Res. 1993, 8, 1179.

(2) Ray, S. S.; Okamoto, M. ng. Polym. Sci. 2003, 28, 1539.

(3) LeBaron, P. C.; Wang, Z.; Pinnavaia, T. J. Appl. Clay Sci. 1999, 15, 11.
(4) Becker, 0.; Simon, G. P. Adv. Polym. Sci. 2005, 179, 29.

(5) Lloyd, S. M.; Lave, L. B. Environ. Sci. Technol. 2003, 37, 3458.

(6) Kommann, X.; Lindberg, H.; Berglund, L. A. Polymer 2001, 42, 1303.

(7) Kloprogge, J.T.; Breukelaar, J.; Jansen, J. B. H.; Geus, J. W. Clays Clay
Miner. 1993, 41, 103.

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

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

(10) Reinholdt, M.; Miehe-Brendle, J.; Delmotte, L.; Tuilier, M.-H.; le Dred, R.;
Cortes, R.; Flank, A.-M. Eur. J. Inorg. Chem. 2001, 11,2831.

(11) Carrado, K. A. Appl. Clay Sci. 2000, 17, 1.

(12) Prihod'ko, R.; Sychev, M.; Hensen, E. J. M.; van Veen, J. A. R.; van Santen,
R. A. Stud. Surf. Sci. Catal. 2002, 142,271.

(13) Chastek, T. T.; Que, E. L.; Shore, J. S.; Lowy, R. J.; Macosko, C.; Stein, A.
Polymer 2005, 46, 4421 .

(14) Chastek, T. T.; Stein, A.; Macosko, C. Polymer, 2005, 46,4431.

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

99

(16) Park, J.; Jana, S. C. Macromolecules 2003, 36,8391.

(17) Zanetti, M.; Camino, G.; Reichert, P.; Mulhaupt, R. Macromol. Rapid
Commun. 2001, 22, 176.

(18) Xie, W.; Gao, Z. M.; Pan, W. P.; Hunter, 0.; Singh, A.; Vaia, R. Chem. Mater.
2001, 13, 2979.

(19) Wang, K.; Chen, L.; Wu, J.; Toh, M. L.; He, 0.; Yee, A. F. Macromolecules
2005, 38, 788.

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

(21) Miyagawa, H.; Foo, K. H.; Daniel, I. M.; Drzal, L. T. J. Appl. Polym. Sci. 2005,
96, 281.

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

(23) Lan, T.; Kaviratna, P. D.; Pinnavaia, T. J. Chem. Mater. 1995, 7,2144.
(24) Park, J. H.; Jana, S. C. Macromolecules 2003, 36,2758.

(25) Zilg, C.; Mulhaupt, R.; Finter, J. Macromol. Chem. Phys. 1999, 200,661.
(26) Wang, K.; Wang, L.; Wu, J.; Chen, L.; He, C. Langmuir 2005, 21 , 3613.

(27) Park, l.; Peng, H.-G.; Gidley, D. W.; Xue, S.; Pinnavaia, T. J. Chem. Mater.
2006, 18,650.

(28) Shah, D.; Maiti, P.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. Adv. Mater.
2005, 17, 525.

100

CHAPTER 4
PALYGORSKITE AS AN EPOXY POLYMER REINFORCING AGENT

4.1 Introduction

Montmorillonite and other members of the smectite family of clay minerals
have been extensively studied as polymer reinforcement agents [1 -4]. In order to
make these intrinsically hydrophilic clay compatible with polymer matrices,
onium-ion exchange reactions normally are used to mediate the surface polarity.
When the organoclay derivatives are dispersed as individual 1-nm thick sheets in
a polymer matrix, resulting nanocomposites usually exhibit excellent mechanical
properties, thermo-stability, barrier property, fire retardant, and other material
properties.

Palygorskite and sepiolite have structures related to the 2:1 layered
structure of smectite clays, except that the layers are pleated to form cross-linked
ribbons [5,6]. This cross-linking feature precludes the possibility of exfoliating the
structure into 1-nm thick nanoparticles. Nevertheless, owing to the relatively high
surface area and lath morphology of these minerals, particularly in the case of
palygorskite, they have found use as absorbents, pigments, catalysts, and
rheological control agents [5, 7-10]. However, palygorskite has received
relatively little attention as a polymer reinforcement agent with only a few studies

101

being devoted to composites of polyolefins, polyurethane, polyimide and
polyamide [11-19].

In the present work we investigate the properties of pristine and silylated
palygorskite for the reinforcement of rubbery and glassy epoxy polymers.
Silylation has been used previously to modify the surface polarity of palygorskite
[20], as well as other mineral [21-23] and synthetic silicates [24] with potential
polymer reinforcement properties. Silylation also is used in the present work to
enhance the hydrophobic character of palygorskite and to improve its dispersion

in an epoxy matrix.

4.2 Experimental
4.2.1 Materials

The commercially available diglycidyl ethers of bisphenol A (DGEBA) with
the structure shown in Scheme 4-1 were provided by the Shell company under
the trade names Epon 828 [n=0 (88%), 1 (10%), 2 (2%)] and Epon 826 [n=0

(99%), n=1 (1%)]:

/°\ 9'11 9" ‘3'“ /°\
CHz—CH-CHz o~.cl:<.o-cu,-cu-cn2 o O (I: Q O-CHz-CH—CH,
CH3 n CH3

(4-1)
The curing agents, shown in Scheme 4-2, were 0.01-polyoxypropylene

diamines provided by Huntsman Chemicals under the trade names Jeffamine

102

02000 [x=33.2, MW~2000] and Jeffamine D230 [x=2.6, MW~230] or the glassy

system respectively (Scheme 4-2).
(EH3 (EH3
HZN-CH-cuzio-CHz-CHithl-I,
(4-2)

Palygorskite was provided by the Active Minerals Company LLC under the
trade name Acti-Gel 208. Other chemicals were purchased from Aldrich
Chemical Co. and were used as received.
4.2.2 Palygorskite purification

The as-received palygorskite was found to contain small amounts of
mineral impurities. A purification process was carried out by preparing a 10 wt%
aqueous suspension of the mineral and allowing the more dense impurities
(quartz and carbonates) to sediment out. After a certain sedimentation time, the
supernatant suspension was decanted off and the sedimentation process
continued. X-ray diffraction analysis of each sediment fraction provided an
indication of the purity of the mineral remaining in suspension. The purification
was repeated until there was no evidence by X-ray diffraction for the presence of
quartz or carbonate minerals in the sediment. The purified mineral remaining in
suspension then was collected and used to form nanocomposites for comparison
with composites made from as-received palygorskite. Survey experiments

revealed no difference in the mechanical properties of epoxy composites

103

prepared from the purified and as-received palygorskite. Accordingly, for use in
all subsequent experiments, the as-received palygorskite was dispersed in
acetone and subjected to sonification for 30 min. The suspension was
evaporated in air to collect the mineral.
4.2.3 Palygorskite silylation

The silylation of palygorskite was carried out in a mixture of the clay and
the silylation reagent in toluene under reflux conditions. A 1- to 2-g amount of
palygorskite and a stoichiometric amount of silylation agent were added to 50 mL
of toluene. The amounts of silylation agent used were determined by assuming
there were 4 silanol sites per nm2 of clay surface area. The mixture was
submitted to ultrasonication for 30 min and then the suspension was heated
under reflux for 4 h. The resulting silylated product was air-dried, ground in a
mortar and pestle, and used for nanocomposite formation. The silylation
reagents used were y-aminopropyltrimethoxysilane (APTMS), N-
dodecyltriethoxysilane (DTES), and 1,1,1,3,3,3—hexamethyldisilazane (HMSZ).
The corresponding silylated palygorskite (PLG) clays were denoted APTMS-PLG,
DTES-PLG, and HMSZ-PLG, respectively.
4.2.4 Composites Preparation.

Both as-received and silylated palygorskite materials were used to make
rubbery or glassy epoxy composites. Equivalent amounts of epoxy monomer
and curing agent were mixed at 75 °C for 10 min, out-gassed at room

104

temperature for 20 min, and transferred into a dog-bone shaped aluminum mold
for curing. The composites were partially cured at 75°C for 3 h and then fully
cured at 125°C for 3 h under nitrogen flow.
4.2.5 Characterization Methods

X-ray diffraction (XRD) patterns were obtained on a Rigaku rotafiex 2003
diffractometer equipped with Cu Ko X-ray radiation and a curved crystal graphite
monochromator, operating at 45 KV and 100 mA. Clay films for X-ray diffraction
analysis were prepared by dripping droplets of aqueous clay suspension onto a
glass slide and allowing the suspensions to dry. The diffraction patterns were
recorded from 1 to 70 ° 29 with a step scan of 0.02 °lpoint and a scan rate of 0.5
°lmin. Cured composite samples were prepared by fitting rectangular flat
specimens into the windows of aluminium holders. The pattems were recorded
from 1 to 40 ° 28 with a step interval of 0.05 °lpoint and a scan rate of 2 °lmin.

TEM images were obtained on a JEOL 2010F 200kV field emission TEM
with an acceleration voltage of 200 W. The clay samples for TEM imaging were
prepared by evaporating clay-ethanol suspensions on a holy carbon coated
copper grid. Composites samples for TEM imaging are in the form of 80~100 nm
thick thin sections supported on a copper grid.

The tensile measurements on dog-bone shaped samples were carried out
at room temperature according to ASTM procedure D3039 using an SFM-20
United Testing System equipped with a laser extensometer. The measurements

105

were conducted with a crosshead speed of 25 mm/min for rubbery samples and
0.5 mm/min for glassy samples. The reported tensile parameters are values
averaged over four independent specimens.

FTIR spectra of samples dispersed in KBr disks were recorded at ambient
temperature on a Mattson Galaxy 3000 FTIR spectrometer over the range 400 to
4000 cm". 29Si MAS NMR spectra were obtained at 79.4 MHz on a Varian VXR-
4008 solid state NMR spectrometer equipped with a magic angle spinning probe.
For each measurement the sample was spun at 4 kHz. The pulse delay was 400
s, and chemical shifts were referenced to talc. Therrnogravimetric analyses
(TGA) were carried out on a Cahn TG System 121 Analyser. The powdered clay

samples were heated to 800 °C at a rate of 5 °C/min.

4.3 Results
4.3.1 Characterization and Purification of Palygorskite

Palygorskite, with an idealized unit cell formula of M2+(x.ytzzy2'nHzO [Mgw
zRyClfl(Si3.xRx)020(OH)2(H20)4, where R is Al(|ll) or Fe(lll), a is an interstitial
vacancy, and M2+ is an exchangeable cation, adopts the pleated layered
structure shown in Figure 4.1 [6, 25]. The linked ribbons represent a 2:1 layer
that is continuous along the a-axis, but of limited lateral extend along the b-axis.
Rectangular channels, formed through the pleating of sheets, contain
exchangeable Ca2+ and Mg2+ cations, zeolitic water, and water molecules bound

106

to coordinatively unsaturated metal ion centers at the edges of the ribbons.
Small polar molecules, such as ammonia, and acetone, can access the tunnels
by displacing the zeolitic water molecules upon partial dehydration of the clay
materials. Further dehydration enables the formation of bonds between ternimal

Mg (II) and the small molecules to afford hybrid materials [26, 27].

 

 

Figure 4.1 A schematic [100] projection of palygorskite structure [6].

Sedimentation of a 10% (w/w) aqueous suspension of the mineral was
used to separate the clay from quartz and carbonate impurity phases [8]. Figure
4.2 Shows the X-ray patterns of the sedimented fractions collected after different
settling times. The impurity phases were almost completely removed from the

suspension after an aging time of 14 days. Owing to the low concentrations of

107

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the impurity phases, composites made from the as-received and purified versions
of the mineral showed no difference in mechanical properties. Thus, the as-
received mineral was used in all nanocomposite experiments.

Nitrogen adsorption measurements indicated the BET surface area of the
as-recived palygorskite to be 182 mzlg, which is substantially larger than the
surface area for pristine montmorillonite (~ 10 mzlg). The TEM images provided
in Figure 4.3 clearly show the presence of bundled lath-like crystals of
palygorskite ~ 5- 20 nm in width and several pm in length. Thermal gravimetric
analysis showed the loss of water over three temperature regions corresponding
to the loss of 7.5 % (w/w) surface and zeolitic water at 25—130 °C, 3 %
coordinated water (130—270 °C), and 6.5 % water due to dehydroxylation of the
2:1 layered structure (>270 °C) [10]. The ”Si MAS NMR spectrum of the as-
received palygorskite (not Shown) contains two Q3 peaks at -98 ppm and —92
ppm, corresponding to SiO4 units at the central and edge positions of the ribbons,
respectively [9]. The absence of a resonance characteristic of SiOH groups (~ -
84 ppm), indicates that the fraction of silicon centers at the external surfaces of

the Iaths is very small (<2 %).

109

 

Figure 4.3 TEM images of as-received palygorskite.

4.3.2 Silylated Paygorskite

The silylation reactions of palygorskite with y-aminopropyltrimethoxysilane
(APTMS), N-dodecyltriethoxysilane (DTES), and 1 ,1,1,3,3,3-
hexamethyldisilazane (HMSZ) resulted in no changes in the XRD pattern or the
”Si MAS NMR spectrum of palygorskite. The N2 adsorption-desorption
isotherms Show that the BET surface area of the HMSZ silylation product,
denoted HMSZ-PLG, was reduced slightly to 122 m2/g. The FTIR spectra of the
as-received palygorskite and the HMSZ-PLG silylation product are compared in
Figure 4.4. The appearance of weak C-H stretching vibrations near 2900 cm'1 are
observed for HMSZ-PLG, indicating that trimethylsilyl groups have been coupled
to the external surfaces of the mineral. Analogous FTIR bands were observed

for the APTMS- and DTES- PLG reaction products. Additional evidence in favor

110

of silylation was provided by the inability of the silylation products to disperse in

 

 

water.
2?
£5, Palygorskite
§
.8
E
E HMSZ-PLG
.—
3900 3400 2900 2400 1900 1400 900 400

Wavenumber(cm'1 )
Figure 4.4 FTIR spectra (KBr) of as-received palygorskite and the HMSZ
silylation product. The inset shows an expansion of the C-H stretching region

between 2800-3000 crn'1 for HMSZ-palygorskite.

4.3.3 Epoxy Nanocomposites

Rubbery and glassy epoxy nanocomposites were prepared by dispersing
as-received and silylated forms of palygorskite in degassed equivalent amounts
of epoxy monomer and curing agent and then curing the mixtures under nitrogen.

The dispersion of the mineral was qualitatively judged by examining the relative

111

 

 

 

z Palygorskite-epoxy
.4?
2
g
E
HMSZ-PLG-epoxy
0 10 20 30 40
28 ( ° ) / Cu K,

Figure 4.5 XRD patterns of the bottom sides of epoxy nanocomposite specimens

prepared from as-received and silylated palygorskite.

 

Figure 4.6 TEM image of a thin section of a glassy epoxy nanocomposite

containing 5% (w/w) of HMSZ-PLG
1 12

intensities of palygorskite diffraction peaks obtained from the bottom and top
surfaces of the composite specimens. As seen by the comparison of 110 and
400 reflections in Figure 4.5, much more palygorskite settles to the bottom of the
composite made from as-received palygorskite in comparison to the silylated
palygorskite composite. The top surfaces of both specimens show no
palygorskite reflections, indicating that there is a graded distribution of
palygorskite from top to bottom in both cases. But the particle gradient is
substantially less for the composite made from silylated palygorskite than for as-
received palygorskite.

A TEM image of a thin-section of a glassy epoxy composite containing 5
% (w/w) HMSZ-PLG is shown in Figure 4.6. In comparison to the image for the
pristine clay in Figure 4.3, the clay laths in the composite appear to be thinner,
suggesting that the aggregation of clay laths is reduced in the composite.

Representative stress - strain curves for HMSZ-PLG reinforced mbbery
epoxy composites are presented in Figure 4.7. The mechanical properties of
rubbery and glassy epoxy composites containing 0 - 10 % (w/w) palygorskite are
given in Table 4.1 and 4.2, respectively. Included in Table 1 for comparison are
the tensile properties of a rubbery epoxy nanocomposite reinforced by an
organo-montmorillonite [28]. The tensile properties of the rubbery composites
are substantially improved in comparison to the pristine polymer, but less so for
the glassy composites.

113

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1.4

I

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1 _ 5% 2%

l

0 °/o

Stress (M Pa)
0
on

0.6 ~
0.4 -

0.2 r

 

0 I l 41 1 l l I

0 5 10 15 20 25 30
Strain (%)

Figure 4.7 Representative strain-stress curves for rubbery epoxy composites

containing different loadings of silylated palygorskite HMSZ-PLG.

4.4 Discussion

The commercial form of palygorskite used in this study contained detectable
amounts of quartz and carbonate minerals (Figure 4.2). However, these impurity
phases were present in such small amounts that their removal by sedimentation
methods did not improve the nanocomposite performance properties of the
mineral. The hydrophilic nature of pristine palygorskite arises from the presence

of coordinated water at the edges of the pleated 2:1 layers and the filling of the

114

 

 

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channels by zeolitic water (Figure 4.1). Thermal dehydration of the mineral
removes the channel water, but dehydration is not expected to alter the
hydrophilic nature of the laths, because terminal hydroxyl groups still decorate
the external edge and base surfaces of the structure. Thus, pristine palygorskite
was not effectively dispersed in the epoxy pre-polymer and tended to sediment
out prior to polymer curing. However, silylation of the mineral substantially
improved dispersion, even though a graded distribution of particles was evident
based on the presence of observable XRD reflections from the bottom surfaces
of a specimen containing 5 % (w/w) of the mineral (Figure 4.5).

The silylation of palygorskite was readily accomplished using 7-
aminopropyltrimethoxysilane (APTMS), N-dodecyltriethoxysilane (DTES), and
1,1,1,3,3,3-hexamethyldisilazane (HMSZ) as silylating agents. The presence of
silylated surfaces was verified by FTIR (Figure 4.4). Bands near 2900 cm'1 were
assigned to C-H bond stretching modes of the grafted organosilane, in accord
with previous assignments [22-24]. It is noteworthy that the broad bands at 3556
cm'1 due to coordinated water and at 3410 cm'1 and 3283 cm‘1 due to zeolitic
water [7] indicated the retention of the channel water after silylation in toluene
under reflux conditions. The sharp band at 3620 cm'1 is assigned to the
stretching mode of hydroxyl groups in the 2:1 ribbons. Thus, the Silylation of the
mineral was confined to hydroxyl groups at the external edge and basal surfaces.
The hydrophobic nature of the silylated products was evidenced in part by the

116

inability to suspend the products in water.

Table 4.2 Tensile properties of glassy epoxy reinforced by silylated palygorskite

HMSZ-PLG. Numbers in the parentheses are relative standard deviations.

 

Clay loading Tensile strength Tensile modulus Tensile elongation

 

(wt%) (MPa) (GPa) (%)
0% 66.1 (1.4%) 2.9 (7.0%) 3.9 (4.0%)
2% 67.1 (2.1%) 3.1 (8.7%) 3.6 (4.6%)
5% 66.0 (2.5%) 3.0 (3.2%) 3.9 (2.8%)
10% 68.9 (4.0%) 3.3 (5.3%) 3.5 (6.1%)

 

The comparison of tensile properties in Table 4.1 for a rubbery epoxy
matrix shows that as-received palygorskite provides some improvement in the
modulus, but little or no improvement in the tensile strength at loadings of 2 and
5 % (w/w). The abrupt increase in tensile strength at a loading of 10 % (w/w) is
associated with an increase in the viscosity of the pre-polymer — palygorskite
mixture. The increase in viscosity retards the settling—out of the as-received
mineral prior to curing and provides a more uniform dispersion of palygorskite
particles in the matrix. In comparison to the pristine mineral, the silylated mineral
improves the degree of dispersion at loadings of 2 and 5 % (w/w). Consequently,
improved tensile properties are observed at these lower loadings for the silylated
mineral. But at a loading level of 10 % (w/w), where the increased viscosity of

the pre-polymer facilitates the dispersion of the pristine mineral, there is little

117

difference in the tensile properties of the composites formed from as-received
and silylated palygorskite.

It is remarkable that the pristine mineral in the absence of any organic
modifiers provides approximately a two-fold increase in strength and modulus for
the rubbery epoxy matrix at 10 % loading. Other pristine clay minerals (eg.,
montmorillonite) invariably weaken an engineering polymer, regardless of the
level of loading. The low layer charge, lath-like morphology and high surface
area (182 mzlg) of palygorskite makes it much more suitable for polymer
reinforcement in comparison to pristine smectite clays which retain their low
surface area tactoidal form (<10 m2/g) when dispersed in a polymer matrix.

In general, the as-received and silylated forms of palygorskite all provide a
substantial benefit in tensile properties at a loading of 10 % (wlw), due in large
part to the improved dispersion of the mineral laths in the matrix (cf., Table 4.1).
However, as shown in Table 4.1 for [CH3(CH2)17NH3]+exchanged montmorillonite
[28], this organoclay provides ~ 5 — 6-fold increases in tensile properties at 10 %
(w/w) loading. In comparison, palygorskite at the same loading level provides
only 2-3-fold improvements in tensile properties. Nevertheless, it also must be
realized that the improvements achieved with palygorskite come at a level of
organic modification that is < 0.2% of the levels needed to achieve the dispersion
of montmorillonite in the polymer matrix. Thus, the benefit to cost ratio lies in
favor of palygorskite.

118

Despite the reinforcement achieved for a rubbery epoxy matrix, little or no
benefit is realized when silylated palygorskite is dispersed in a glassy matrix (cf.,
Table 4.2). In view of the pleated layer structure of palygorskite, the tensile
properties of the mineral may be approaching those of the polymer matrix itself,
thus precluding the possibility of reinforcement. Further studies are needed to
explore this possibility.

Although palygorskite has been investigated as a reinforcement agent for
several other polymer systems [11-19], the benefits seldom approach those
observed in the present work. For instance, the tensile strength of polyethylene-
palygorskite formed by in-situ polymerization was increased by a maximum of
only 2 0 % in comparison to the pristine polymer [15, 1 6, 1 9]. Forthe related
polypropylene-palygorskite composite system, the clay served as a good
nucleating agent for isotactic polypropylene, and the relative crystallinity and
crystallization temperature of the composite materials generally increased with
filler content [17] An improvement of only 25% in mechanical properties was
found for polyimide composites containing 5% palygorskite [13].

Among the best improvements previously reported for any palygorskite
composite was the 220% boost in tensile strength found for polyurethane
reinforced by 10% (w/w) of the mineral silylated by N-[3-
(trimethoxysilyl)propyl]ethylenediamine [14]. As note above, comparable levels

of reinforcement are obtained for the rubbery epoxy composites reported in the

119

present work, even without silylation of the mineral. Thus, the key to achieving
polymer reinforcement benefits with palygorskite lies first in achieving optimal
particulate dispersion, whether the dispersions are achieved through viscosity

increases in the pre-polymer or through surface modification of the mineral itself.

4.5 Conclusion

Palygorskite, a silicate clay with a pleated 2:1 layered structure, has a
lath-like particle morphology that makes the clay, in both pristine and silylated
forms, an attractive candidate for the formation of polymer nanocomposites.
Three silylation reagents were used for surface modification of the mineral,
namely, aminopropyltrimethoxysilane (APTMS), N-dodecyltriethoxysilane (DTES),
and 1,1,1,3,3,3-hexamethyldisilazane (HMSZ). The silylated palygorskite
derivatives provided better dispersions in rubbery epoxy matrices than the
pristine mineral, affording improvements in mechanical properties at low loadings
levels of 2 and 5% (w/w). But at higher loadings where increases in the viscosity
of the pre-polymer helps to stabilizes the mineral dispersion, little or no
differences were observed for the reinforcement benefits provided by the pristine
and silylated forms of palygorskite. Glassy epoxy nanocomposites formed from
both pristine and silylated palygorskite exhibited marginal improvements in
tensile properties regardless of the mineral loading level, suggesting that the

tensile strength of the pleated sheet Silicate may be approaching that of the

120

polymer matrix.

121

References

(1) Becker 0.; Simon G. P. Adv Polym. Sci. 2005, 179, 29.

(2) D'Souza N. A. Encyclopedia of Nanoscience and Nanotechnology 2004, 3,
253.

(3) Okamoto M. Encyclopedia of Nanoscience and Nanotechnology 2004, 8, 791 .
(4) Ray S. S.; Okamoto M. Prog. Polym. Sci. 2003, 28, 1539.
(5) Murray H. H. Appl. Clay Sci. 2000, 17, 207.

(6) US. Geological Survey Open-File Report 01-041: http:/lpubs.usgs.govlof/
of01 -041lhtmldocs/clays/seppaly.htm.

(7) Augsburger M. S.; Strasser E.; Perino E.; Mercader R. C.; Pedregosa J. C. J
Phys. Chem. Solids 1998, 59, 175.

(8) Chipera S. J.; Bish D. L. Clays Clay Miner. 2001, 49, 398.

(9) D'Espinose de la Caillerie J-B, Fripiat J. J. Clay Miner. 1994, 29, 313-318.
(10) Guggenheim S.; van Groos A. F. K. Clays Clay Miner. 2001, 49, 433.
(11) Shen L.; Lin Y.; Du 0.; Zhong W.; Yang Y. Polymer 2005, 46, 5758.

(12) Wang L.; Sheng J.; Polymer 2005, 46, 6243.

(13) Lai S. Q.; Yue L.; Li T. S.; Liu X. J.; Lv R. G. Macromol. Mater. Eng. 2005,
290, 195.

(14) Ni R; Li J.; Suo J.; Li S. J. Mater. Sci. 2004, 39,4671.
(15) Du Z.; Rong J.; Zhang W.; Jing Z.; Li H. J. Mater. Sci. 2003, 38, 4863.

(16) Rong J.; Sheng M.; Li H.; Ruckenstein E. Polym. Compos. 2002, 23, 658.

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(17) Medeiros E. S.; Tocchetto R. S.; Carvalho L. H.; Conceicao M. M.; Souza A.
G. J. Ther. Anal. Calorim. 2002, 67, 279.

(18) Rong J.; Jing Z.; Li H.; Sheng M. Macromol. Rapid Commun. 2001, 22, 329.

(19) Rong J.; Li H.; Jing 2.; Hong X.; Sheng M. J. Appl. Polym. Sci. 2001, 82,
1829.

(20) Wang L.; Sheng J. J. Macromol. Sci. Pure Appl. Chem. A40 2003, 11 , 1135.
(21) Grandjean J.; Bujdak J.; Komadel P. Clay Miner. 2003, 38, 367.

(22) Johnson L. M.; Pinnavaia T. J. Langmuir 1991, 7, 2636.

(23) Okutomo S.; Kuroda K.; Ogawa M. Appl. Clay Sci. 1999, 15, 253.

(24) He J.; Shen Y.; Yang J.; Evans D. G.; Duan X. Chem. Mater. 2003, 15, 3894.

(25) Fernandez M. E.; Ascencio J. A.; Mendoza-Anaya D.; Rodriguez Lugo V.;
Jose-yacaman M. J. Mater. Sci. 1999, 34, 5243.

(26) Facey G. A.; Kuang W.; Detellier C. J. Phys. Chem. B 2005, 109, 22359.
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(28) Lan T.; Pinnavaia T. J. Chem. Mater. 1994, 6, 2216.

123

CHAPTER 5

MESOCELLULAR SILICA FOAM FOR POLYETHYLENE REINFORCEMENT

5.1 Introduction

Nanoparticle dispersion in polymer matrices provides an innovative
pathway to combine the properties of polymers and filler materials (often
inorganic particles). The enormous interacting surface area between polymer
matrix and filler nanoparticles makes it possible to transfer the properties of the
fillers, such as high strength, high thermal stability and low permeability, to the
polymer even at low filler loadings [1]. More interestingly, when the nanoparticle
fillers are capable of binding with the polymer matrix via chemical or physical
bonding, such as hydrogen bonding, they may affect the crystallinity and phase
behavior of polymers [2]. Nanocomposites can be divided into three categories
depending on whether the filler has one, two, or three dimensions in the
nanometer range. Clay minerals in exfoliated form are representative of the first
type of nanoparticle filler, since the silicate sheets are 1-nm thickness, but up to
several microns in diameter [3, 4]. Carbon nanotubes and metal nanoparticles
are representative of the second and third types of nanoparticle fillers,

respectively [5-7].

124

It is almost always necessary to modify inorganic filler particles for
polymer reinforcement purposes. The use of organic modifiers, which usually
have low molecular weight, may affect the properties of the composites via
plasticizer effects [8, 9]. The organic modifiers also are susceptible to thermal
degradation, and may initiate the restacking of clay layers or the decomposition
of the polymer matrix [10, 11]. The complicated effects of modifier on
nanocomposite formation make it difficult to correlate process-structure-property
relationships to optimize nanocomposite reinforcement. Additionally, using
organic modifiers increase the cost of composite processing.

Polyolefins are of critical significance in composite research, because they
represent the most commonly used thermoplastic polymers. Conventional fillers,
such as carbon fibers and glass fibers, are able to improve the mechanical
properties of polyethylene at high filler loadings [12, 13]. However, the fibers
have to be weaved to achieve uniform alignment, and, consequently, the
composites are anisotropic and are reinforced only in the direction of the oriented
fibers. More recently, clay minerals have been studied as polyolefin reinforcing
agents [14, 15]. Since polyolefins are highly hydrophobic polymers, the
preparation of polyolefin — clay nanocomposites requires modification to both
clays and polymers to achieve compatibility. The former is usually via organic
ion exchange reaction of the clay, and the latter is achieved by grafting maleic
anhydride to the polymer.

125

The term “polymer mesocomposite” was used to distinguish composites
reinforced by mesoporous silicas [16]. Mesoporous silicas can be synthesized
using supermolecular assembly techniques. They exhibit controlled pores in the
meso scale range (2-50 nm) and surface areas that are comparable to those of
exfoliated clays [17]. Attempts have been made to intercalate polymer chains
within the mesopores of silica particles using different techniques, such as in-situ
monomer polymerization and employing high-pressure as a driving force for
intercalation [18, 19, 20]. In a recent report, large pore worrnhole mesostructured
MSU-J silica, with and without the removal of the amine surfactant porogen, was
used to reinforce epoxy polymers. Systematic improvements were achieved for
the tensile strength, modulus, toughness, and elongation-at-break of the resulting
composites. [16]

In the present study we investigated mesocellular foam-stmctured MSU-F
silicas with different pore sizes for the preparation of polyolefin — silica
mesocomposites. The MSU-F silicas have pore sizes greater than 20 nm and
electrically neutral surfaces, making it possible to readily intercalate polymer
even without surface modification of the silica or the polymer. We investigated
both low density polyethylene (LDPE), and high density polyethylene (HDPE) as
the polymer matrix. Both LDPE - and HDPE - silica mesocomposites exhibited
improved tensile strength and modulus that are analogous to the best reported
values for polyethylene - organo-clay nanocomposites.

126

5.2 Experimental
5.2.1 Materials

The LDPE pellets were Dow Polyethylene 6401 with a density of 0.922
g/cc, and the HDPE pellets were Dow Polyethylene 25455N with a density of
0.955 g/cc. Pluronic P123 Block Copolymer was purchased from BASF. All
other chemicals were purchased from Aldrich, and used without further
punficafion.
5.2.2 Mesoporous silica synthesis

Mesocellular foam silicas with large pore sizes, denoted MSU-F, were
assembled using sodium silicate solution as the silica source and a triblock
copolymer P123 as the stmcture-directing poregen. The MSU-F used in the
present study was provided by Dr. Seong-Su Kim. Detailed preparation method
was reported in a previous report [21]. When ethanol as a co-solvent substitutes
in part for water at a concentration of 1 5 vol%, the product, denoted MSU-F2
exhibited further expanded pore size [22]. In a typical synthesis, 9.6 g (1.6 mmol)
porogen P123 was dissolved in a mixed solution of 80 mL 1M acetic acid, 20 mL
water and 60 mL ethanol in a shaking bath at room temperature for 1.5 h. After
the introduction of 7.968 g (0.06 mol) tetramethylbezene, the solution was again
mixed in a shaking bath at room temperature for 1 hour. A solution of sodium
silicate (21.6 9 containing 27% silica and 14% sodium hydroxide) was added to
the porogen solution forming a milky suspension. The suspension was shaked at

127

35 °C for 20 h, and heated at 100 °C under static conditions for 20 h. The
product was filtered, washed with water, dried in air and denoted MSU-F2. Both
MSU-F and MSU-F2 were calcined at 600 °C for 4 hours to remove P123 and
dried at 125 °C overnight prior to composite processing.
5.2.3 Mesocomposites preparation

Low density polyethylene (LDPE) and high density polyethylene (HDPE)
were used as polymer matrices for the mesocomposites. The LDPE and HDPE
pellets were dried at 75 °C under vacuum overnight beforehand. The composites
were processed using a twin-screw extruder (DSM Research 15 cc Micro
Extruder Injection molder), and the hot melt was transferred into a hot cylinder

and injection molded. The processing conditions are listed in Table 5.1.

Table 5.1 The extrusion and injection molding conditions for pure polymers and

their mesocomposites.

 

 

 

Extrusion Extrusion Cylinder Mold
temperature (°C) duration (min) temperature (°C) temperature (°C)
LDPE 185 10 185 70
HDPE 195 10 195 70
5.2.4 Characterization

N2 adsorption-desorption isotherms were recorded at -196 °C on a

Micrometritcs ASAP 2010 sorptometer. Prior to analysis, samples were

128

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 (T EM) images were obtained on a
JEOL 2200FS field emission microscope equipped 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. For
composites, thin sections (~100 nm thick) of the composites were put on naked
coppergfids.

Tensile mechanical 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 LDPE and its
composites and 1.5 mm/min for HDPE and its composites. The reported tensile
values are the ayerages for 4 specimens with dog-bone shapes. The standard
deviation for the modulus and strength was +/- 3%.

DSC measurements were conducted on a 2920 modulated DSC (TA
Instruments). An 8 mg quantity of sample was heated to 185 °C, and held at the
same temperature for 10 min, before cooling down to room temperature. The

heating and cooling rates were both 10 °C/min.

129

The polarized light microscope images were obtained using a LSM 5
Pascal confocal microscopy (Carl Zeiss, Inc). A 1 mg quantity of polymer or
composite specimen was melted between two glass slides to form a film for

microscopy imaging.

5.3 Results and Discussion

The surface properties of MSU-F and MSU-F2 were analyzed using N2
adsorption-desorption isotherms. The isotherm curves and pore Size
distributions are shown in Figure 5.1, and the surface properties are listed in
Table 5.2. MSU-F has a BET surface area of 544 mzlg, a cell size of 22 nm, and
a window size of 11 nm. The cell size and window size were determined from
the adsorption and desorption isotherms, respectively. MSU-F2 has a lower BET
surface area, but larger pore sizes and pore volume, especially cell size,
compared with MSU-F. The decrease in surface area of MSU-F2 might be due
to the fact that it has a greater wall thickness. The foam mesostnlcture was
clearly exhibited in the TEM images of both MSU-F and MSU-F2 (Figure 5.2).
The cell size of MSU-F is centered around 20 nm, and the cell sizes of MSU-F2
is centered over a broad range around 50 nm, consistent with the N2 isotherm

results.

130

Figure 5.1 (a) The N2 adsorption-desorption isotherms of MSU-F and MSU-F2.
The curves were offset by 200 for clarity. (b) Cell size (solid lines) and window
size (dash lines) distributions of MSU-F and MSU-F2 obtained from the
adsorption and desorption isotherm curves, respectively. The distributions were

determined by applying BJH model.

131

dvmo(unflgnm)

(a) 1800
1600
1400
,2 1200
O.
5
. 1000
2’
E, 800
8
a
> 600
400.
200.
0
040,7
0351
030‘

 

 

 

 

 

 

 

+ MSU-F2
- —I— MSU-F
0.0 0.2 0.4 0.6 0.8 1.0
PIPo
+ MSU-F cell size

- - --- - - MSU-F window size
+ MSU-F2 cell size
...k - - MSU-F2 window size

 

40
Pore diameter (nm)

50 60 70 80

132

Table 5.2 The surface properties of MSU-F and MSU-F2

 

 

BET S.A Cell size Window size Pore volume
(mzlg) (nm) (nm) (ems/g)
MSU-F 544 22 11 2.24
MSU-F2 410 50~70 16 2.34

 

The relative standard deviations of the BET surface areas are less than 0.5%.

 

Figure 5.2 The TEM images of (a) MSU-F and (b) MSU-F2.

Both MSU-F and MSU-F2 were extruded with PE to make the
mesocomposites. The optimum blending time was 10 min. Other attempts
showed that too short an extrusion time (5 min) resulted in the presence of visible
white particles in the products, while too much extrusion (15 min) lowered the
mechanical performance of the composites. The TEM images of composite thin-
sections (Figure 5.3) exhibit the presence of silica particles of semi-micron sizes

and large blank polymer regions (not shown in the images). Despite of the non-

133

uniform silica dispersion, the polymer chains filled the pores of the silica particles,
as revealed by the TEM EDS analysis showing the presence of C element in the

silica porous region.

 

Figure 5.3 TEM images of thin-sections of a HDPE — MSU-F 4.2 wt%

mesocomposites.

The crystallinity of the LDPE and its silica mesocomposites were studied
using DSC (Table 5.3). The melting or crystallization temperatures of the
composites are comparable with those of pure LDPE. However, the degree of
crystallinity was slightly reduced with the addition of silica, especially at high
loading of 7.4 wt%. Polarized light microscopy (PLM) is also a powerful tool to
study polymer crystallinity, such as spherulite sizes and nucleation effects. The

polarized light can only pass through the polymer crystal Iamellas with a

134

matching chain orientation direction, and, consequently, gives bright spots.
Polymer chains that are perpendicular to the polarize plane will appear as dark
spots on the image. LDPE spherulites can be clearly seen in the PLM image
(Figure 5.4). The diameters of the spherulites have a wide range from several
microns to over 20 microns. In a LDPE mesocomposites containing 5.5 wt%
MSU-F, the range of spherulite sizes is similar as in LDPE, despite the lose in
contrast due to the scattering of light by the silica particles. HDPE has a higher
crystallinity compared with LDPE, and, consequently, the addition of MSU-F had
a greater effect in lowering the degree of crystallinity. A more than 10%
decrease in degree of crystallinity was observed upon the addition of 8.6 wt%
MSU-F. A decrease in crystallinity is commonly observed for polymer
composites, especially those reinforced by nano- or sub-micron particles. The
lose of crystallinity is related to the fast crystallization rate and the nucleation
effects of the filler particles [23].

Due to the compatible surfaces and good dispersion of silica in the
polymer matrix, the foam-structured mesoporous silicas provided substantial
reinforcement in mechanical properties to both LDPE and HDPE. For LDPE, the
addition of 8.3 wt% MSU-F improved the strength and modulus by 55% and
100%, respectively (Table 5.4). This improvement is comparable to that
achieved in PE-organo-clay systems that require extensive modification of both
polymer and clay [14, 15]. However, at low MSU-F loading (5.5 wt%), the

135

improvement in strength and modulus was less Significant. The elongation at
break decreased with the addition of MSU-F. The large pore sizes of MSU-F
allow polymer intercalation inside the pores, and, consequently, provided
reinforcement to the polymer. Similar improvement was achieved using MSU-F2,
which had larger pore sizes than MSU-F, indicating that a further expansion of

pore Sizes of the silicas has negligible effects on the reinforcement.

Table 5.3 The crystallinity of LDPE and HDPE mesocomposites.

 

 

Polymer Filler and Melting Crystallization AH", a Degree of

loading (wt%) temp. (°C) temp. (°C) (J/g) crystallinity b (%)

LDPE Pristine 111.6 97.0 111.1 38.3

MSU-F 5.5 111.6 96.9 112.8 38.9

MSU-F 8.3 111.6 97.1 105.3 36.3

MSU-F2 7.4 112.0 97.4 101.7 35.1

HDPE Pristine 133.1 1 15.4 200.2 69.0

MSU-F 4.2 133.1 116.0 187.0 67.3

MSU-F 8.6 132.7 117.0 162.0 61.1

 

a The heat of fusion was obtained from the area under the melting curve from DSC scan.
b The degree of crystallinity was obtained by comparing the AH", with the 290 J/g heat of fusion
for a 100% crystalline polyethylene.

When HDPE was used as the polymer matrix, substantial improvement
was also observed with the addition of MSU-F (Table 5.5). Specifically, the
strength and modulus were improved by 40% and 60% respectively at 8.6 wt%
MSU-F loading. Unlike the LDPE system, improvement was also achieved at a

low loading 4.2%. Both tensile elongation and elongation at break were reduced

136

with the addition of MSU-F. However, they were not significantly weakened at

the low loading (4.2 wt%).

 

Figure 5.4 PLM images of (a) LDPE and (b) a LDPE - MSU-F 5.5 wt% loaded

mesocomposite.

Compared with other commonly used inorganic filler materials for
polyethylene reinforcement, there are advantages of large pore mesostructured
silicas. Firstly, no surface modification is necessary for both the silica and the
polymer; secondly, the reinforcement in mechanical properties far exceeds those
of composites made from regular sphere silicas [24]; thirdly, due to the
mesoporous structure of MSU-F, the reinforcement benefit is comparable to the

best reported PE - organo-clay nanocomposites; fourthly, the isotropic nature of

137

the MSU-F pore system eliminates the issue of orientation, which is a common
occurrence for fibrous fillers.

The oxygen permeability of LDPE was increased by 15.7% with the
addition of 5.5 wt% MSU-F, and by 42.6% with the addition of 7.4 wt% MSU-F2
(Table 5.4). The reduction in crystallinity may contribute to the increased
permeability of the mesocomposites. This phenomenon is consistent with the
literature studies on composites membrane permeability and selectivity, which
attributed the increase in permeability to increases in free volume between the
polymer matrix and filler surfaces [25]. This result further indicates that even
greater improvement in material properties can be expected once better

compatibility of PE and silica surfaces is achieved.

Table 5.4 The mechanical properties and oxygen permeabilities of LDPE - silica

 

 

mesocom posites.
Loading Tensile Tensile Elongation at Oxygen
(wt%) strength modulus break (%) permeability
(MPa) (MPa) (cc.mil/m2d)
none - 14.1 119.7 104.0 5456.7
MSU-F 5.5 15.8 156.8 42.3 6311.0
8.3 21.8 240.3 32.2 -
MSU-F2 3.0 13.7 164.6 49.5 -
7.4 20.3 243.7 41.1 7779.5

 

The relative standard deviations are less than 3% for strength and modulus, and are less than
20% for the elongation.

138

Table 5.5 The mechanical properties of HDPE - silica mesocomposites.

 

 

Loading Tensile Tensile Tensile Breaking
(wt%) strength modulus Elongation elongation
(M Pa) (M Pa) (%) (%)
pristine - 21.9 733 1 1 .7 830
MSU-F 4.2 26.8 891 9.5 670
MSU-F 8.6 31.1 1165 8.2 28

 

The relative standard deviations are less than 3% for strength and modulus, and are less than
20% for the elongation.

5.4 Conclusion

Mesoporous silica MSU-F with large pore sizes allow polyethylene
intercalation inside the porous regions without the need for surface modifiers.
LDPE- or HDPE- mesocomposites formed from MSU-F silica exhibited
significantly improved strength and modulus, despite the in-homogeneous overall
dispersion of the silica in the polyethylene matrix. However, the crystallinity of
the mesocomposites was slightly reduced, as were the barrier properties,
indicating some weakness in surface affinity between polyethylene and silica
particles. Chemical modification might be able to further improve the dispersion

of silica, and to enhance the compatibility between polymer and silica particles.

139

References

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(19) Ji, X.; Hampsey, J. E.; Hu, 0.; He, J.; Yang, 2.; Lu, Y. Chem. Mater. 2003,
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(21) Kim, S. S.; Pauly, T. R.; Pinnavaia, T. J. Chem. Comm. 2000, 17, 1661.
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141

CHAPTER 6

CONCLUSIONS AND FUTURE DIRECTIONS

6.1 Comparison of Various Polymer Reinforcing Agents

In the proceeding chapters, three inorganic particles, namely, synthetic
saponites, palygorskite, and mesocellular silica foam (MSU-F), have been
studied as polymer reinforcing agents with little or no need for a surface organic
modifier. The polymer matrices included thermoset epoxy and thermoplastic
polyethylene. In this chapter, as a conclusion to the thesis, epoxy - synthetic
saponite nanocomposites and polyethylene (PE) - MSU-F mesocomposites will
be compared with literature reports on nanocomposites based on epoxy and PE
made from different reinforcing particles to evaluate the advantages of synthetic
saponite and MSU-F as polymer reinforcing agents.
6.1 .1 Nanoparticles for Epoxy Reinforcement

Spherical alumina nanoparticles have been used for epoxy reinforcement
in both untreated or silylation forms. At loadings in the range of 1 to 10 wt%, the
inorganic alumina had negligible effects on the tensile strength and modulus of
epoxy [1]. In a separate study, it has been shown that inorganic alumina
improved the tensile modulus by 40% ata loading of 15%, but weakened the

strength and elongation, while [3-(2-aminoethylamino)—propyI]-triethoxysilane

142

modified alumina could only manage to maintain the strength and modulus of
pristine epoxy [2].

Silica nanoparticles in epoxy polymers can be formed in-Situ by mixing
alkoxysilane (TEOS) aqueous solution with epoxy monomer in organic solvent to
form a homogenous mixture [3-5]. The gelation reaction of hydrolyzed TEOS
occurs simultaneously with the evaporation of solvents, resulting in a colloidal
dispersion of silica nanoparticles in the epoxy resin. In a representative report,
epoxy - silica nanocomposites made through the in-situ sol-gel method exhibit a
15% increase in tensile modulus, a 40% increase in fracture toughness and
same strength, compared with pristine epoxy, at a filler loading of 8 wt% [4].
Fumed Silica can also be directly introduced to a mixture of epoxy monomer and
curing agent to form nanocomposites upon epoxy curing. However, dramatic
decreases in stress and elongation were observed [6].

Ever since lijima first discovered carbon nanotubes in 1991, carbon
nanotubes have been attractive polymer reinforcing agents due to the
combination of their superb mechanical properties, electrical conductivities, and
thermal stability and conductivity [7]. Molecular modeling approaches have
predicted greater improvement in material properties [8]. Even after purification,
carbon nanotubes still have a great tendency to form bundles through
aggregations. Therefore, in order to achieve good dispersion in an epoxy matrix,
carbon nanotubes need to be either surface modified or to be put into a

143

homogenous suspension in a solvent, followed by mixing with epoxy prepolymer.
Plasma treatment and maleic anhydride grafting of the carbon nanotubes
provided a 50% increase in tensile strength, a 100% increase in tensile modulus,
as well as improvement in elongation and glass transition temperature at a
loading of only 1 wt%[9]. However, such improvements were based on a rubbery
epoxy system, which is relatively soft and easy to reinforce. The reinforcement
of carbon nanotubes became less significant for rigid glassy epoxies. For
example, in a study that utilized acetone to achieve carbon nanotube dispersion,
the addition of 0.5 wt% carbon nanotubes improve the tensile strength by 100%
for an soft epoxy (tensile strength was 5.4 MP3 for pristine epoxy), but had no
effects on a glassy epoxy with a tensile strength of 45.5 MPa. The reason for the
latter results was attributed to poor interfacial interaction between the carbon
nanotubes and the polymer [10].

Carbon nanofibers also have been used to reinforce epoxy. At a 2 wt%
nanofiber loading, increase of 17% and 14% can be achieved in strength and
modulus, respectively [11]. However, the strength tends to decrease with
increase in carbon nanofiber loadings [11, 12].

Epoxy - organo-clay nanocomposites have been extensively studied.
Great improvements were reported for rubbery epoxy matrices. More than a 10-
fold increase in strength and modulus is realized by the addition of 15 wt% of the
exfoliated organoclay [13]. However, for clays dispersed in glassy epoxies, the

144

strength of the nanocomposites was not significantly improved [14-16]. Even
when uniform dispersion of exfoliated or intercalated clay nanolayer is achieved
in epoxy, a 25% decrease in strength compared with pristine epoxy is observed
with the addition of 5 wt% organoclay, although the modulus is improved by 40%
[16]. It was proposed that the presence of intercalated clay nanolayers may
initiate the rupture of the composites under stress and, consequently, lower the
mechanical properties.

In chapter 3, a synthetic saponite was investigated for rigid glassy epoxy
reinforcement. Compared with the other nanoparticles for glassy epoxy
reinforcement, saponite is the only example that improves the tensile strength,
modulus, elongation and toughness of the polymer. Additionally, facile
dispersion of the saponite nanoparticles in the polymer matrix can be achieved
without the need for an organic surface modifier.

6.1.2 Nanoparticles for Polyethylene Reinforcement

Polyethylene (PE) is one of the most extensively used thermoplastic
polymers. Property enhancement of PE using various nanoparticles has been
studied, and the representative results are summarized in Table 6.1. The results
obtained for LDPE - MSU-F mesocomposites described in chapter 5 also are

included for comparison.

145

 

Table 6.1 Comparison of various inorganic particles for PE reinforcement.

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Composite system Filler Tensile Tensile Breaking
Loading strength modulus elongation
(MPa) (MPa) (%)
LDPE - MSU-F - 14.08 119.7 104.0
8.3 wt% 21.80 240.3 32.2
(3.8 vol%)
HDPE - oriented - 30.5 1076 600
carbon fiber [17] 50 vol% 317 3.1*10‘ -
LDPE - CNT [18] - 10.7 235 380
10 wt% 15.6 444 12
LLDPE - glass fiber ' - 10.4 74 34
I19] 20 wt% 8.9 79 26
HDPE - 29 870 440
- [20] silica particle 0.75 vol% 30 900 440
g-PBA silica 1 0.75 vol% 29 1100 330
LLDPE - 1 1.8 190 >400
- [21] organo-MMT ' 7 wt% 15.3 435 >400
organo-MMT, 7 wt% 18.8 569 221
gMAH-PE

 

' LLDPE, g-PBA, MMT, and g-MAH are abbreviations for linear LDPE, polybutyl acrylate-grafted,
montmorillonite, maleic anhydride grafted, respectively.

Mesocellular silica foam MSU-F provides far better reinforcement for PE in

comparison to glass fiber and spherical silica nanoparticles. The reason may be

due to the mesoporous structure of MSU-F which isolates the polymer in the

mesopores. The reinforcements provided by carbon nanotubes and organo-clay

are comparable with those achieved by MSU-F. However, the former two

composites required complicated processing methods. For example, 10 cycles

of extrusion and molding were needed to achieve good dispersion of CNT in PE.

146

In the PE - organo-clay system, both clay and polymer must be modified to
enhance surface compatibility. Carbon fibers provided significant reinforcement
to PE, but a loading as high as 50 vol% was required. Additionally, the carbon
fibers had to be weaved and oriented in PE, making the composites anisotropic.
In comparison to other particles, MSU-F is a good reinforcing agent for PE, and

the composites can be made readily through conventional melt blending.

6.2 Future Directions

Synthetic saponite with an irregularly stacked morphology has proven to
be an effective reinforcing agent for epoxy polymers. However, the composite
preparation requires on the assistance of ultrasonication in solvent suspension to
break down saponite particles to the nanoscale prior to composite preparation.
Therefore, epoxy - saponite nanocomposite deserves further exploration to
simplify the preparation procedures by avoiding the use of solvent and
ultrasonication. Possible methods include saponite edge decoration to reduce
the edge-face interactions between clay nanolayers.

The use of saponite for polymer reinforcement deserves further expansion
into polymer systems other than epoxy. Preliminary studies on PE -saponite
composites prepared through melt-blending showed no improvement in the
strength of the composites, possibly due to (i) incompatibility between the

saponite surfaces and the highly hydrophobic PE matrix; (ii) the presence of sub-

147

micron saponite particles. The first limitation may be solved by employing a
small amount of maleic anhydride grafted PE into the system to enhance surface
compatibility. Regarding the second limitation, a new technique has been
reported that has a ultrasonic slit die attached to a compounding extruder [22].
This technique has been used in enhancing organo-clay dispersion in polyolefin,
and may have a promising ability to disperse saponite in polyolefin as well.

One of the advantages of synthetic saponite is that it has an irregularly
stacked layered morphology. Therefore, organic modification of the clay
surfaces no longer is necessary to achieve delamination of the layers. It will be
interesting to use saponite to form nanocomposites based on polymers that
require high processing temperature, such as nylon, polystyrene, polyimides and
poly(ethylene terephthalate). For these polymers, it has been reported that
alkylammonium exchanged clays have the modifiers decomposed at the polymer
processing temperatures [23]. Therefore, a completely inorganic reinforcing
agent, such as synthetic saponite, can be a great advantage in developing
polymer nanocomposites.

Mesocellular MSU-F foam has the ability to reinforce polyethylene without
any surface modification. The questions remaining in this research are whether
or not PE chains intercalate into MSU-F pores and, if so, whether the structure or
crystallinity of the PE chains is altered in the MSU-F pores. Preliminary
investigation has been carried out by measuring composite density. It is known

148

that pristine LDPE and silica have a density of 0.92 glcc and 2.15 g/cc,
respectively, and that MSU-F has a pore volume of 2.3 cc/g. The density of a 5.5
wt% L DPE - M SU-F mesocomposites s hould be either 0 .95 g/cc o r 0.85 g/cc,
assuming that the porous region on MSU-F surfaces is filled with LDPE or empty,
respectively. Experimentally, a 1 cm X 6 mm X 2 mm pristine LDPE block and a
5.5 wt% LDPE - MSU-F mesocomposite block were both emerged into ethanol in
a 10 mL graduate cylinder, and the mass and volume of the material were
recorded to calculate the density. The densities of LDPE and the 5.5 wt% LDPE
- MSU-F mesocomposite were 0.93 g/cc and 0.95 g/cc, respectively, agreeing
with the assumption that the mesopores on MSU-F were filled with LDPE chains.
However, more elaborate evidence for pore filling is needed. Such evidence may
be obtained by using positron annihilation lifetime spectroscopy for the
measurement of free volume in the composite [24]. TEM-EDS line scan is a
promising way to image the 3-D structure of the composites. Another approach
is to use HF to etch the silica portion of the composite and image the remaining
polymer portion. It may be possible to fill the silica voids with heavy scattering
materials to enhance the contrast. However, maintaining empty pores during
TEM sample preparation, particularly under microtome conditions, will be a

challenging task.

149

Preliminary studies indicate that MSU-F provides little reinforcement
benefit for polypropylene. Further exploration of the PP - MSU-F system will
require the study of polymer structure in silica pores, as described above.

To conclude, this thesis research hopefully has demonstrated that
inorganic particles with suitable surface properties can be used to make polymer
nanocomposites that are as good, if not better, than those made from organically
modified particles. Saponite and MSU-F are two examples. A wide range of
inorganic nanoparticles, such as layered and fibrous alumina, layered and large
pore zeolites, silica nanoparticles, carbon nanotubes, nano clay, makes the field

of polymer nanocomposites research full of excitements and opportunities.

150

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152

   
 
 

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