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Nanoclay-reinforced polyurethane

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Nanoclay-reinforced polyurethane

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NANOCLAY-REINFORCED POLYURETHANE

by

Chia-Hao Wang

A Dissertation Presented to the
FACULTY OF THE GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(MATERIALS SCIENCE)

December 2009
Copyright 2009 Chia-Hao Wang

ii
Acknowledgements

First and foremost, I would like to thank my advisor, Professor Steven R. Nutt, for
the guidance, encouragement, kindness, and support that he provided as a mentor. I
believe that this work would not been accomplished without his valuable advice and
endless patience. He always provided unlimited accessibility to discussion and immediate
assistance in facilitating research works. He has helped me significantly to develop
professional skills and knowledge.
I would also like to sincerely thank Dr. Yeong-Tarng Shieh for taking the time to
work with me through the later stage of my graduate career. I would also like to extend
the gratitude to all my committee members: Dr. Katherine Shing and Dr. Charles
Sammis for their agreement to serve as my PhD guidance committee and taking the time
and efforts to evaluate my work.
I would also thank Dr. Maria Lujan Auad and Dr. Hongbin Lu who lead me into
the fascinating world of materials science. In particular, Dr. Maria Lujan Auad guided me
immensely and was always prompt in coordinating efforts. I would also thank Mr.
Warren Haby, a laboratory manager at University of Southern California, for his great
advice and technical support facilitating experiments. Also thanks to the folks at the USC
Composites Center where I had the privilege of working with group of most wonderful
colleagues. I wish them the best in all their future endeavors.
Last, but not least, I would express my sincere gratitude and respect to my parents
in Taiwan, brother and sister in-law in US for the un-ending support and encouragement.
iii
Table of Contents
Acknowledgements............................................................................................................. ii
List of Tables ..................................................................................................................... vi
List of Figures................................................................................................................... vii
Abstract.................. …………………………………………….………………………….x

Chapter 1 Introduction................................................................................................... 1
1.1 Polymer Nanocompsoites ......................................................................................... 1
1.2 Layered Silicates....................................................................................................... 3
1.3 Dispersion of Layered Silicates ................................................................................ 7
1.4 Research Objectives.................................................................................................. 9
Chapter 1 References................................................................................................... 11

Chapter 2 2-D Nanoclay-Reinforced Polyurethane..................................................... 13
2.1 Motivation............................................................................................................... 13
2.2 Experimental........................................................................................................... 15
2.2.1 Materials .......................................................................................................... 15
2.2.2 Synthesis of PU elastomer and PU/MMT nanocomposites............................. 16
2.2.3 Characterization ............................................................................................... 20
2.3 Results and Discussion ........................................................................................... 21
2.3.1 Morphologies of organic MMTs and PU/MMT nanocomposites ................... 21
2.3.2 FTIR characterization ...................................................................................... 26
2.3.3 Tensile properties............................................................................................. 32
2.3.3 DSC measurements.......................................................................................... 35
2.4 Conclusions............................................................................................................. 38
Chapter 2 References................................................................................................... 39

iv
Chapter 3 1-D Nanoclay-Reinforced Polyurethane..................................................... 42
3.1 Motivation............................................................................................................... 42
3.2 Experimental........................................................................................................... 43
3.2.1 Materials .......................................................................................................... 43
3.2.2 Preparation of ATT-OH and ATT-MDI .......................................................... 43
3.2.3 Synthesis of PU elastomer and ATT/PU composites ...................................... 46
3.2.3 Characterization ............................................................................................... 49
3.3 Results and Discussion ........................................................................................... 50
3.3.1 Characterization of ATT.................................................................................. 50
3.3.2 ATT/PU composites......................................................................................... 56
3.3.2.1 DSC measurements................................................................................... 56
3.3.2.2 Morphology of PU elastomer and ATT-MDI/PU nanocomposites.......... 58
3.3.2.3 Dynamic Mechanical Analysis (DMA) .................................................... 61
3.3.2.4 Tensile properties...................................................................................... 64
3.3.2.5 Thermal degradation ................................................................................. 68
3.4 Conclusions............................................................................................................. 70
Chapter 3 References................................................................................................... 71

Chapter 4 Effects of Organophilic Modified Attapulgite on Polyurethane................. 73
4.1 Motivation............................................................................................................... 73
4.2 Experimental........................................................................................................... 74
4.2.1 Materials .......................................................................................................... 74
4.2.2 DOE, heat treatment, acid treatment, and grafting of ATT ............................. 74
4.2.3 Synthesis of PU elastomer and PU/ATT-MDI nanocomposites...................... 75
4.2.4 Characterization ............................................................................................... 78
4.3 Results and Discussion ........................................................................................... 79
4.3.1 Characterization of ATT.................................................................................. 79
4.3.2 PU/ATT-MDI nanocomposites........................................................................ 86
4.3.2.1 Morphology of PU/ATT-MDI nanocomposites ............................................... 86
4.3.2.2 DSC measurements........................................................................................... 88
4.3.2.3 Dynamic Mechanical Analysis (DMA) ............................................................ 90
4.3.2.4 Tensile properties.............................................................................................. 92
4.4 Conclusions............................................................................................................. 94
Chapter 4 References................................................................................................... 95

v
Chapter 5 Effects of Heat and Acid Treatments on Attapulgite.................................. 97
5.1 Motivation............................................................................................................... 97
5.2 Experimental........................................................................................................... 98
5.2.1 Materials, heat- and acid-treated ATT................................................................. 98
5.2.2 Characterization ................................................................................................... 99
5.3 Results and Discussion ......................................................................................... 100
5.3.1 TGA measurement for heat-treated ATT....................................................... 100
5.3.2 FTIR characterization for heat-treated ATT.................................................. 103
5.3.3 XRD analysis for heat-treated ATT............................................................... 106
5.3.4 Morphology for acid-treated ATT ................................................................. 109 T 5.3.5 FTIR characterization for acid-treated ATT.................................................. 112
5.3.6 XRD analysis for acid-treated ATT............................................................... 116
5.4 Conclusions........................................................................................................... 118
Chapter 5 References................................................................................................. 119

Chapter 6 Conclusions and Suggestions for Future Works ....................................... 120
6.1 Conclusions........................................................................................................... 120
6.2 Suggestions for Futuer Works .............................................................................. 120

Bibliography ....................................................................................................................123

List of Tables
Table 2.1 The composition and hard segment content of PU and PU/MMT
nanocomposites.

18
Table 2.2 The area ratio of the IR absorption peaks of N–H
stretching (A
NH
) and C–H stretching (A
CH
) and the area
ratio of the IR absorption peaks of the hydrogen-bonded
C=O (A
hCO
) and free C=O (A
fCO
).

31
Table 2.3 Tensile properties of PU and PU/MMT nanocomposite.

34
Table 2.4 Melting temperature and enthalpy of soft segment in PU and
PU/MMT nanocomposites.

37
Table 3.1 Neat PU and ATT/PU Nanocomposites.

47
Table 3.2 Thermal Properties of Neat PU and the ATT/PU
Composites

57
Table 3.3 Storage modulus values (G’) at 40°C and 100°C and the
peak in tan δ vs. temperature curve for hard segment.

63
Table 3.4 Young’s modulus and yield stress for ATT/PU composites
with different functional fillers.

67
Table 4.1 Grafting amount of MDI as obtained from TGA heating at
20
o
C/min.
84

Table 4.2 Melting temperature and enthalpy of soft segment in PU
and PU/ATT-MDI nanocomposites.
89

Table 4.3 Storage modulus values (G’) at 50°C and 100°C and the peak 91
temperatures of the tan δ curves for soft and hard
segments.

Table 4.4 Tensile properties of PU and PU/ATT-MDI
nanocomposite.

93
Table 5.1 Thermal analysis of attapulgite at various stages.

101
vi
vii
List of Figures
Figure 1.1 Structure of montmorillonite

5
Figure 1.2 Structure of attapulgite

6
Figure 1.3 Scheme of three types of morphologies in polymer-
layered silicate composites. (a) phase separated
microcomposites (b) intercalated nanocomposite (c)
exfoliated nanocomposite.

8
Figure 2.1

Synthesis and chemical structure of PU/MMT
nanocomposite.

19
Figure 2.2 WAXD patterns for (a) PU/MMT-30B and (b) PU/MMT-I30E
nanocomposites.

24
Figure 2.3 TEM images of (a) PU2900/MMT-30B and (b)
PU2900/MMT-I30E nanocomposites.

25
Figure 2.4 FTIR spectra of (a) MMT-30B and (b) MMT-I30E.

29
Figure 2.5 FTIR spectra of (1) PU1000 series, (2) PU2000 series, and
(3) PU2900 series in the range from 1650 to 1800 cm
-1
. (a:
PU; b: PU/MMT-I30E; c: PU/MMT-30B.)

30
Figure 3.1 Modification and chemical structure of the ATT nanorods.

45
Figure 3.2 Synthesis and chemical structure of the ATT–MDI/PU
nanocomposite.

48
Figure 3.3 TEM image of the original ATT. 53
Figure 3.4 TGA curves for the ATT nanorods before and after
modification.

54
Figure 3.5 FTIR spectra for the (a) original ATT, (b) ATT–OH, and (c)
ATT–MDI.

55
Figure 3.6 TEM images of ATT–MDI/PU nanocomposite with (a) 2.5 and
(b) 10 wt % nanorods.

59
Figure 3.7 SEM images of (a) neat PU (b) ATT-MDI/PU nanocomposite
with 5%wt nano-rods.
60
viii
Figure 3.8 Yield stress and Young’s modulud for ATT-MDI/PU
nanocomposites with different amounts of modified ATT.

65
Figure 3.9 Yield stress and Young’s modulus for different functional fillers
of ATT-MDI/PU composites.

66
Figure 3.10 Thermal degradation behaviors of PU and ATT-MDI/PU
nanocomposites.

69
Figure 4.1 Synthesis of PU/ATT-MDI nanocomposite.

77
Figure 4.2 TEM images for (a) untreated ATT and (b) heat and acid-
treated ATT.

82
Figure 4.3 TGA curves for the ATT nano-rods before and after
grafting.

83
Figure 4.4 FTIR curves for (a) ATT (b) H700A3M2H ATT-OH (c)
H700A3M2H ATT-MDI.

85
Figure 4.5 TEM images of (a) PU/ATT-MDI-H500A1M1H (b)
PU/ATT-MDI-H600A5M4H and (c) PU/ATT-MDI-
H700A3M2H nanocomposites.

87
Figure 5.1 TGA curves for the ATT nano-rods after heat treatments at
various temperatures.

102
Figure 5.2 FTIR curves for the ATT nano-rods (a) before, and after
heat treatments at (b) 500
o
C, (c) 600
o
C, (d) 700
o
C, and (e)
800
o
C.

105
Figure 5.3 WAXD patterns for ATT (a) without heat treatment, with
heat treatments at (b) 500
o
C, (c) 600
o
C, (d) 700
o
C, and (e)
800
o
C.

107
Figure 5.4 Morphological changes of ATT after heat treatment at
high temperatures.

108
Figure 5.5 AFM image for acid-treated sample HT500AT1H.

110
Figure 5.6 TEM image for acid-treated sample HT800AT8H.

111
Figure 5.7 FTIR spectra for ATT nano-rods after heat and acid treatments:
(a) HT500AT1H (b) HT500AT8H (c) HT800AT1H (d)
HT800AT8H.
114
ix
Figure 5.8 The mechanism of acid-activation for Attapulgite. 115
Figure 5.9 WAXD patterns for (a) HT500AT1H (b) HT500AT8H (c)
HT800AT1H (d) HT800AT8H.

117
Figure 6.1 Scheme of the microwave-exfoliation nanocomposite (a)
Clay with polar molecules in the galleries (b) Popcorn
effect of polar molecules under microwave treatment (c)
Exfoliated silicate layers in nanocomposite.

122
x
Abstract

Polyurethane-layered silicate nanocomposites with different layered silicates,
montmorillonite (MMT) and attapulgite (ATT), were investigated to gain fundamental
understanding of the role of nanofillers, the chemistry of the modifiers, and the physics of
polymer nanocomposites.
Polyurethane-based nanocomposites with compositions that included soft segments
with number average molecular weights of 1000, 2000, and 2900, and organic-modified
MMT (including MMT-30B and MMT-I30E) were prepared by in situ polymerization.
Transmission electron microscopy (TEM) and wide-angle X-ray diffraction (WAXD)
results revealed that both MMT-30B and MMT-I30E were intercalated and partially
exfoliated by the PU. Mechanical tests showed that the PU1000 series in soft-segment
molecular weight yielded superior tensile properties compared to the PU2000 and
PU2900 series. Also, for a given molecular weight of soft segment in PU, the MMT-30B
nanocomposites exhibited greater increases in Young’s modulus, tensile strength and
elongation at break than the MMT-I30E counterpart, and the crystallinity of PU was
enhanced by the clays.
In pristine PU and composites containing ATT nano-rods, the effects of
functionalization on dispersion and on thermal, physical, and mechanical properties were
characterized by Fourier transform infrared spectroscopy (FTIR), thermogravimetric
analysis (TGA), transmission electron microscopy (TEM), differential scanning
calorimetry (DSC), and mechanical testing. Functionalization of ATT nano-rods proved
beneficial in terms of dispersion, degree of crystallinity, thermal stability, and tensile
xi
strength. A higher grafting content of 4,4’-methlenebis(phenyl isocyanate)-modified
attapulgite (ATT-MDI) led to more uniform dispersion and enhanced tensile properties of
the nanocomposites. However, the grafting content had negligible effects on the thermal
behavior of PU/ATT-MDI nanocomposites.
To extend the applications of attapulgite to specific polymer systems and to
heterogeneous catalytic reactions, the physicochemical changes resulting from thermal
and acid treatments were investigated. TGA analysis of the untreated attapulgite revealed
that dehydration and dehydroxylation occurred during heating. The dehydration and
dehydroxylation led to changes in attapulgite structure, as demonstrated by FTIR spectra
of heat-treated samples. The FTIR spectra also revealed that acid treatment led to
hydroxylation of the heat-treated attapulgite. WAXD patterns of heat-treated samples
showed a loss of crystal structure of attapulgite. Atomic force microscopy (AFM) and
transmission electron microscopy (TEM) images of acid-treated attapulgite nano-rods
revealed a rough surface with perforations.

1
Chapter 1 Introduction

1.1 Polymer Nanocomposites
A polymer is a large molecule composed of repeating structural units typically
connected by covalent chemical bonds. Due to the wide spectrum of properties accessible
in polymeric materials [1], they have become an essential and ubiquitous role in daily life
- from plastics and elastomers to natural biopolymers such as DNA and proteins.
However, the inherent properties of the polymer alone can’t be satisfied with the
demands of applications with the development of time. Thus, a stronger or stiffer material
was incorporated with neat polymer to improve the mechanical and thermal performance
and become a new type of material “composite.”
The development of composite materials has broadened the range of material
properties by expanding possible combinations. Composite materials are made from two
or more components with significantly different physical or chemical properties which
remain separate and distinct on a macroscopic level within the finished structure. Many
composites offer property trade-offs as well. The magnitude of the property change in
composites not only depends on the filler composition but also includes the particle size,
shape, and surface chemistry [2]. Particle size, shape, and bonding condition between
filler and polymer matrix are all important factors in determining the performance of
composite.
As new technologies continue to increases the demands on the high performance
of polymeric materials, the traditional polymeric composites can not meet these
requirements. As a result, new polymeric composite materials with nano-scale materials
2
were developed. Such materials not only enhance the mechanical properties but also
provide structural integrity, as well as serve additional functions, for example conduct
electricity, gas barrier, and optical properties. Hussain et al.[3] also ascribed the dramatic
effects on the physical properties of polymers to the transition from micro-particles to
nano-particles. This transition created large surface area that nano-scale fillers have for a
given volume.
Generally speaking, polymer nanocomposites are composites materials that
consist of a polymer matrix with well dispersed fillers, which have at least one dimension
below 100nm. Since organic-inorganic nanocomposites based on layered silicates were
first introduced by Toyota Research Center (combining nylon and montmorillonite) [4-6],
nanofillers have been widely studied, creating the field of polymer nanocomposites [7-
10]. For example, the introduction of small amount of nanofillers has been shown to
cause increases in strength, stiffness, and heat resistance of polymers [11-15].
According to the number of dimensions of the dispersed particles in the
nanometer range, the nanocomposites can be classified into following three types. (1)
Three dimensions are in the order of nanometers, forming a cubic structure. This includes
isodimensional nanoparticles (e.g. spherical silica nanoparticles) [16-17], semiconductor
nanoclusters [18], and many others [19]. (2) Two dimensions are in the order of
nanometers and the third is larger, forming an elongated structure. This includes carbon
nanotubes [20] or cellulose whiskers [21] which have been extensively studied as
reinforcing nanofillers with excellent properties. (3) One dimension in the order of
nanometer only. In this case, the fillers are present in the form of sheets of one to a few
nanometers in thick and hundreds to thousands nanometers in long. This family of
3
composites can be named the polymer-layered crystal nanocomposites, which are almost
exclusively obtained by the intercalation of the polymer inside the galleries of layered
host crystal [22]. These crystalline hosts can be natural or synthetic, for example the
graphite [23], graphite oxide [24], layered double hydroxides [25], and layered silicates
[22]. Among them, clay and layered silicates have been most widely studied because the
clay minerals are abundance in the earth and easy to process.

1.2 Layered Silicates
The commonly used layered silicates in nanocomposites are the structural family of
2:1 phyllosilicates, which consists of two silica tetrahedral layers and one alumina or
magnesia octahedral layer. The silica tetrahedral groups (SiO
4
) linked together to form a
hexagonal network of the repeating units of composition Si
4
O
10
. The alumina or
magnesia octahedral layer is composed of two sheets of close packed oxygens or
hydroxyls between which the coordinated aluminum or magnesium atoms are imbedded.
The combination of two tetrahedral sheets and one octahedral sheet is referred to as a unit
layer. The unit layer thickness is around 1 nm and the lateral dimensions may range from
100 nm to 1 μm.
Within the tetrahedral sheet of layered silicate, the Si
4+
is partly replaced by Al
3+
.
In the octahedral sheet, the Al
3+
can be replaced by Mg
2+
or Fe
2+
without complete filling
of the third vacant octahedral position. When an atom of lower positive valence replaces
one of higher valence, the clay unit will show a net negative charge and naturally
compensated by the adsorption of the cations on the layer surface. This type of silicate
4
can be characterized by a moderate negative surface charge known as the cation
exchange capacity (CEC, mequiv/100g). Due to the charge of the layer varies from layer
to layer, it must be considered as an average value over the whole crystal.
Montmorillonite (MMT) is the most commonly used layered silicates in
polymer/clay nanocomposites. The unit layers are regularly stacked and bound together
by weak inter-atomic forces to form 2D plate-like layers as shown on Fig 1.1 MMT is
described by the chemical formula (Al
2-y
Mg
y
)(Si
4-x
Al
x
)O
10
(OH)
2
,M
y
+
in which M
+
is the
exchangeable cation (Na
+
, Ca
2+
, Li
+
), which can be substituted by an organic cation via
ion exchange reaction in the galleries, and y is the degree of substitution [26]. Due to
isomorphic substitution between metals, negative charged MMT can adsorb cations with
an equal electric quantity to balance the potential in the galleries. MMT is expandable
clay with specific surface area as large as 800 m
2
/g. The whole surface of layers
including the internal surface and external surface can be hydrated and ion-exchanged
[27].
Another layered silicate, attapulgite (ATT), has a similar structure to MMT but 1D
fiber-like morphology also has been considered as a new candidate to reinforce polymer
because of the good mechanical strength and thermal stability. ATT is a hydrated
magnesium aluminum silicate with chemical formula Si
8
O
20
Al
2
Mg
2
(OH)
2
(OH
2
)
4
．4H
2
O
[28]. Although ATT features ribbons of a 2:1 phyllosilicate structure, it differs from other
layered silicates in that it lacks continuous octahedral sheets. Each ribbon is linked to the
next by inversion of SiO
4
tetrahedra along a set of Si-O-Si bonds which extend parallel to
the X-axis, forming rectangular channels that contain zeolitic water (Fig 1.2) [29].
However, the Si-O-Si bonds are weak and are easily broken by shear stress to form
fibrous crystals [29]. The average external and internal surface areas have been estimated
to be 300 and 600 m
2
/g, respectively [29]. Because of the unique morphology and
structure, ATT has been used for various commercial applications, such as adsorbents,
catalysts, rheological agents, and fillers [29-33].

Fig 1.1 Structure of montmorillonite [26].

5

Fig 1.2 Structure of attapulgite [29].

6
7
1.3 Dispersion of Layered Silicates
Three generalized morphologies can be introduced when a layered silicate is
associated with a polymeric system. Those morphologies are phase separation,
intercalation, and exfoliation, respectively (Fig 1.3).
First of all, the phase separated morphology is considered as traditional
microcomposites. In these cases, the polymeric molecules are unable to penetrate into the
galleries between silicate layers. The spacing of layered silicates remains the same as
original status. As a result, the properties of phase separated microcomposites stay in the
same stage as traditional microcomposites. Beyond the classical family of composites,
both intercalated and exfoliated morphologies are considered as nanocomposites. The
intercalated structure is caused by one or several polymer chains inserted into the
galleries between silicate layers resulting in a well ordered multilayer morphology. The
exfoliated structure which consists of individual nano-scale silicates uniformly dispersed
in a continuous polymer matrix and has the largest spacing between the silicate layers.
Besides, the exfoliation has become an idealized morphology in the nanocomposites
because it is expected to lead to dramatic improvements by the largest interface between
polymer and silicate layers.
In addition, there is an intermediate morphology between the intercalation and
exfoliation, which is called partial exfoliation. The partial exfoliated structure can be
found with the dispersion of small stack of 2-4 layers in the polymer matrix. In practice,
however, the non-expandable layered silicates (e.g. attapulgite or sepiolite) also can be
dispersed in polymer matrix in nano-scale fashion but can’t be described by previous
classification. Therefore, we will use the adequate description of nano-scale
morphologies (e.g. uniform dispersion) for the no-expandable layered silicates in this
contribution.

Phase separated
Microcomposite
Intercalated
Nanocomposite
Exfoliated
Nanocomposite
＋
(a)
(b) (c)
Layered silicate Polymer
Fig 1.3 Scheme of three types of morphologies in polymer-layered silicate composites. (a)
Phase separated microcomposites (b) intercalated nanocomposite (c) exfoliated
nanocomposite.

8
9
1.4 Research Objectives
Academic and applied interests provided motivation for these studies. Efforts were
spent on synthesis of nanoclay-reinforced polyurethane elastomers with 2D plate-like
layered silicate (montmorillonite), as well as 1D fiber-like layered silicate (attapulgite).
Commercial montmorillonite (Cloisite 30B and Nanomer I.30E) and self organophilically
modified attapulgite were used to reinforce polyurethane elastomers. The nanocomposites
were characterized and confirm with nano-scale dispersion of silicates by using wide-
angle X-ray diffraction (WAXD), transmission electron microscopy (TEM), and scanning
electron microscopy (SEM). Thermal properties were investigated via differential
scanning calorimetry (DSC), and thermogravimetric analysis (TGA). Dynamic
mechanical analysis and tensile tests were performed to understand the effects of
nanocomposites morphology on mechanical properties. Fourier transform infrared
spectroscopy (FTIR) has been used to characterize the self organically modified
attapulgite and investigate the effects of nano-clay additions on chemical bonding. These
studies of PU/clay nanocomposites have advanced our understanding of the physics of
polymer nanocomposites, and have provided a scientific basis for future applications of
polymer-layered silicate nanocomposites.
Another basic research on non-expandable layered silicate, attapulgite, by heat and
acid treatments were also investigated. Thermal analysis of the untreated attapulgite
revealed dehydration and dehydroxylation, and clarified the controversy surrounding the
specific composition of water molecules (e.g., bound water versus zeolitic water) released
during successive stages of heating. The dehydration and dehydroxylation during heat
treatment of attapulgite led to structural changes, as demonstrated by FTIR spectra and
10
WAXD patterns. FTIR spectra also revealed that acid treatment hydroxylated the heat-
treated attapulgite. Atomic force microscopy (AFM) and transmission electron
microscopy (TEM) images demonstrated that acid treatment caused surface roughening
and pitting, although the fibrous structure was retained. This work provides insight into
physicochemical changes resulting from thermal and acid treatments. This understanding
is necessary to extend applications of attapulgite to specific polymer systems and to
heterogeneous catalytic reactions.

11
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22. Alexandre M., P. Dubois / Materials Science and Engineering 2000, 28, 1.
23. Shioyama H., Polymerization of isoprene and styrene in the interlayer spacing of
graphite, Carbon 1997, 35, 1664.
24. Matsuo Y., Tahara K., Sugie Y., Synthesis of poly(ethylene oxide)-intercalated
graphite oxide, Carbon 1996, 34, 672.
25. Oriakhi C.O., Farr I.V., Lerner M.M., Thermal characterization of poly(styrene
sulfonate)/layered double hydroxide nanocomposites, Clays and Clay Minerals
1997, 45, 194.
26. Pinnavaia, T. J., Science 1983, 220, 365.
27. Ke Y.C., Stroeve P., Polymer-Layered Silicate and Silica Nanocomposites,
Elsevier, Amsterdam 2005.
28. Bradley Amer., Mineral 1940, 25, 405.
29. Galan E., Clay Miner 1996, 31,443.
30. Cao Jian-Liang., Shaoa Gao-Song., Wanga Yan., Liua Yuping., Yuan Zhong-
Yong., Catalysis Communications 2008, 9, 2555.

31. Alexander Neaman., Arieh Singer., Applied Clay Science 2004, 25, 121.
32. Li An., Wang Aiqin., Eur Polym J 2005, 41, 1630.
33. Shen Liang., Lin Yijian., Du Qiangguo., Zhong Wei., Yang Yuliang., Polymer
2005, 46, 5758.
13
Chapter 2 2-D Nanoclay-Reinforced Polyurethane [1]

2.1 Motivation
Segmented polyurethane (PU) elastomer is a commercially important with
widespread applications. Its attributes include high abrasion resistance, shock absorption,
flexibility, elasticity, and resistance to chemicals [2]. These properties originate primarily
from the tendency of PU to form discrete regions of micro-domains. The linear chain
structure of PU can be expressed in the form of (A-B)
n
, where the hard segment A is
composed of low molecular weight diol or diamine (chain extender) with diisocyanate,
and the soft segment B is composed of high molecular weight polyester or polyether
polyol. Because of the different chemical structure of the hard and soft segments,
repulsive interactions and thermodynamic incompatibility lead to microphase segregation
[3], and formation of hard- and soft-segment domains. Moreover, the hard segments form
micro-domains by intermolecular hydrogen bonding in the PU [4]. Two approaches are
commonly used to enhance the mechanical properties and thermal stability of
polyurethanes. The first is to alter the molecular structure of the polyurethane (e.g., molar
ratio of ingredients, molecular weight of soft segment, etc.), while the second approach is
to introduce inorganic fillers into the polyurethane matrix.
Organic-inorganic nanocomposites based on layered silicates were first introduced
by Toyota Research Center (combining nylon and montmorillonite) in 1993 [5-7]. Since
then, the use of inorganic nanofillers in polymers has developed into the field of
polymer/inorganic nanocomposites. In these materials, the introduction of small amounts
of nanofillers can result in marked increases in strength, stiffness, and heat resistance
14
compared with unreinforced (neat) polymer counterparts [8-12]. One common nanofiller
is montmorillonite (MMT) clays, 2:1 layered silicates composed of two silica tetrahedral
sheets and one central octahedral sheet of alumina. MMT additions can significantly
improve mechanical and thermal properties for polymer/inorganic nanocomposites,
including polyimide, polycaprolactone, and polypropylene [13-15]. The layers are
regularly stacked and bound together by weak inter-atomic forces. MMT is described by
the chemical formula (Al
2-y
Mg
y
)(Si
4-x
Al
x
)O
10
(OH)
2
,M
y
+
in which M
+
is the exchangeable
cation (Na
+
, Ca
2+
, Li
+
), which can be substituted by an organic cation via ion exchange
reaction in the galleries, and y is the degree of substitution [16].
Because of the poor compatibility of the hydrophilic inorganic silicates with
hydrophobic polymer matrices, organic modification of the silicates has become the most
common and efficient way to increase the hydrophobicity and thereby improve the
compatibility of the composite pair. In this organic modification, the cations in the layer
galleries are exchanged with cationic organic compounds (e.g. alkylammonium and
alkylphosphonium). After the organic modification, the MMT not only becomes
hydrophobic and thus more compatible with the polymer, but the spacing of the galleries
is expanded, facilitating intercalation and/or exfoliation in polymer/MMT
nanocomposites. Exfoliation of the nanoclays generally results in improved mechanical
properties [17-18].
Early work with PU/organically modified, layered silicate nanocomposites with
intercalated morphology showed a large enhancement in tensile properties [19]. Similar
work was followed by several studies of nanoclay composites using solution and bulk
polymerization methods to intercalate the clay with soft segment polyols prior to reaction
15
with isocyanates [20-26]. For example, Tien and Wei [27-28] prepared PU/clay
nanocomposites by allowing pre-polymer chain terminated with –NCO groups to react
with primary ammonium modifier carrying 1-3 –CH
2
OH groups. They reported increases
in tensile properties, glass transition temperature, and resistance to dynamic mechanical
and thermal degradation with the addition of small amounts of organophilic MMT.
However, the effects of the different organic modifiers and molecular weight of soft
segments in PU/clay composites on composite properties have been rarely reported. In
this chapter, the effects of two functionally organic-modified montmorillonites, MMT-
30B (Cl
-
N
+
(CH
2
CH
2
OH)
2
(CH
3
)T modified MMT) and MMT-I30E (CH
3
(CH
2
)
17
NH
3
+
Cl
-

modified MMT), and three molecular weights of the polyether soft segments in PU, on
the properties of PU/MMT nanocomposites were investigated.

2.2 Experimental
2.2.1 Materials
The nanoclays used in the composite synthesis included two MMTs (Cloisite 30B,
from Southern Clay Products, Inc, and Nanomer I.30E, from Nanocor
®
, Inc). Cloisite
30B is a natural MMT modified by quaternary ammonium salt. The quaternary
ammonium ion has the structure, Cl
-
N
+
(CH
2
CH
2
OH)
2
(CH
3
)T, where T represents an
alkyl group of approximately 65% C
18
H
37
, 30% C
16
H
33
, and 5%C
14
H
29
[29]. Nanomer
I.30E, on the other hand, is an octadecylammonium modified MMT. The
octadecylammonium modifier has the structure CH
3
(CH
2
)
17
NH
3
+
Cl
-
. All nanoclays were
dried in a vacuum oven at 80
o
C for 48h prior to use. After dehydration, PU was prepared
16
using 1,4 butanediol (BDO, Avocado Research Chemicals Ltd.) and polytetrahydrofuran
(PTHF, Sigma-Aldrich, M.W 1000, 2000, and 2900), which are the hydroxyl-terminated
monomer and oligomer, respectively. The 1,4 butanediol was dried over calcium hydride
for 48h and then was vacuum distilled. Polytetrahydrofuran was dehydrated in a vacuum
oven at 60
o
C for 48h. The diisocyanate, MDI (98%, Sigma-Aldrich), was purified by
filtration of molten MDI liquid at 70
o
C. Dimethylformamide (DMF, Sigma-Aldrich) was
dehydrated and used as a solvent in the polymerization reaction.

2.2.2 Synthesis of PU elastomer and PU/MMT nanocomposites
PU elastomer was synthesized with a molar ratio of 4:1:2.64 (MDI-to-PTHF-to-
1,4BDO), using excess diisocyanate to produce partially cross-linked networks and –
NCO end groups [30]. First, the oligomeric polyol (PTHF) was dissolved in DMF and
reacted with MDI at 80
o
C for 30min to obtain a pre-polymer in a round bottom flask with
continuous stirring under vacuum. Second, the chain extender (1,4 butanediol) was added
to build up the polyurethane network, and the mixture was allowed to react at 80
o
C for an
additional 2 min. In initial experiments, the network building reaction was active within 2
min (after the addition of 1,4 BDO). The reaction was not completed in 2 min, but if the
reaction was allowed to continue beyond 2 min, the viscosity increased rapidly, making it
impossible to pour out of the round bottom flask. Finally, the mixture was poured into a
Teflon
®
mold and cured at 80
o
C for 24h to obtain the PU elastomers. To prevent moisture
absorption, the reaction was performed under vacuum and all chemicals were dehydrated
beforehand.
17
Six batches of PU/MMT nanocomposites with different soft-segment molecular
weight of PTHF (M.W 1000, 2000, and 2900) but identical nanoclay contents (2.5 wt%)
were also prepared with identical molar ratios, as shown in Table 2.1. To prepare the
PU/MMT nanocomposites, two organically modified clays, including Cloisite 30B
(MMT-30B) and Nanomer I.30E (MMT-I30E) were used. The polyol was dissolved in
DMF and reacted with excess MDI at 80
o
C under vacuum for 30min to form the pre-
polymer. Next, the organically modified MMT was added to the PU pre-polymer, which
was dissolved in the DMF solvent, using a high-speed dual-axis mixer (HM-500,
Keyence, USA) for 30min. Subsequently, the chain extender 1,4 butanediol was added to
the pre-polymer/MMT system and stirred vigorously for 2 min to complete the reaction.
Finally, the viscous polymer was poured into a Teflon
®
mold and cured at 80
o
C for 24h
to form PU/MMT nanocomposites. Figure 2.1 illustrates the steps in the synthesis route
and the associated chemical structures. The samples will be referred to as PU1000 series,
which includes PU1000, PU1000/MMT-30B and PU1000/MMT-I30E for soft-segment
molecular weight of 1000; PU2000 series, which includes PU2000, PU2000/MMT-30B
and PU2000/MMT-I30E for soft-segment molecular weight of 2000; and PU2900 series,
which includes PU2900, PU2900/MMT-30B and PU2900/MMT-I30E for soft-segment
molecular weight of 2900.

18
Table2.1 The composition and hard segment content of PU and PU/MMT
nanocomposites.
Samples Soft Segment
MW (g/mol)
PTHF
(mole)
MDI
(mole)
BDO
(mole)
Hard segment
content
*
(wt%)
PU1000 1000 1 4 2.64 55.86
PU1000/MMT-30B 1000 1 4 2.64 55.86
PU1000/MMT-I30E 1000 1 4 2.64 55.86
PU2000 2000 1 4 2.64 38.72
PU2000/MMT-30B 2000 1 4 2.64 38.72
PU2000/MMT-I30E 2000 1 4 2.64 38.72
PU2900 2900 1 4 2.64 30.00
PU2900/MMT-30B 2900 1 4 2.64 30.00
PU2900/MMT-I30E 2900 1 4 2.64 30.00
* (W
MDI
+W
BDO
)/(W
PTHF
+W
MDI
+W
BDO
)

Figure 2.1 Synthesis and chemical structure of PU/MMT nanocomposite.

19
20
2.2.3 Characterization
Wide-angle X-ray diffraction (WAXD) was performed to analyze the organic
modified MMT powders (MMT-30B and MMT-I30E) and the PU/MMT nanocomposites
(Bruker X-ray diffractometer, model D8 Advance). The analysis was carried out with
0.154 nm radiation, and a 2 θ scan from 2.5° to 30°.
The dispersion of the MMT-30B and MMT-I30E in the PU matrix was observed
using transmission electron microscopy (TEM, Philips EM420) at an acceleration voltage
of 120kV. Ultra-thin samples were sectioned by cryogenic ultramicrotoming.
IR spectroscopy was performed on the nanoclays and the nanocomposites (Nicolet
4700). For the nanoclays, KBr/ MMT disks were prepared and used to obtain IR spectra
for the MMT. The organically modified MMT clay (0.6 mg) and the KBr (200 mg) were
mixed and pressed to make disks. The IR spectra of the MMT were obtained in
transmission mode by FTIR spectroscopy (Nicolet 4700), and 128 scans were recorded
with a resolution of 4 cm
-1
. For the PU and PU/MMT nanocomposites, IR spectra were
obtained using the attenuated total reflection (ATR) technique, and 32 scans were
collected at a resolution of 4 cm
-1
.
The glass transition temperatures of the soft-segment components in polyurethane
and PU/MMT nanocomposites were measured by differential scanning calorimetry (DSC
2920, TA Instruments). Measurements were performed by scanning twice to avoid signal
noise and to erase thermal history. In the first scanning, the sample was heated to 250°C
at 10°C/min, then cooled to -65°C. In the second scanning, the sample was heated to 300°
at a heating rate of 10°/min. The second scanning was used to record transition
temperatures.
Tensile strength, modulus, and elongation at break were measured using a
universal testing machine (Instron 8531). Tests were performed on dog-bone shaped
specimens using a crosshead speed of 10 mm/min at room temperature, in accordance
with ASTM D 638-94b.

2.3 Results and Discussion
2.3.1 Morphologies of organic MMTs and PU/MMT nanocomposites
The WAXD patterns of MMT-30B, MMT-I30E, and PU/MMT nanocomposites
are shown in Fig 2.2. The (001) diffraction peak of MMT-30B and MMT-I30E appear at
diffraction angle 2 θ=4.85° and 3.45°, respectively. Using the Bragg equation, the d
001

spacing of MMT-30B was determined to be 1.82 nm, while d
001
of MMT-I30E was 2.57
nm. Comparing the d spacings for MMT-30B and MMT-I30E, MMT-I30E showed a
larger d spacing. The larger d spacing is associated with the higher cation exchange
capacity for MMT-I30E. In addition, the arrangement of modifier molecules in MMT
also affects the layer separation. Note that a small shoulder appears near 5° in the MMT-
I30E pattern, and the small shoulder is attributed to the presence of nanoclay particles
without organic modification. Also, the (001) peak for MMT-30B in the PU/MMT-30B
nanocomposite pattern is not present, indicating that exfoliation of the silicate layer
structure of the organo-clay has occurred in PU.
Similar results were obtained for the (001) peaks of PU1000/MMT-I30E and
PU2000/MMT-I30E, although PU2900/MMT-I30E showed a broad hump near 2 θ=3°.
The absence of the basal reflection in the PU1000/MMT-I30E and PU2000/MMT-I30E
nanocomposites also suggests extensive exfoliation of the nano-clays in the PU matrix.
21
22
However, the PU2900/MMT-I30E nanocomposite exhibited a small peak near 2 θ=3°,
suggesting an intercalated morphology with partial exfoliation. The diffraction peak near
2θ=5° for MMT-I30 and PU/MMT-I30E nanocomposites was attributed to the
incompletely modified MMT-I30E. The XRD patterns for PU/MMT-30B and PU/MMT-
I30E nanocomposites may not fully reveal the levels of exfoliation or intercalation,
because the concentration and order of the nano-clay can influence the XRD patterns [31].
Nevertheless, despite some limitations, the XRD results provided a useful approximation
to the nanostructure from a global perspective. More detailed (and local perspective) was
obtained from TEM analysis, which provided direct images of the morphology and
spatial distribution of clay platelets.
Figure 2.3 shows TEM images of both PU2900/MMT-30B and PU2900/MMT-
I30E nanocomposites with 2.5 wt% nano-clay loadings. The images show intercalated
and partially exfoliated morphologies. The PU2900/MMT-30B nanocomposite, however,
exhibits more extensive exfoliation than PU2900/MMT-I30E, which retains ordered
layers in which the galleries are extensively expanded. From these observations of
exfoliated nanoclay, the thickness of plate-like MMT-30B and MMT-I30E is ~1 nm,
while the aspect ratio (length-to-thickness) is ~150-200 in both PU2900/MMT-30B and
PU2900/MMT-I30E nanocomposites.
The organophilic MMT had an expandable layer structure and a plate-like shape.
When the chemical modifier reacted with the PU chains, the clay gallery spacings were
expanded or peeled, resulting in a mixture of intercalation and partial exfoliation. This
observation indicates that the MMT-30B had a stronger chemical driving force for
exfoliation than the MMT-I30E when reacted with the -NCO end groups of PU. The
23
combination of XRD patterns and morphologies observed by TEM confirm that both
MMT-30B and MMT-I30E formed intercalated and partially exfoliated morphologies in
PU. However, MMT-30B has two hydroxyl chain ends that are more reactive to the NCO
groups of PU than MMT-I30E, in which the ammonium cation has low reactivity to the
NCO groups of PU and is much less nucleophilic than the amine group. Thus,
PU2900/MMT-30B exhibited more extensive exfoliation than PU2900/MMT-I30E.

24
0 2 468 1
PU2900/MMT-30B
PU2000/MMT-30B
PU1000/MMT-30B
MMT-30B

Intensity
2 Theta(degree)

(a)
246 8 10
PU2900/MMT-I30E
PU2000/MMT-I30E
PU1000/MMT-I30E
MMT-I30E

Intensity
2 Theta(degree)

(b)

Figure 2.2 WAXD patterns for (a) PU/MMT-30B and (b) PU/MMT-I30E
nanocomposites.

(a)

(b)
Figure 2.3 TEM images of (a) PU2900/MMT-30B and (b) PU2900/MMT-I30E
nanocomposites.
25
26
2.3.2 FTIR characterization
The FTIR spectra of MMT-30B and MMT-I30E reveal characteristic bands of Al-
OH stretching at 3627 cm
-1
and a broad band of H-bonded H-O-H stretching at 3429 cm
-1

(Figure 2.4). The two small peaks in the spectrum of MMT-I30E near 3300 cm
-1

correspond to N-H stretching vibrations, an indication that the modifier in the clay has
ammonium groups. Two peaks at 2926 and 2854 cm
-1
correspond to methylene groups in
hydrocarbon chains of the organic ammonium ions present in both MMT-30B and MMT-
I30E. The peak at 1620 cm
-1
corresponds to H-O-H bending vibrations [32]. The
characteristic bands of MMT at 1049, 523, and 465 cm
-1
are the stretching vibration of Si-
O bonds, the bending vibration of Si-O-Al, and the Si-O-Si bending vibration,
respectively.
The FTIR spectra for PU and the PU/MMT nanocomposites are shown in Figure
2.5. Although the spectra are qualitatively similar, deviations in peak positions and
relative peak intensities of certain bands were evident among the samples that were
synthesized with identical molar ratios but with different soft-segment molecular weight
and functional MMT. In particular, consider the NH and C=O stretching, which are the
two most interesting regions in this study. The NH absorption peak located at 3322 cm
-1

corresponds to hydrogen-bonded NH groups of urethane linkages [33]. This hydrogen
bonding can form with hard-segment carbonyl and with soft-segment ether linkages.
The carbonyl bands can be divided into three regions. In Fig 2.5, the peak at 1729–
1731 cm
-1
is assigned to free urethane carbonyl, the peak at 1700–1704 cm
-1
is the
hydrogen-bonded urethane carbonyl, and the peak at 1643–1648 cm
-1
is the hydrogen-
bonded urea carbonyl [34]. For each sample, the absorbencies of N-H and C=O reflect
27
the respective concentration of these functional groups. The absence of the absorbance
due to stretching of NCO groups at 2270 cm
-1
[33] indicates that the NCO end-groups
were completely reacted after the synthesis. The splitting of the carbonyl stretching
absorption peak in the PU and PU/MMT nanocomposites was attributed to microphase
separation on both materials [33]. Overall, the main features of the spectra for the
PU/MMT nanocomposites were the same as those from the neat PU, indicating similar
bond characteristics. This observation supports the conclusion that there were no major
chemical structural changes in PU/MMT nanocomposites.
Peak area ratios in IR absorption spectra provided additional insights into the
effects of the nano-clay additions on chemical bonding. Table 2.2 shows the peak area
ratios for the free carbonyl at 1731 cm
-1
(A
fCO
) to the hydrogen-bonded carbonyl at 1700
cm
-1
(A
hCO
), and for the hydrogen-bonded NH at 3322 cm
-1
(A
NH
) to the CH stretching
(A
CH
) between 2811 and 2984 cm
-1
are summarized on Table 2.2. The value of A
CH
was
used as an internal standard. As shown in Table 2.2, the A
NH
/A
CH
ratio for each PU series
was insignificantly affected by the presence of MMT. The higher A
NH
/A
CH
ratio for the
PU series of a lower soft-segment molecular weight indicates that the content of the hard
segment was greater. This observation is consistent with results shown in Table 2.1. (This
also was reflected in the tensile properties, where the PU1000 series showed greater
tensile strength and modulus than the PU2000 and PU2900 series.) In addition, the A
hCO
/
A
fCO
ratio was decreased by additions of MMT in the PU series. This result indicates that
the content of the hydrogen-bonded carbonyls was decreased by MMT-30B and MMT-
I30E to a greater extent than that of the free carbonyls.
28
Comparing the two nano-clays, MMT-30B had a smaller effect on the A
hCO
/ A
fCO

ratio than MMT-30B. The smaller effect of MMT-30B is attributed to the two hydroxyl
groups in the organic compound 30B that was used to modify the MMT. The more
extensive hydrogen bonding between PU and the organic compound in MMT-30B was a
major cause of the more extensive exfoliation in the PU/MMT-30B nanocomposite (as
observed in the TEM images in Figure 2.3). A second possible contributing factor is that
the MMT loading (2.5 wt%) was insufficient to affect the A
hCO
/ A
fCO
ratio in the
PU/MMT nanocomposites. In related work, Pattanayak et al [35] reported that when 5
wt% MMT30B was added to PU, the hydrogen-bonded carbonyl increased. From their
observations, they inferred that the urethane carbonyls reacted with residual CH
2
CH
2
OH
groups of the quaternary ammonium ions and formed additional hydrogen bonds.

4000 3500 3000 2500 2000 1500 1000 500
3627
465
523
2854
2926
1620
(a)
(b)
Transmittance(%)
Wavenumber(cm
-1
)

Figure 2.4 FTIR spectra of (a) MMT-30B and (b) MMT-I30E.

29

Figure 2.5 FTIR spectra of (1) PU1000 series, (2) PU2000 series, and (3) PU2900 series
in the range from 1650 to 1800 cm
-1
. (a: PU; b: PU/MMT-I30E; c: PU/MMT-30B.)

30
Table 2.2 The area ratio of the IR absorption peaks of N–H stretching (A
NH
) and C–H
stretching (A
CH
) and the area ratio of the IR absorption peaks of the hydrogen-bonded
C=O (A
hCO
) and free C=O (A
fCO
).
Samples A
NH
/A
CH
A
hCO
/A
fCO
PU1000 0.44 19.56
PU1000/MMT-30B 0.37 17.57
PU1000/MMT-I30E 0.37 13.23
PU2000 0.22 5.14
PU2000/MMT-30B 0.2 4.21
PU2000/MMT-I30E 0.17 1.72
PU2900 0.08 1.76
PU2900/MMT-30B 0.07 0.94
PU2900/MMT-I30E 0.11

0.6

31
32
2.3.3 Tensile properties
The tensile yield strength (YS), Young’s modulus (E) and elongation at break for
PU and PU/MMT nanocomposites are presented in Table 2.3. For a given soft-segment
molecular weight, the yield strength, modulus, and elongation at break of PU were all
increased by the clay additions, and MMT-30B resulted in a larger increase than MMT-
I30E. Table 2.3 shows that the Young’s modulus, tensile yield strength, and elongation at
break of the PU2900/MMT-30B nanocomposite increased by 142%, 36%, and 49%
respectively, while increases for the PU2900/MMT-I30E nanocomposites were more
modest - 21%, 9% and 11%, respectively. The PU1000 series resulted in the greatest
Young’s modulus and yield strength values among the three PU series, as expected,
because of the highest hard-segment content synthesized within the PU1000 series.
However, the value of elongation at break in PU1000 series is the lowest among the three
PU series because of the increased hard-segment content and the associated loss of
elasticity [36-37].
The finding that, for a given molecular weight of soft segment, MMT-30B resulted
in greater increases in YS and E than MMT-I30E is explained by the stronger interactions
and higher reactivity of the two hydroxyl chain ends (MMT-30B) with the NCO groups
of PU than the ammonium cation (MMT-I30E) to the NCO groups of PU. In addition,
because of the large specific surface area of organophilic MMT, a large interface zone
between MMT particles and PU matrix was created in both PU/MMT-30B and
PU/MMT-I30E nanocomposites. As a result, when load was applied to the PU/MMT-
30B and PU/MMT-I30E nanocomposites, it was transferred effectively from the matrix
to the nearby nanoclay particles [38]. In this way, the load was transferred efficiently via
33
the large interface zone for PU/MMT-30B and PU/MMT-I30E nanocomposites and the
tensile properties increased substantially.
From the TEM observations (Fig 2.3), the PU/MMT-30B nanocomposites were
more exfoliated than the PU/MMT-I30E nanocomposites. The greater extent of
exfoliation created a larger, thicker interface zone. Brinson et al [39] have argued that
this interface zone is effectively hardened relative to the neat polymer. To summarize,
these results illustrate that uniform dispersion of the clay particles into the polymer
matrix is necessary but not sufficient to enhance mechanical properties, and that surface
or chemical modifications that lead to an improved interfacial adhesion between filler and
matrix are also necessary.

Table 2.3 Tensile properties of PU and PU/MMT nanocomposite.

Samples Yield Strength
(MPa)
Young’s Modulus
(MPa)
Elongation at
break (%)
PU1000 2.08 57.10 286
PU1000/MMT-30B 2.42 102.00 423
PU1000/MMT-I30E 2.17 59.17 338
PU2000 1.53 21.83 351
PU2000/MMT-30B 1.82 35.16 505
PU2000/MMT-I30E 1.62 23.80 393
PU2900 1.08 10.10 416
PU2900/MMT-30B 1.47 24.40 620
PU2900/MMT-I30E 1.18 12.21 462

34
2.3.4 DSC measurements
The DSC results for neat PU and PU/MMT nanocomposites with different soft-
segment molecular weights are summarized in Table 2.4. Because the structure of PU
includes hard and soft segments, two melting temperatures are expected. However, the
DSC curve (not shown) revealed only one peak, and this was associated with the melting
temperature of the soft segment (T
m
, SS). The peak associated with the melting
temperature of the hard segment (T
m
, HS) was absent from the DSC curve [40-42]. The
absence was attributed to the inactive movement of hard segment, which has a small
[43], and to the widely dispersed HS micro-domains within the PU matrix [44]. Cp Δ
The melting temperatures for the PU2000 and PU2900 series in the region between
16 and 32
o
C (Table 2.4) are characteristic of ordered soft-segment structures. The values
also indicate that polytetrahydrofuran crystallized in both neat PU and PU/MMT
nanocomposites with soft-segment (SS) molecular weights (MW) of 2000 and 2900. For
these SS MWs, the enthalpy changes of PU2000 were increased in the presence of MMT-
30B and MMT-I30E by about 24.2 and 14.6 J/g, respectively. A similar trend also
occurred in the PU2900/MMT-30B and PU2900/MMT-I30E nanocomposites, which
showed increases of 22.4 and 12 J/g, respectively. These increases in melting temperature
and SS enthalpy in PU2000 and PU2900 were attributed to the enhancement of
crystallinity by the clays acting as nucleating agents [45]. However, in the PU1000 series,
no SS melting transition (T
m
, SS) was detected. This phenomenon indicated that the SS in
the PU1000 series did not crystallize because of insufficient SS content. As shown in
Table 2.4, the magnitude of the enthalpy change increased with increasing molecular
weight of the soft segment for both PU2000 and PU2900 series. This result indicates that
35
36
the degree of SS crystallinity is greater for the high MW SS, i.e., a higher content of the
soft segment in the PU.
Korley et al [46] reported that the presence of crystallites within the soft segment
could absorb strain energy during deformation. Thus, semicrystalline soft segments
would act as a load-bearing phase during deformation. As shown by the DSC
measurements, the MMT-30B and MMT-I30E acted as nucleation seeds, increasing the
degree of crystallinity of the soft segment. In addition, the rigid crystalline MMT clay
also contributed to the increased hardness by virtue of simple reinforcement. As a result,
the PU/MMT-30B and PU/MMT-I30E nanocomposites exhibited greater tensile strength
and modulus values than neat PU because of the crystallization induced by the MMT clay.
Comparing the enthalpy changes in PU/MMT nanocomposites with similar SS
MWs, the PU/MMT-30B nanocomposites exhibited greater crystallinity than the
PU/MMT-I30E nanocomposites in the PU2900 and PU2000 series. The greater
crystallinity can be explained by the fact that the -CH
2
CH
2
OH groups attached to the
MMT-30B had stronger interactions and higher reactivity with PU than the –NH groups
in MMT-I30E. While the -CH
2
CH
2
OH groups reacted with hard segments in the PU, the
hard segment mobility was restricted by the MMT-30B clay. In contrast, the soft
segments had greater mobility and free volume to produce greater crystallization for the
PU/MMT-30B nanocomposites. However, as discussed previously, the soft segments in
the PU1000 series did not crystallize, yet resulted in superior tensile properties relative to
the PU2000 and PU2900 series because of the highest hard-segment content synthesized
within the PU1000 series. This supports the assertion that the hard segment concentration
affected the tensile properties more strongly than the degree of SS crystallinity for the
PU1000 series.

Table 2.4 Melting temperature and enthalpy of soft segment in PU and PU/MMT
nanocomposites.

Samples T
m, SS
(°C)
a
ΔH (J/g)
PU1000 N/A N/A
PU1000/MMT-30B N/A N/A
PU1000/MMT-I30E N/A N/A
PU2000 16.13 12.97
PU2000/MMT-30B 19.82 37.13
PU2000/MMT-I30E 20.58 27.57
PU2900 25.93 32.75
PU2900/MMT-30B 31.72 55.10
PU2900/MMT-I30E 31.92
a
Melting temperature of the soft segment. ΔH: Enthalpy change.
44.79

37
38
2.4 Conclusions
PU/MMT nanocomposites were prepared by in-situ polymerization reactions. The
resulting nanocomposites exhibited increases in degree of crystallinity and mechanical
strength compared with the neat PU. These findings were attributed to the chemical
modification of MMT, especially for -CH
2
CH
2
OH groups on MMT-30B nanoclay, which
led to stronger interactions and reactivity with the matrix system. In particular, the
modifications led to strong chemical bonding between the dispersed and continuous
phases, as well as intercalation and partial exfoliation of the nanoclays, both of which
contributed to the enhancement in strength, modulus and elongation at break. Generally,
the effect of different SS MWs on the mechanical and thermal properties of PU/MMT
nanocomposites resembled the tendencies observed for neat PU. That is, as the SS MW
was increased, the tensile strength and modulus decreased for both the neat PU and the
PU/MMT nanocomposites. However, the WAXD data pattern indicated that when the SS
MW was increased in the nanocomposites (particularly PU2900/MMT-I30E), the
increase in tangled SS molecular chains altered the morphology of the MMT in PU/MMT
nanocomposites, reducing the interaction between nanoclays and the soft segments. Thus,
while the two approaches – varying SS MWs and adding modified nanoclay – can be
combined, interference can occur in some circumstances, diminishing the beneficial
effects. This finding highlights the complexity of designing nanocomposites which
combine strengthening approaches. Nevertheless, the findings indicate that suitable
modification of nanoclays and judicious selection of SS MW of components can be
combined to achieve additive effects, further expanding the design space for
nanocomposites.
39
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42
Chapter 3 1-D Nanoclay-Reinforced Polyurethane [1]

3.1 Motivation
In the chapter 2, we have found that the resulting PU/MMT nanocomposites
exhibited increases in degree of crystallinity and mechanical strength compared with the
neat PU. These findings were attributed to the chemical modification of MMT which led
to stronger interactions and reactivity with the matrix system. However, can this marked
reinforcement in PU/MMT system be only found in the presence of the 2D platelet-like
silicate? How about other layered silicates, especially for the 1D rod-like attapulgite with
a similar microstructure.
Attapulgite (ATT) is a hydrated magnesium aluminum silicate with chemical
formula Si
8
O
20
Al
2
Mg
2
(OH)
2
(OH
2
)
4
(4H
2
O) [2]. Although ATT features ribbons of a 2:1
phyllosilicate structure, it differs from other layered silicates in that it lacks continuous
octahedral sheets. Each ribbon is linked to the next by inversion of SiO
4
tetrahedra along
a set of Si-O-Si bonds which extend parallel to the X-axis, forming rectangular channels
that contain zeolitic water [3]. However, the Si-O-Si bonds are weak and are broken
easily by shear stress to form fiber crystals [3]. The average external and internal surface
areas have been estimated to be 300 and 600m
2
/g, respectively [3]. Because of the unique
morphology and structure, ATT has been used for various commercial applications, such
as adsorbents, catalysts, rheological agents and fillers [3-7].
In this chapter, 1D ATT nano-rods were used as fillers to reinforce PU elastomers.
The effect of the ATT surface treatment on composite mechanical behavior and thermal
stability was evaluated. ATT nano-rods were treated with chemical agents to generate an
43
organophilic character on the hydrophilic silicate surface to improve the compatibility
with the polyurethane matrix. The chemical modifications were confirmed by TGA and
FTIR studies. Subsequently, ATT-MDI nano-rods were incorporated into the
polyurethane matrix by in situ polymerization, and ATT-MDI/PU nanocomposites were
produced. The physical, mechanical and thermal properties of ATT-MDI/PU
nanocomposites with different ATT loadings and chemical affinities were investigated.

3.2 Experimental
3.2.1 Materials
Attapulgite (Engelhard Co.), was subjected to purification with a polymeric
dispersant, sodium polyacrylate (molecular mass ~ 4,000-5,000) [8]. After dehydration,
1,4 butanediol (BDO, Research Chemicals Ltd.) and polytetrahydrofuran (PTHF, BASF,
M.W 2900), which are the hydroxyl-terminated monomer and oligomer, respectively,
were used for the preparation of PU networks. The isocyanate (4,4’-methlenebis(phenyl
isocyanate), MDI, 98%, Sigma-Aldrich) was purified by filtration of molten MDI liquid
at 70
o
C. Dimethylformamide (DMF, Sigma-Aldrich) was used as a media solvent in the
polymerization reaction.

3.2.2 Preparation of ATT-OH and ATT-MDI
Figure 3.1 shows the modification and chemical structure of ATT nano-rods. The
polymeric dispersant was removed from ATT surfaces by heating at 500
o
C for 1h. Next,
44
ATT nano-rods were treated with 5 M hydrochloric acid at 80
o
C for 1h to produce
hydroxyl groups on the nanorod surfaces and remove residual sodium ions and impurities.
After exhaustive washing with deionized water and acetone, the ATT-OH nano-rods were
dried at 80
o
C under vacuum for 24h. Subsequently, the dried ATT-OH was ground to a
powder and dispersed in 600mL acetone by ultrasonic treatment. To graft isocyanate
molecules onto nano-rods surface, excess (8-10 g) dehydrated 4,4’-methlenebis(phenyl
isocyanate) (MDI) was added to 10 g of ATT-OH [9] and the mixture was refluxed at
80
o
C for 1h. After grafting, the ATT-MDI nano-rods were washed 3-4 times with acetone
and centrifuged at 4000rpm to remove ungrafted MDI, and finally dried at 80
o
C under
vacuum for 24h. The dried cake was ground and screened through a 325-mesh sieve to
obtain the organically modified ATT (ATT-MDI). Thermogravimetric analysis (TGA)
and Fourier transform infrared spectroscopy (FTIR) were used to determine the amount
of MDI grafted onto the surface of ATT and to identify the chemical bonds.

Figure 3.1 Modification and chemical structure of the ATT nanorods.

45
46
3.2.3 Synthesis of PU elastomer and ATT/PU composites
Polyurethane elastomer was synthesized with a molar ratio of 4:1:2.64 (MDI-to-
PTHF-to-1,4BDO) using excess diisocyanate to produce partially cross-linked networks
[10]. First, the oligomeic polyol (PTHF) was dissolved in DMF and reacted with MDI at
80
o
C for 30min to obtain a pre-polymer. Second, the chain extender (1,4 butanediol) was
added to build up the polyurethane network, and allowed to react at 80
o
C for another 2
additional min. Finally, the mixture was poured into a Teflon
®
mold and cured at 80
o
C
for 24h to obtain the PU elastomers.
Three types of ATT/PU composites were also prepared with identical molar ratios,
as shown in Table 3.1. To prepare ATT-MDI/PU nanocomposites, different amounts of
ATT-MDI nano-rods (0.5, 2.5, 5, 10 wt%) were directly blended with
polytetrahydrofuran (PTHF), which was dissolved in the DMF solvent, using a high-
speed dual-axis mixer (HM-500, Keyence, USA) for 30min to form OH-terminated
urethane oligomers. Next, excess MDI was added to the mixture, forming the pre-
polymer at 80
o
C under vacuum for 30min. Subsequently, the chain extender 1,4
butanediol was added to the pre-polymer with vigorous stirring for 2 min to complete the
reaction. Finally, the viscous polymer was poured into a Teflon
®
mold and cured at 80
o
C
for 24h to form ATT-MDI/PU nanocomposites. Figure 3.2 illustrates the steps in the
synthesis route and the associated chemical structures. In addition, acid-treated ATT
(ATT-OH) and original ATT were also incorporated into the PU polymer for comparison
with the 2.5 wt% ATT-MDI nanocomposite. These composites were prepared by the
same procedure as the former polymerization reaction.
Table 3.1 Neat PU and ATT/PU Nanocomposites.

Samples PTHF
(mole)
MDI
(mole)
BD
(mole)
Clay Content
(wt%)
Neat PU 1 4 2.64 0
0.5 wt% ATT-MDI/PU 1 4 2.64 0.5
2.5 wt% ATT-MDI/PU 1 4 2.64 2.5
5 wt% ATT-MDI/PU 1 4 2.64 5
10 wt% ATT-MDI/PU 1 4 2.64 10
2.5 wt% ATT-OH/PU 1 4 2.64 2.5
2.5 wt% ATT/PU 1 4 2.64 2.5

47

Figure 3.2 Synthesis and chemical structure of the ATT–MDI/PU nanocomposite.

48
49
3.2.4 Characterization
The ATT morphology was observed using transmission electron microscopy (TEM,
Philips 420) at 120KV. Thermogravimetric analysis (TGA 2050, TA Instruments) was
performed to determine the amount of diisocyanate grafted on the surface of ATT nano-
rods and to identify the thermal decomposition temperatures of ATT-MDI/PU
nanocomposites. In both cases, samples were heated to 800
o
C at 10
o
C/min under nitrogen
flow.
The grafting of functional groups was confirmed by FTIR spectroscopy in
transmission mode (Nicolet 4700).
The transition temperatures of soft and hard segments in ATT/PU composites were
measured by differential scanning calorimetry (DSC 2920, TA Instruments).
Measurements were performed by scanning twice to avoid signal noise and to erase
thermal history. In the first scanning, the sample was heated to 250° (at 10°/min), then
cooled to -65°C. In the second scanning, the sample was heated to 300° at a heating rate
10°/min. The second scanning was used to record transition temperatures.
The morphology of ATT-MDI/PU nanocomposite fracture surfaces was
investigated by scanning electronic microscopy (SEM, Cambridge 360, 15 kV). The
samples were gold-coated prior to SEM examination. The dispersion of the ATT-MDI
nano-rods in the PU matrix was observed using transmission electron microscopy (TEM,
Philips 420) at 120KV. Samples were sectioned using a microtome.
The storage modulus (G’) was assessed by dynamic mechanical analysis (DMA
2980, TA Instruments). Tests were conducted using the temperature scan mode, from
room temperature to 200°C, at a 10°/min heating rate. The frequency of the forced
oscillations was set to 1Hz and the applied deformation was 0.05%.
Tensile strength and modulus were measured using a universal testing machine
(Instron 8531). Tests were performed on dog-bone shaped specimens using a crosshead
speed of 10 mm/min at room temperature, in accordance with ASTM D 638-94b.

3.3 Results and Discussion
3.3.1 Characterization of ATT
The fibrous morphology of unmodified attapulgite was revealed by TEM (Figure
3.3). In the pristine condition, the fibers were rod-shaped and formed a randomly oriented,
densely packed network. The attapulgite nano-rods are approximately 15-20 nm in
diameter and several microns in length. The high aspect ratio (length-to-diameter) results
in a high specific surface area to interact with the continuous polymeric matrix. Thus,
strong interactions between ATT particles and PU are expected.
Figure 3.4 shows the thermogravimetric curves for pristine (ATT) and modified
(ATT-OH and ATT-MDI) attapulgite. The mass loss of the original ATT below 120
°
C is
attributed to evaporation of a small amount of adsorbed moisture. Discounting the mass
loss of original ATT, the mass fraction of organic groups eliminated at 600
o
C was
approximately 5.5 wt% for ATT-OH and 17 wt% for ATT-MDI. These values indicate
that the two-step surface modification process led to a significant mass fraction of
organic groups bonded to ATT. Additional insight can be gained by comparing
experimental mass fractions with those expected from a complete reaction of OH with
50
51
MDI groups. Approximately 5.5 g of OH attached to 94.5 g of the original ATT can
incorporate 80.9 g MDI, leading to a theoretical organic mass fraction of 86% at full
conversion, which is much greater than the measured value of 17%. The difference was
attributed to the consumption of NCO groups by moisture (possibly residual zeolitic
water) to form urea.
In Figure 3.5(a), the spectrum for the original ATT reveals characteristic
absorption bands at 3632 cm
-1
(from O-H stretching) and 1044 cm
-1
(for Si-O in-plane
stretching) [11]. The acid treatment of ATT (ATT-OH) and the grafting of functional
groups on ATT (ATT-MDI) were also verified by FTIR, as shown in Figures 3.5(b)-(c).
In the high wave-number region, a broad band at 3600-3200 cm
-1
was attributed to
adsorbed water molecules, and the band increased after the acid attack. (The peak at
3600-3550 cm
-1
was attributed to OH-stretching in H-bonded hydroxyls, and the peak at
3400-3200 cm
-1
was attributed to OH-stretching in free or weakly H-bonded hydroxyls.).
The band that appeared at 1625 cm
-1
in original ATT and shifted to 1600 cm
-1
with
increased intensity in ATT-OH, was caused by the bending vibration mode of adsorbed
water [12]. In addition, after the acid treatment, the absorption band at 880 cm
-1

disappeared, and the characteristic band that appeared at 960-920 cm
-1
was attributed to
the Si-O-H angle deformation vibration for acid-treated ATT [13].
The two main regions of interest in the ATT-MDI spectrum are NH absorption and
C=O stretching. The NH absorption peak at 3320 cm
-1
, can be attributed to hydrogen-
bonded NH groups of urethane linkages [14]. The small shoulder, seen at 3420 cm
-1
, is
characteristic of stretching of free NH groups. Moreover, the broad band arising from
adsorbed water in the high wave-number region is absent in this spectrum, indicating that
52
the hydrophilic character of the ATT-MDI strongly decreased. On the other hand, peaks
in the 1735–1690 cm
-1
region, expected from free and hydrogen-bonded urethane
carbonyl are almost imperceptible at this scale. However, the absorbance due to
stretching of NCO groups at 2270 cm
-1
[14] is clearly discerned. The amide bands at
1645, 1511 and 1309cm
-1
, assigned to CO (Amide I), NH (Amide II) and CNH (Amide
III), respectively, which are consistent with polyurea formation [14-16], are also clearly
observed. These results indicate that part of the MDI molecules were grafted onto the
ATT surface through chemical bonding (according to Figure 3.1), although the remainder
may be consumed in the formation of urea, which should be strongly adsorbed onto the
ATT surface.

Figure 3.3 TEM image of the original ATT.

53

Figure 3.4 TGA curves for the ATT nanorods before and after modification.

54

Figure 3.5 FTIR spectra for the (a) original ATT, (b) ATT–OH, and (c) ATT–MDI.

55
3.3.2 ATT/PU Composites
3.3.2.1 DSC measurements
The DSC results from neat PU and from ATT/PU composites are summarized in
Table 3.2. The data show evidence of profound differences in structure resulting from the
addition of ATT to the PU networks. Because the structure of PU includes hard and soft
segments, two melting temperatures are expected. However, the DSC curve revealed only
one peak, and this was associated with the melting temperature of soft segment (T
m
, SS).
The peak associated with the melting temperature of the hard segment (T
m
, HS) was
absent from the DSC curve [17-19]. The absence was attributed to the inactive movement
of hard segment, which has a small Cp Δ [20], and to the widely dispersed HS micro-
domains within the PU matrix [21].
The melting points in the region between 21.6 and 27.9
o
C are characteristic of
ordered soft segment structures and indicate that polytetrahydrofuran crystallizes in both
neat PU and ATT/PU composites. For these soft segments (SS), the enthalpy change of
ATT-MDI/PU nanocomposites increases from 43.2 J/g to 46.5 J/g with increasing ATT-
MDI content. This increase in soft segment melting temperature and enthalpy change can
be attributed to the effects produced by the rigid nanofillers in the soft segment domains,
which act as nucleation seeds, increasing the degree of crystallinity of the PU networks
[22]. Moreover, these values increase with ATT-MDI concentrations up to 5%, then
decrease for the 10% sample. This behavior is attributed to the reduction in effective
surface area caused by particle agglomeration at high ATT-MDI contents. Similarly, for
both ATT-OH/PU and ATT/PU composites, the soft segment melting temperature and
enthalpy change increased relative to neat PU, indicating that the crystalline nano-rods of
56
ATT-OH and original ATT also acted as seeds to nucleate and enhance the crystallinity
of the PU matrix. On the other hand, when the ATT concentration was held constant at
2.5%, the ΔH values differed with ATT treatment, indicating that the compatibility of the
filler with the matrix also played a significant role in the PU crystallization process.

Table 3.2 Thermal Properties of Neat PU and the ATT/PU Composites

Samples T
m,SS
(
o
C)
a
ΔH (J/g)
Neat PU 21.6 37.8
0.5 wt% ATT-MDI/PU 25.5 43.2
2.5 wt% ATT-MDI/PU 25.3 44.3
5 wt% ATT-MDI/PU 27.9 46.5
10 wt% ATT-MDI/PU 25.0 38.5
2.5 wt% ATT-OH/PU 23.2 43.6
a
Melting temperature of the soft segment. ΔH: Enthalpy change.
2.5 wt% ATT/PU 25.5 40.2

57
58
3.3.2.2 Morphology of PU elastomer and ATT-MDI/PU nanocomposites
The morphology of ATT-MDI/PU nanocomposites was investigated by TEM and
SEM, respectively. TEM observation of ATT-MDI/PU nanocomposite with 2.5wt%nano-
rods indicated the modified ATT nano-rods were dispersed quite uniformly in the PU
matrix as single crystals or crystal bundles with diameters smaller than 100nm (Fig.
3.6(a)). The relatively good dispersion of nano-rods in the PU matrix can be the result of
a stable hindrance developed between the modified nano-rods caused by the grafted MDI
chains on the ATT surface as well as the filler increased compatibility with the polymer.
Both facts contributed to the nano-rods dispersion and prevented agglomeration. The
uniform dispersion illustrates the effectiveness of the two-step modification of ATT for
the PU system. However, when the loading of ATT-MDI nano-rods was increased to
10wt%, the dispersion became less uniform and agglomeration was observed (Fig 3.6(b)).
The agglomeration reduced the effective surface area of the ATT-MDI and diminished
the interfacial adhesion.
Tensile fracture surfaces of the PU elastomer and the ATT-MDI/PU
nanocomposites were examined by SEM. Figure 3.7 shows typical SEM images of neat
PU and ATT-MDI/PU nanocomposites with 5wt% nano-rods. The fracture surface of
neat PU (Figure 3.7(a)) was uneven but relatively smooth (at 10,000X). In contrast, the
fracture surface of the 5wt% ATT-MDI nanocomposite exhibited pronounced micro-
roughness (Figure 3.7(b)). The micro-rough morphology evidenced in Figure 3.7(b) was
attributed to the inherent anisotropy and rigidity of the ATT-MDI, which caused micro-
scale crack deflection. While the nano-rods were well-dispersed in the PU matrix, the
interfacial adhesion with the polymeric matrix was enhanced, and the mobility of
polymeric molecules in close proximity was restricted [23]. The uneven and rough
surface morphology observed for the 5wt% of ATT-MDI/PU nanocomposite derived
from the presence of ATT and the associated constraints on chain mobility in the PU
matrix.

Figure 3.6 TEM images of ATT–MDI/PU nanocomposite with (a) 2.5 and (b) 10 wt %
nanorods.

59

Figure 3.7 SEM images of (a) neat PU (b) ATT-MDI/PU nanocomposite with 5%wt
nano-rods.

60
61
3.3.2.3 Dynamic Mechanical Analysis (DMA)
Table 3.3 lists the values of dynamic storage modulus (G’) at 40° and 100°C as a
function of weight% of ATT-MDI, as well as the transition temperature of hard segment
(T
g
, HS). The storage modulus values at both temperatures (40°C and 100°C) increased
with increased loadings of ATT-MDI nano-rods. An increase in storage modulus of
approximately 20% over neat PU was observed for 10 wt% ATT-MDI nanocomposites at
40°C. However, at 100°C, the values of storage modulus for all samples dropped by
roughly half, underscoring the heat sensitivity of both neat PU and ATT-MDI/PU
nanocomposites within this temperature range. Despite the decrease in storage modulus
values at 100°C, the G’ values for ATT-MDI/PU nanocomposites still exceeded those of
neat PU for most samples. In addition, the storage modulus values for ATT-OH/PU and
ATT/PU composites at 40° and 100°C were similar to neat PU, but less than ATT-
MDI/PU nanocomposites with the same filler contents. This behavior arises from the
chemical modification of ATT by grafted MDI molecules and the enhanced interfacial
adhesion to the PU matrices.
The transition temperature of hard segment (T
g
, HS) can be observed from the
peaks in the tan δ vs. temperature curve. An increase in transition temperature of hard
segment (T
g
, HS) with increasing ATT-MDI loading is evident in the data summarized in
Table 3.3. This increase of T
g
can be ascribed, at least in part, to the restricted mobility of
amorphous polymer chains that results from the physical cross-links induced by
crystallization, as noted by other authors [24-25]. However, the shift in T
g
with increasing
additions of ATT-MDI nano-rods is so large that it can only be explained as the direct
62
consequence of the decreased flexibility of polymer chains in the presence of stiff
crystalline ATT.
Note that both the ATT-OH/PU and ATT/PU nano-composites have hard segment
transition temperatures similar to neat PU. This observation suggests that ATT-OH and
original ATT do not substantially restrict the mobility of polymer chains, despite the high
filler surface area of these crystalline nano-fillers. This is attributed to weak interfacial
adhesion in these materials. From the DSC and DMA results, we conclude that the ATT-
MDI nano-rods, in contrast, induce additional crystallinity within the soft segment, but
restrict the movement of polymer chains in the hard segment. This behavior is caused by
the heterogeneous nano-rods, the enhanced interfacial adhesion to the PU matrix, and co-
reaction between filler and polymer.

Table 3.3 Storage modulus values (G’) at 40°C and 100°C and the peak in tan δ vs.
temperature curve for hard segment.

a
Glass-transition temperature of the hard segment.

Samples
G' (MPa) at 40
o
C G' (MPa) at 100
o
C T
g
, HS (
o
C)
a
Neat PU 21.96 10.42 162.6
0.5 wt% ATT-MDI/PU 22.83 9.91 162.8
2.5 wt% ATT-MDI/PU
24.24 12.00 173.8
5 wt% ATT-MDI/PU 24.90 12.48 174.6
10 wt% ATT-MDI/PU 26.18 14.93 189.8
2.5 wt% ATT-OH/PU
23.25 10.82 163.7
2.5 wt% ATT/PU 22.85 10.53 161.3
63
64
3.3.2.4 Tensile properties
Yield strength and elastic modulus values for different ATT-MDI nanocomposites
are plotted as a function of nano-rod filler content in Figure 3.8. In comparison with neat
PU, the tensile properties of the ATT-MDI/PU nanocomposites increased significantly.
Figures 3.8 shows that the Young’s modulus and tensile yield strength increased 36% and
12% respectively after addition of 2.5wt% ATT-MDI nano-rods. The 10wt% ATT-
MDI/PU nanocomposite showed increases in Young’s modulus and tensile strength of
more than 75% relative to neat PU. The increased modulus arises from the reinforcing
effect provided by the ATT-MDI nano-rods.
Note that the melting point of soft segment (T
m
, SS) is close to room temperature.
Thus, the tensile properties at room temperature are directly related to the soft segment
characteristics. Korley et al reported that the presence of crystallites within the soft
segment might absorb strain energy upon deformation. Thus, semicrystalline soft
segments can act as an effective load-bearing phase during deformation [26]. As shown
by the DSC measurements, the ATT-MDI rods act as nucleation seeds, increasing the
degree of crystallinity of the soft segment. In addition, the rigid crystalline ATT-MDI
nano-rods also contribute to the increased hardness by virtue of simple reinforcement.
Nanocomposite specimens containing hydroxyl functional groups on the ATT
surface (ATT-OH) and composites prepared from the as-received ATT were also
mechanically tested. Unlike the ATT-MDI nano-rods, the addition of ATT-OH and ATT
nano-rods to the polymeric matrix caused decreases in Young’s modulus and yield
strength (Figure 3.9). As summarized in Table 3.4, the ATT-MDI/PU nanocomposite
exhibits a 40% increase in Young’s modulus and a 15 to 30% increase in yield strength
compared to the other composites with similar loading. These results illustrate that the
adequate dispersion of the rigid filler into the polymer is necessary but not sufficient to
enhance composite mechanical properties, and that surface or chemical modifications that
lead to an improved interfacial adhesion between filler and matrix are also necessary.

Figure 3.8 Yield stress and Young’s modulus for ATT-MDI/PU nanocomposites with
different amounts of modified ATT.

65

Figure 3.9 Yield stress and Young’s modulus for different functional fillers of ATT/PU
composites.

66
Table 3.4 Young’s modulus and yield stress for ATT/PU composites with different
functional fillers.

Samples Young’s Modulus (MPa) Yield Strength (MPa)
Neat PU 6.69 1.22
2.5 wt% ATT-MDI/PU 9.12 1.37
2.5 wt% ATT-OH/PU 6.51 1.05
2.5 wt% ATT/PU 6.46

1.19

67
68
3.3.2.5 Thermal degradation
Conventional PU is regarded as having poor thermal stability. However, the
addition of ATT-MDI nano-rods is expected to enhance thermal stability. Decomposition
temperature was measured by thermogravimetric analysis (TGA) for the ATT-MDI/PU
nanocomposites and the neat PU, and the results are shown in Figure 3.10. The profiles of
weight loss for PU and ATT-MDI/PU nanocomposites can be divided into two stages
(ignoring the small weight loss below 200°C, resulting from adsorbed moisture
evaporation). In the first stage (280° ～370°C), a 25% weight loss occurs, and this is
associated with decomposition of hard segment. The second stage, starting at 370°C for
PU and at 380°C for ATT-MDI/PU nanocomposites, involves major weight loss
stemming from soft segment decomposition, which is consistent with the work of
Petrovic et al [27]. Both stages exhibit a difference of at least 10°C in decomposition
temperature when comparing the ATT-MDI/PU nanocomposites to neat PU. This
difference arises because the ATT-MDI nano-rods act as a physical barrier in the
polymeric network and limit the gas diffusion at the interface, thus affecting the kinetics
of the degradation reactions [28].

Figure 3.10 Thermal degradation behaviors of PU and ATT-MDI/PU nanocomposites.

69
70
3.4 Conclusions
Attapulgite/polyurethane nanocomposites were synthesized by copolymerization
reaction after a two-step functionalization of ATT. The resulting nanocomposites
exhibited increases in degree of crystallinity, thermal stability and mechanical strength
compared with the neat PU. The findings are attributed to the chemical modification of
ATT, which led to suitable surface compatibility and reactivity with the matrix system.
Relatively uniform dispersion and strong interfacial adhesion were created by chemical
bonding between dispersed and continuous phases. The chemical modification of the
attapulgite also facilitated the development of the nanocomposite network. Further
improvements and applications in specific polymer matrices are likely to be possible with
optimized surface treatments. The addition of small amounts of nano-rods will afford
opportunities to design polymers with enhanced properties for adhesives and composite
materials.

71
Chapter 3 References
1. Wang Chia-Hao., Auad María L., Marcovich Norma E., Nutt Steven., J Appl Polym
Sci 2008, 109, 2562.

2. Bradley Amer., Mineral 1940, 25, 405.
3. Galan E., Clay Miner 1996, 31,443-453.
4. Phillips Ken A.. Arizona Department of Mines and Mineral Resources 1989,
September, Report 89-3.

5. Haydn Murray., MMSD 2002, March, No.64.
6. Li An., Wang Aiqin., Eur Polym J 2005, 41, 1630..
7. Shen Liang., Lin Yijian., Du Qiangguo., Zhong Wei., Yang Yuliang., Polymer 2005
46, 5758.

8. Purcell Jr R.J., Parker D.C., US patent 2002, 6, 444, 601.
9. Guo Zhao-xia., Liu Wen-Fang., Li Ying., Yu Jian., J. Macromolecular Science Part
A-Pure and Applied Chemistry 2005, 42, 221.

10. Prisacariu C., Olley R.H., Caraculacu A.A., Bassett D.C., Martin C., Polymer 2003,
44, 5407.

11. Chen-Yang Y.W., Lee Y.K., Chen Y.T., Wu J.C., Polymer 2007, 48, 2969.
12. Myriam M., Suarez M., Martin Pozas J.M., Clays and Clay Minerals 1998, 46, 3, 225.
13. Fripiat J.J., Leonard A., Barake N., Bull Soc Chim 1963, 1, 122.
14. Pattanayak Asim., Jana Sadhan C., Polymer 2005, 46, 5183.
15. Kuptsov A.H., Zhizhin G.N., Handbook of fourier transform Raman and spectra of
polymers, Elsevier, New York, 1988.

16. Chattopadhyay D.K., Raju K.V.S.N., Polym Sci 2007, 32, 352.
17. Speckhard T.A., Hwang K.K.S., Yang C.Z., Laupan W.R., Cooper S.L., J Macromol
Sci Phys B 1984, 23,175.

18. MacKnight W.J., Yang M., Kajiyama T., Polym Prepr (Am Chem Soc, Div Polym
Chem) 1968, 9, 860.
72
19. Yang C.Z., Hwang K.K.S., Cooper S.L., Makromol Chem 1983,184, 651.
20. Yilgor I., Riffle J.S., Wilkes G.L., McGrath J.E., Polym Bull (Berlin) 1982, 8535.
21. Li Y., Tong G., Liu J., Linliu K., Desper C.R., Chu B., Macromolecules 1992, 25
7365.

22. Sun Li., Yang Jin-Tao., Lin Gen-Yao., Zhong Ming-Qiang., Materials Letters 2007,
61, 3963.

23. Etan A., Fisher F.T., Andrews R., Brinson L.C., Schadler L.S., Composites Science
and Technology 2006, 66, 1159.

24. Mathew A.P., Dufresne A., Biomacromolecules 2002, 3, 609.
25. Auad María L., Contos Vasili S., Nutt Steven., Aranguren Mirta I., Marcovich Norma
E., Polymer International 2008, 57, 651.

26. James Korley., LaShanda T., Pate Brian D., Thomas Edwin L., Hammond Paula T.,
Polymer 2006, 47, 9, 3073.

27. Petrovic Z.S., Zavargo Z., Flynn J.H., Macknight W.J., J Appl Polym Sci 1994, 51,
1087.

28. Golebiewski Jan., Galeski Andrzej., Composites Science and Technology 2007, 67,
3442.

73
Chapter 4 Effects of Organophilic Modified Attapulgite on Polyurethane [1]

4.1 Motivation
In the chapter 3, the polyurethane/Attapulgite nanocomposites were synthesized
by copolymerization reaction after a two-step functionalization of ATT. The resulting
nanocomposites also exhibited increases in degree of crystallinity, thermal stability and
mechanical strength compared with the neat PU. The findings are attributed to the
chemical modification of ATT, which led to suitable surface compatibility and reactivity
with the matrix system. However, the measured amount of organic molecules grafted
onto ATT was much lower than the predictions of theoretical calculation in the chapter 3
and this may affect the efficiency of the reinforcement.
Researchers have studied the heat treatment [2-7] and acid treatment of ATT [8-12].
These treatments are commonly used to increase surface area and enhance active sites for
ATT, although the combination of heat and acid treatment has rarely been considered.
Additional surface modification is accomplished through coupling agents, which only are
grafted onto the inorganic silicate surface and reacts with polymer chains. The idea to use
diisocyanate molecule as a coupling agent on silicates surface via reaction with silanol
groups (Si-OH) was first introduced by Yosomiya et al [13], who grafted various
diisocyanate compounds via silanol groups onto the glass fiber. Their work indicated that
the silanol group (Si-OH) had reactivity similar to alcohol (OH) when reacted with
diisocyanate, and attributed this to a comparable activation energy.
In this chapter, we use heat and acid treatments to form Si-OH groups on ATT
followed by grafting MDI to investigate the effects of different grafting contents on
74
thermal and mechanical properties of the in-situ synthesized PU/ATT-MDI
nanocomposites.

4.2 Experimental
4.2.1 Materials
Attapulgite (Attagel 50, Engelhard Co., USA) was selected for experiments and
was used as-received. The diisocyanate, MDI (98%, Sigma-Aldrich), was purified by
filtration of molten MDI liquid at 70°C. Dibutyltin dilaurate (95%, Sigma-Aldrich
Chemical Co) was used as a catalyst for grafting MDI molecules on the surface of
attapulgite. PU and PU/ATT-MDI nanocomposites were prepared using 1,4 butanediol
(BDO, Avocado Research Chemicals Ltd.) and polytetrahydrofuran (PTHF, Sigma-
Aldrich, M.W 2900), which are the hydroxyl-terminated monomer and oligomer,
respectively. The 1,4 butanediol was dried over calcium hydride for 48h and then was
vacuum distilled. Polytetrahydrofuran was dehydrated in a vacuum oven at 60°C for 48h.
Dimethylformamide (DMF, Sigma-Aldrich) was dehydrated and used as a solvent in the
polymerization reaction.

4.2.2 DOE, heat treatment, acid treatment, and grafting of ATT
The design of experiment was guided by software (S-Matrix CARD DOE) and was
based on three parameters - temperature of heat treatment (500, 600, 700, and 800
°
C),
concentration of hydrochloric acid (1, 3, 5, 7, and 9M), and reaction time for acid
treatment (1, 2, 4, 6, and 8h). A total of 17 experimental preparations were specified by
75
the DOE. The 17 heat- and acid-treated ATT samples were prepared by the following
steps. (1) 20 grams of ATT powder were heated to 500, 600, 700, and 800°C for 1h, then
cooled to room temperature in a desiccator. The heat-treated samples were designated
H500, H600, H700, and H800, respectively. (2) Heat-treated ATT samples were soaked
in hydrochloric acid (1, 3, 5, or 9M) at 80°C for 1, 2, 4, 6, or 8h to produce hydroxyl
groups on the nano-rod surfaces (ATT-OH) and to remove residual sodium ions and
impurities. After exhaustive washing with deionized water and acetone, the ATT-OH
nano-rods were dried at 80°C under vacuum for 24h. After heat and acid treatments,
sample designations were modified. For example, H500 became H500A5M1H, for heat
treatment at 500
o
C followed by acid treatment in 5M HCl solution for 1h. (3) The dried
ATT-OH cake was ground to a powder (10 g) and dispersed in a 600ml acetone bath
under ultrasonic agitation. To graft isocyanate molecules onto nano-rod surfaces, 8-10 g
of dehydrated 4,4’-methlenebis(phenyl isocyanate) (MDI) was added to 10 g of ATT-OH
and 2 to 3 drops of dibutyltin dilaurate were added as a catalyst. The mixture was then
refluxed at 80°C for 1h. After grafting, the nano-rods were washed 3-4 times with
acetone and centrifuged at 4000rpm to remove ungrafted MDI, and finally dried at 80
°
C
under vacuum for 24h. (4) The dried cake was ground and screened through a 325-mesh
sieve to obtain the organically modified ATT (ATT-MDI).

4.2.3 Synthesis of PU elastomer and PU/ATT-MDI nanocomposites
PU elastomer was synthesized with a molar ratio of 4:1:2.64 (MDI-to-PTHF-to-
1,4BDO), using excess diisocyanate to produce partially cross-linked networks and –
NCO end groups [14]. First, the oligomeric polyol (PTHF) was dissolved in DMF and
76
reacted with MDI at 80°C for 30min to obtain a pre-polymer in a round bottom flask with
continuous stirring under vacuum. Second, the chain extender (1,4 butanediol) was added
to build up the polyurethane network, and the mixture was allowed to react at 80°C for an
additional 2 min. In initial experiments, the network building reaction was active within 2
min (after the addition of 1,4 BDO). If the reaction was allowed to continue beyond 2min,
the viscosity increased rapidly, making it impossible to pour out of the round bottom
flask. Finally, the mixture was poured into a Teflon
®
mold and cured at 80°C for 24h to
obtain the PU elastomers. To prevent moisture absorption, the reaction was performed
under vacuum, and all chemicals were dehydrated beforehand.
Three different grafting amounts of ATT-MDI nano-rods (ATT-MDI-
H500A1M1H, ATT-MDI-H600A5M4H, and ATT-MDI-H700A3M2H), were chosen to
synthesize the PU/ATT-MDI nanocomposites, representing high, medium, and low
grafting contents, respectively. The PU/ATT-MDI nanocomposites were prepared with
2.5 wt% ATT-MDI loading in identical molar ratios. Figure 4.1 illustrates the steps in the
synthesis route and the associated chemical structures. First, the ATT-MDI nano-rods
were directly blended with polytetrahydrofuran (PTHF), which was dissolved in the DMF
solvent using a high-speed dual-axis mixer (HM-500, Keyence, USA) for 30min to form
OH-terminated urethane oligomers. Second, excess MDI was added to the mixture,
forming the pre-polymer at 80°C under vacuum for 30min. Subsequently, the chain
extender 1,4 butanediol was added to the pre-polymer with vigorous stirring for 2 min to
complete the reaction. Finally, the viscous polymer was poured into a Teflon
®
mold and
cured at 80°C for 24h to form PU/ATT-MDI nanocomposites. The three samples will be
referred to as PU/ATT-MDI-H500A1M1H, PU/ATT-MDI-H600A5M4H, and PU/ATT-
MDI-H700A3M2H, respectively.

Figure 4.1 Synthesis of PU/ATT-MDI nanocomposite.
77
78
4.2.4 Characterization
Thermogravimetric analysis (TGA 2050, TA Instruments) was performed to
determine the amount of diisocyanate grafted on the surface of ATT nano-rods. Samples
were heated to 800°C at 10
o
C/min under nitrogen flow.
FTIR spectroscopy (Nicolet 4700) with a sensitivity of 128 scans and a resolution
of 4 cm
-1
was performed on the untreated ATT, heat- and acid-treated ATT (ATT-OH),
and grafted ATT (ATT-MDI), respectively. KBr/ATT disks were prepared and used to
obtain the IR spectra.
The morphology of untreated ATT and the heat and acid-treated ATT nano-rods
were observed by transmission electron microscopy (TEM Philips EM420). The TEM
samples were prepared using gentle sonication in deionized water to create a colloidal
suspension. A drop of the colloidal was placed on the grid with dry-analyzed at an
acceleration voltage of 120kV. The dispersion of the ATT-MDI nano-rods in the PU
matrix was also observed with ultra-thin samples which were sectioned by cryogenic
ultramicrotoming.
The melting temperatures of the soft and hard segments in polyurethane and
PU/ATT-MDI nanocomposites were measured by differential scanning calorimetry (DSC
2920, TA Instruments). Measurements were performed by scanning twice to avoid signal
noise and to erase thermal history. In the first scan, the sample was heated to 250° (at
10°/min), then cooled to -65°C. In the second scan, the sample was heated to 300°C at a
heating rate of 10°/min. The second scan was used to record transition temperatures.
The storage modulus (G’) was assessed by dynamic mechanical thermal analysis
(DMTA Q800, TA Instruments). Tests were conducted using the temperature scan mode,
79
from -100 to 250°C, at a 10°/min heating rate. The frequency of the forced oscillations
was set to 1Hz and the applied deformation was 0.05%.
Tensile strength, modulus, and elongation at break were measured using a
universal testing machine (Instron 8531). Tests were performed on dog-bone shaped
specimens using a crosshead speed of 10 mm/min at room temperature, in accordance
with ASTM D 638-94b.

4.3 Results and Discussion
4.3.1 Characterization of ATT
The fibrous morphology of untreated and treated ATT were revealed by TEM
(Figure 4.2). In the untreated condition (Fig 4.2(a)), the fibers were rod-shaped and
formed a randomly oriented, densely packed network. The ATT nano-rods were
approximately 15-25 nm in diameter and 0.2 to several microns in length. The high
aspect ratio (length-to-diameter) resulted in high specific surface area to interact with the
continuous polymeric matrix. After heat and acid treatment, the external morphology of
ATT was preserved (Fig 4.2(b)), although rough surface with pits was evident in the ATT
nano-rods. The persistence of the fibrous morphology indirectly shows that the heat and
acid treatment affected structure throughout the internal microchannels.
Figure 4.3 shows two thermogravimetric curves - one for pristine ATT after heat
(500
o
C) and acid (1M HCl for 1 h) treatment (ATT-OH), and one for MDI-grafted ATT
(ATT-MDI). The mass loss of the ATT-OH below 120°C is attributed to evaporation of a
small amount of adsorbed moisture. Below 120°C, the MDI-grafted ATT curve shows
much less weight loss than the heat- and acid-treated ATT. The ATT-MDI had become
80
hydrophobic clay and consequently adsorbed less moisture than ATT-OH. Discounting
the mass loss of ATT-OH, the mass fraction of organic groups eliminated at 600°C was
approximately 52 wt% for ATT-MDI. The mass fraction lost indicates that the two-step
surface modification process led to a significant mass fraction of organic groups bonded
to ATT. After the TGA measurement, the content of the grafted diisocyanate on the
surface of ATT nano-rods based on DOE was measured, as summarized in Table 4.1.
H500A1M1H exhibited the maximum amount of MDI grafting, while H700A3M2H
exhibited the minimum amount of MDI grafting. The difference between two samples in
the amount of MDI grafting was more than 20 wt%, and was attributed to changes in the
structure and composition in ATT after the heat and acid treatment.
The FTIR spectrum of the untreated ATT, H700A3M2H ATT-OH, and
H700A3M2H ATT-MDI (which had the minimum amount of MDI grafting) are shown in
Fig 4.4. In the high wave-number region (3700-3200 cm
-1
), the untreated ATT (Fig 4.4(a))
revealed two characteristic absorption bands at 3615 and 3548 cm
-1
, and two shoulders at
3400-3420 and 3260-3290 cm
-1
. The first two bands were assigned to Al
3+
-OH and (Fe
3+
,
Mg)-OH, while the two shoulders were attributed to hydrogen bound to oxygen in Si-O-
Si and Si-O-Al groups, respectively [15]. After heat and acid treatment, the high wave-
number region became broad and the band that appeared at 1655 cm
-1
in untreated ATT
shifted to 1636 cm
-1
with decreased intensity. Mendelovici suggested that this shift was
related to energetically different hydrogen bondings [8]. In addition, after the heat and
acid treatment, the absorption band at 878 cm
-1
disappeared, and a new shoulder appeared
at 960-955 cm
-1
. Also, the three absorptions at 985, 1031, and 1195 cm
-1
were merged
into one characteristic band at 1072 cm
-1
. The new shoulder signified that new silanol
81
groups (Si-OH) had formed and the merged band indicated that the ATT crystal structure
had been modified.
Two of the main regions of interest in the spectrum of H700A3M2H ATT-MDI are
the NH absorption and C=O stretching. The NH absorption peak at 3316 cm
-1
, can be
attributed to hydrogen-bonded NH groups of urethane linkages [16], while the small
shoulder at 3420 cm
-1
is characteristic of stretching of free NH groups. Moreover, the
broad band arising from adsorbed water in the high wave-number region is absent (Fig
4.4(c)), indicating that the hydrophilic character of the ATT-MDI strongly decreased. On
the other hand, the absorbance due to stretching of NCO groups at 2275 cm
-1
[16] is
clearly visible. The amide bands at 1645, 1511 and 1309cm
-1
, assigned to CO (Amide I),
NH (Amide II) and CNH (Amide III), respectively, are also clearly observed, and are
consistent with polyurea formation [16-18]. These results indicate that the MDI
molecules were grafted onto the ATT surface through chemical bonding even in the
H700A3M2H ATT-MDI, which had the lowest grafting amount.

(a)

(b)

Figure 4.2 TEM images for (a) untreated ATT and (b) heat and acid-treated ATT.

82

Figure 4.3 TGA curves for the ATT nano-rods before and after grafting.

83

Table 4.1 Grafting amount of MDI as obtained from TGA heating at 20
o
C/min.

ATT-OH
samples
Weight loss
(wt%)
ATT-MDI
samples
Weight loss
(wt%)
Grafting amount of
MDI (wt%)
H500A1M1H 88.99 H500A1M1H 36.98 52.01
H500A1M1H 89.02 H500A1M1H 37.16 51.86
H500A5M4H 89.30 H500A5M4H 43.36 45.94
H500A5M8H 91.41 H500A5M8H 52.08 39.33
H500A9M1H 91.45 H500A9M1H 52.34 39.11
H500A9M8H 91.84 H500A9M8H 54.57 37.27
H600A3M6H 91.08 H600A3M6H 48.44 42.64
H600A5M4H 92.13 H600A5M4H 46.61 45.52
H600A5M4H 91.57 H600A5M4H 46.31 45.26
H600A9M4H 92.14 H600A9M4H 49.40 42.74
H700A3M2H 90.62 H700A3M2H 59.91 30.71
H800A1M4H 93.88 H800A1M4H 42.36 51.52
H800A1M8H 92.79 H800A1M8H 47.94 44.85
H800A5M1H 92.90 H800A5M1H 43.39 49.51
H800A5M8H 92.65 H800A5M8H 44.61 48.04
H800A9M4H 91.67 H800A9M4H 43.13 48.54
H800A9M4H 91.53 H800A9M4H 43.36 48.17

84

Figure 4.4 FTIR curves for (a) ATT (b) H700A3M2H ATT-OH (c) H700A3M2H ATT-
MDI.

85
86
4.3.2 PU/ATT-MDI nanocomposites
4.3.2.1 Morphology of PU/ATT-MDI nanocomposites
The morphology of PU/ATT-MDI-H500A1M1H, PU/ATT-MDI-H600A5M4H,
and PU/ATT-MDI-H700A3M2H nanocomposites which were synthesized from ATT-
MDI of high, medium, and low grafting amounts, respectively, were inspected by TEM.
TEM images (Figure 4.5) of PU/ATT-MDI nanocomposites with 2.5wt% ATT-MDI
nano-rods indicated that the MDI-modified ATT nano-rods were dispersed in the PU
matrix as bundles with diameters smaller than 100nm in PU/ATT-MDI-H500A1M1H
and PU/ATT-MDI-H600A5M4H nanocomposites (Fig 4.5(a)-(b)). However, the ATT
nano-rods showed relatively poor dispersion in the PU/ATT-MDI-H700A3M2H
nanocomposite (Fig 4.5(c)). The superior dispersion of ATT nano-rods in the PU matrix
can be attributed to compatibility between the MDI-modified ATT and PU matrix. This
suggests that the H700A3M2H ATT-MDI contained insufficiently modified nano-rods
which have affinity to cause agglomeration. The agglomeration reduced the effective
surface area of the ATT-MDI and diminished the interfacial adhesion in PU/ATT-MDI-
H700A3M2H nanocomposite.

(a)

(b)

(c)

Figure 4.5 TEM images of (a) PU/ATT-MDI-H500A1M1H (b) PU/ATT-MDI-
H600A5M4H and (c) PU/ATT-MDI-H700A3M2H nanocomposites.

87
4.3.2.2 DSC measurements
The DSC results for neat PU and PU/ATT-MDI nanocomposites are summarized in
Table 4.2. Because the structure of PU includes hard and soft segments, two melting
temperatures are expected. However, the DSC curve (not shown) exhibited only one peak
and this was associated with the melting of the soft segment (T
m
, SS). The peak
associated with the melting temperature of the hard segment (T
m
, HS) was absent from
the DSC curve [19-21]. The absence was attributed to the inactive movement of hard
segment, which has a small Cp Δ [22], and to the widely dispersed HS micro-domains
within the PU matrix [23].
The melting temperatures in the region between 21.6 and 25.84°C are characteristic
of ordered soft segment structures. The values also indicate that polytetrahydrofuran
crystallized in both neat PU and PU/ATT-MDI nanocomposites. For these soft segments
(SS), the melting enthalpy increases from 37.8 JU/g for PU to ~ 44 J/g for all three
PU/ATT-MDI nanocomposites. The increases in melting temperature and enthalpy of the
soft segment are attributed to the effects produced by the rigid nano-rods in the soft
segment domains. The nano-rods act as nucleation seeds, increasing the degree of
crystallinity of the PU networks. Nunes et al. [24] observed that soft segments interacted
preferentially with silica, resulting in reduced interactions between soft and hard phases.
In other words, the silica increased phase segregation and facilitated orientation and
crystallization. In addition, the melting temperature and enthalpy of the soft segment
were not significantly altered in PU/ATT-MDI-H500A1M1H, PU/ATT-MDI-
H600A5M4H and PU/ATT-MDI-H700A3M2H nanocomposites. This suggests that the
88
grafting content had negligible effect on the thermal behavior, although the modified
ATT-MDI nano-rods enhanced the crystallinity of soft segments in the PU system.

Table 4.2 Melting temperature and enthalpy of soft segment in PU and PU/ATT-MDI
nanocomposites.
Samples T
m
,
SS
(°C) ΔH (J/g)
Neat PU 21.6 37.8
PU/ATT-MDI-H500A1M1H 25.84 44.98
PU/ATT-MDI-H600A5M4H 25.63 44.71
PU/ATT-MDI-H700A3M2H 25.49

44.44

89
90
4.3.2.3 Dynamic Mechanical Analysis (DMA)
The values of dynamic storage modulus (G’) at 50° and 100°C and, the glass
transition temperature of soft segment (T
g
, SS) and hard segment (T
g
, HS) are
summarized in Table 4.3. The storage modulus values at both temperatures (50°C and
100°C) increased with the addition of 2.5 wt% of ATT-MDI nano-rods. An increase in
storage modulus of approximately 17% (over neat PU) was observed for PU/ATT-MDI-
H500A1M1H nanocomposites at 50°C. However, at 100°C, the values of storage
modulus for all samples dropped by roughly half, underscoring the heat sensitivity of
both neat PU and PU/ATT-MDI nanocomposites within this temperature range. Despite
the decrease in storage modulus values at 100°C, the G’ values for PU/ATT-MDI
nanocomposites still exceeded those of neat PU. In addition, the storage modulus values
for PU/ATT-MDI-H500A1M1H, PU/ATT-MDI-H600A5M4H, and PU/ATT-MDI-
H700A3M2H nanocomposites at both 50°C and 100°C were similar, suggesting that the
dynamic mechanical thermal behavior was not strongly dependent on the grafting amount
of ATT-MDI nano-rods above the melting temperature of soft segment (Table 4.2). The
DMA data was consistent with the DSC data.
The glass transition temperature of soft segment (T
g
, SS) and hard segment (T
g
, HS)
can be observed from the peaks in the tan δ vs. temperature curve. The glass transition
temperature values for soft segment (T
g
, SS) was unchanged relative to the neat PU and
PU/ATT-MDI nanocomposites. The glass transition temperature of soft segment was not
influenced by the addition of nano-rods or the different grafting amount of ATT-MDI. An
increase in the T
g
of hard segment (T
g
, HS) with 2.5 wt% loading of ATT-MDI is evident
in Table 4.3. This increase of T
g
is ascribed, at least in part, to the restricted mobility of
amorphous polymer chains that results from the physical cross-links induced by
crystallization, as noted by other authors [25-26]. However, there is no obvious
dependence of T
g
, HS on grafting amount of ATT-MDI.
From the DSC and DMA results, we conclude that the ATT-MDI nano-rods
enhance crystallization in the soft segment but reduce the mobility of polymer chain in
the hard segment. This can be attributed to the heterogeneous nano-rods, the enhanced
interfacial adhesion to the PU matrix, and co-reaction between filler and polymer.

Table 4.3 Storage modulus values (G’) at 50°C and 100°C and the peak temperatures of
the tan δ curves for soft and hard segments.
G' (MPa) at
50
o
C
Samples
G' (MPa) at
100
o
C
T
g
, SS
(
o
C)
T
g
, HS
(
o
C)
Neat PU 21.06 10.37 -54.95 190
PU/ATT-MDI-H500A1M1H 24.54 12.21 -55.63 196.7
PU/ATT-MDI-H600A5M4H 23.83
91

11.96 -55.52 195.8
PU/ATT-MDI-H700A3M2H 23.63 11.86 -55.50 195.1

92
4.3.2.4 Tensile properties
The tensile yield strength, Young’s modulus, and elongation at break values for
PU and PU/ATT-MDI nanocomposites are tabulated in Table 4.4. Compared with neat
PU, the yield strength and modulus of PU/ATT-MDI nanocomposites were increased by
the clay additions, although elongation at break values were decreased. Table 4.4 shows
that the Young’s modulus and tensile yield strength of the PU/ATT-MDI-H500A1M1H
nanocomposite increased by 60%, and 19% respectively, while increases for the
PU/ATT-MDI-H700A3M2H nanocomposite were more modest - 38%, and 0%,
respectively. The PU/ATT-MDI-H500A1M1H nanocomposite resulted in the greatest
yield strength and Young’s modulus values among the three PU/ATT-MDI
nanocomposites, as expected, because of the highest grafting amount of MDI molecules
onto the ATT-MDI nano-rods. However, the values of elongation at break for PU/ATT-
MDI nanocomposites were lower than neat PU and decreased with the decreasing of
grafting content. In particular, the PU/ATT-MDI-H700A3M2H nanocomposite resulted
in the lowest value of elongation at break. This low elongation is attributed to an
insufficient amount of grafting on H700A3M2H ATT-MDI nano-rods and the tendency
to agglomerate in the PU matrix (Fig 4.5(c)). This agglomeration in PU/ATT-MDI-
H700A3M2H nanocomposite was responsible for the marked reduction in elongation.
Note that the melting point of soft segment (T
m
, SS) is close to room temperature.
Thus, the tensile properties at room temperature are directly related to the soft segment
characteristics. Korley et al. [27] reported that the presence of crystallites within the soft
segment would absorb strain energy upon deformation. Thus, semicrystalline soft
segments can act as an effective load-bearing phase during deformation [27]. As shown
by the DSC measurements, the ATT-MDI nano-rods acted as nucleation seeds and
induced crystallinity in the soft segment, resulting in an increase in the hard segment
domain size and a decrease in the soft segment domain region. Consequently, the tensile
yield strength and Young’s modulus of PU/ATT-MDI nanocomposites were increased,
but the elongation was reduced as the crystallinity of the soft segment increased.

Table 4.4 Tensile properties of PU and PU/ATT-MDI nanocomposite.

Young’s Modulus
(MPa)
Yield Strength
(MPa)
Elongation at
break (%)
Samples
Neat PU 7.14 0.93 335
PU/ATT-MDI-H500A1M1H 11.4 1.11 297
PU/ATT-MDI-H600A5M4H 10.35 1.01 214
PU/ATT-MDI-H700A3M2H 9.82 0.92

177

93
94
4.4 Conclusions
ATT-MDI nano-rods with different grafting contents were prepared by heat and
acid treatment of ATT followed by grafting with MDI. The resulting PU/ATT-MDI
nanocomposites exhibited increased crystallinity and strength compared with the neat PU.
These findings were attributed to the chemical modification of ATT, especially for
H500A1M1H ATT-MDI nano-rods, which led to improved dispersion in the PU matrix
and a strong chemical bonding between the dispersed and continuous phases. A higher
grafting content on the nano-rod surfaces led to more uniform dispersion and enhanced
tensile properties of the nanocomposites. However, the grafting content had negligible
effect on the thermal behavior of PU/ATT-MDI nanocomposites. The results demonstrate
that small additions of nano-rods will afford opportunities to design polymers with
enhanced properties for adhesives and composite materials. Further improvements and
applications in specific polymer matrices are likely to be possible with optimized surface
treatments and suitable modifiers.

95
Chapter 4 References
1. Wang Chia-Hao., Shieh Yeong-Tarng., Guo Gangjian., Nutt Steven., Euro Polym J
(submitted).

2. Mifsud A., Rautureau M., Fornes M., Clay Miner 1978, 13, 367.
3. Khorami J., Lemieux A., Thermochim. Acta 1989, 138, 97.
4. Kiyohiro T., Otsuka R., Thermochim. Acta 1989, 147, 127.
5. Shuali U., Yariv S., Steinberg M., Muller-Von-moons M., Kahr G., Rub A., Clay
Miner 1990, 25, 107.

6. Shuali U., Yariv S., Steinberg M., Muller-Von-moons M., Kahr G., Rub A., Clay
Miner 1991, 26, 497.

7. Frost R.L., Ding Z., Thermochim. Acta 2003, 397, 119.
8. Mendelovici E., Clays Clay Miner 1973, 21, 115.
9. Myriam M., Saurez M., Martin-Pozas J.M., Clays Clay Miner 1998, 46, 225.
10. Vicente-Rodríguez M.A., Lopez-Gonzalez J.D., Bañares-Muñoz M.A., Clay Miner
1994, 29, 361.

11. Vicente-Rodríguez M.A., Suarez Barrios M., Lopez-Gonzalez J.D., Bañares-Muñoz
M.A., Clays Clay Miner 1994, 42, 724.

12. Vicente-Rodríguez M.A., Suarez M., Bañares-Muñoz M.A., Lopez-Gonzalez J.D.,
Spectrochim. Acta 1996, 52A, 1685.

13. Yosomiya R., Morimoto K., Suzuki T., J. Appl. Polym. Sci 1984, 29, 671.
14. Prisacariu C., Olley R.H., Caraculacu A.A., Bassett D.C., Martin C., Polymer 2003,
44, 5407.

15. Blanco C., Herrero J., Mendioroz S., Pajares J.A., Clays Clay Miner. 1988, 36, 364
16. Pattanayak Asim., Jana Sadhan C., Polymer 2005, 46, 5183
17. Kuptsov A.H., Zhizhin G.N., Handbook of fourier transform Raman and spectra of
polymers, Elsevier, New York, 1988.

18. Chattopadhyay D.K., Raju K.V.S.N., Polym Sci 2007, 32, 352
96
19. Speckhard T.A., Hwang K.K.S., Yang C.Z., Laupan W.R., Cooper S.L., J Macromol
Sci Phys B 1984, 23,175.

20. MacKnight W.J., Yang M., Kajiyama T., Polym Prepr (Am Chem Soc, Div Polym
Chem) 1968, 9, 860.

21. Yang C.Z., Hwang K.K.S., Cooper S.L., Makromol Chem 1983,184, 651.
22. Yilgor I., Riffle J.S., Wilkes G.L., McGrath J.E., Polym Bull (Berlin) 1982, 8535.
23. Li Y., Tong G., Liu J., Linliu K., Desper C.R., Chu B., Macromolecules 1992, 25
7365.

24. Nunes R.C.R., Pereira R.A., Fonseca J.L.C., Pereira M.R., Polym Testing 2001, 20,
707.

25. Mathew A.P., Dufresne A., Biomacromolecules 2002, 3, 609.
26. Auad María L., Contos Vasili S., Nutt Steven., Aranguren Mirta I., Marcovich Norma
E., Polymer International 2008, 57, 651.

27. James Korley., LaShanda T., Pate Brian D., Thomas Edwin L., Hammond Paula T.,
Polymer 2006, 47, 9, 3073.

97
Chapter 5 Effects of Heat and Acid Treatments on Attapulgite [1]

5.1 Motivation
In the chapter 4, we have found that after heat and acid treatment at various
conditions, ATT showed different grafting contents with MDI and this was attributed to
the different contents of active site on ATT. In addition to be fillers, ATT was also
usually applied to heterogeneous catalytic reactions at elevated temperature and acid
environment. Moreover, there is a matter of some controversy on the specific
composition of water for each state in the thermal analysis. Caillere et al [2] described
water loss in the following three stages: (1) hygroscopic and zeolitic water is lost below
200
o
C, (2) bound water is lost between 250 and 400
o
C, and (3) hydroxyl groups are lost
above 400
o
C. In contrast, Prost [3] indicated that the loss of bound water in ATT occurs
in two steps. In the first step, half of the bound water is removed between 150 and 300
o
C,
and subsequently the remainder is removed along with the hydroxyl groups at
temperatures above 300
o
C. In an attempt to clarify the disparate descriptions offered in
the two previous studies [2, 3], and understand changes in the physicochemical properties
resulting from thermal and acid treatments, the effects of heat and acid treatments on
ATT were investigated in this chapter by using multiple analytical techniques.

98
5.2 Experimental
5.2.1 Materials, heat- and acid- treated ATT
Attapulgite (Attagel 50, Engelhard Co., USA) was subjected to purification with
sodium polyacrylate [4] having molecular mass of 4,000-5,000. The loss of hygroscopic
water and zeolitic water which occurred at low temperatures had less effect than bound
water and hydroxyl group on the structure of ATT. Thus, the present study was
performed at high temperatures (above 500
o
C) to investigate the effect of the loss of
bound water and hydroxyl group on the structure and composition of ATT. Heat and acid
treatments of ATT samples were conducted according to the following steps. First, 20-
gram quantities of ATT powder were each heated to 500, 600, 700, and 800
o
C for 1h,
then cooled to room temperature in a desiccator. For the convenience, the samples heat-
treated at 500, 600, 700, and 800
o
C were denoted as HT500, HT600, HT700, and HT800,
respectively. Characterization by thermogravimetric analysis (TGA), Fourier transform
infrared spectroscopy (FTIR), and wide-angle X-ray diffraction (WAXD) revealed that
the HT500, HT600, and HT700 exhibited similar characteristics, but the HT800 showed a
drastic change in crystal structure. Therefore, HT500 and HT800 were chosen to
investigate the effects of acid treatment. Samples of each were soaked in 5M
hydrochloric acid at 80
o
C for 1h or 8h to produce hydroxyl groups on the nano-rod
surfaces and to remove residual sodium ions and impurities. After heat and acid
treatments, sample designations were modified. For example, HT500 became
HT500AT1H, for heat treatment at 500
o
C followed by acid treatment for 1h. After
exhaustive washing with de-ionized water, the ATT nano-rods were dried at 80
o
C under
99
vacuum for 24h. The dried ATT was ground to powders and screened through a 325-
mesh sieve for further analysis.

5.2.2 Characterization
TGA (TGA 2050, TA Instruments) was performed to analyze the dehydration and
dehydroxylation for untreated and heat-treated ATT samples. In TGA, samples were
heated from room temperature to 800
o
C at 10
o
C/min under flowing nitrogen.
FTIR spectroscopy (Nicolet 4700) with a sensitivity of 128 scans and a resolution
of 4 cm
-1
was performed on the heat-treated and acid-treated ATT samples. KBr/ATT
disks were prepared and used to obtain the IR spectra.
WAXD (Bruker X-ray diffractometer, model D8 Advance) was performed to
analyze the structure of the heat-treated and acid-treated ATT. The analysis was carried
out with Cu K- α radiation (0.154 nm), and a 2 θ scan from 2.5° to 40°.
The morphologies of acid-treated ATT were observed by atomic force microscopy
(AFM NT-MDT Solver PRO-M) and transmission electron microscopy (TEM Philips
EM420). The AFM samples were prepared using gentle sonication in de-ionized water to
create a colloidal suspension. A drop of colloid was placed on a substrate and then dry-
analyzed under tapping mode of AFM. The TEM samples were also made from the
colloidal suspension and placed on a carbon film support grid and dry-analyzed at an
acceleration voltage of 120kV.

100
5.3 Results and Discussion
5.3.1 TGA measurement for heat-treated ATT
The weight loss of the untreated ATT measured by TGA is summarized in Table
5.1. The weight loss values indicate that dehydration and dehydroxylation during heating
from 25 to 800
o
C occurred in 4 stages. The first stage, below 120
o
C, resulted in water
loss of 10.1%, and was associated with the evaporation of hygroscopic water. During the
second stage, the 3.1% weight loss from 120 to 250
o
C was associated with the loss of
zeolitic water. In the third stage, ranging from 250 to 550
o
C, the 4.5% weight loss was
attributed to the dehydration of zeolitc water and bound water. During the fourth and
final stage from 550-800
o
C, the more strongly bound water and the hydroxyl groups were
lost.
Theoretically, an ATT molecule contains four zeolitic water molecules, four bound
water molecules, and two hydroxyl groups [5]. Assuming that the hydroxyl groups were
totally lost during heat treatment at 800
o
C, the 4.9% weight loss in the final stage (550-
800
o
C) can be assigned to the release of 2 molecules of hydroxyl groups and 1.99
molecules of bound water on the basis of stoichiometric calculations. The 4.5% weight
loss in the third stage between 250-550
o
C can also be stoichiometrically calculated to be
the loss of 2.01 bound water molecules and 1.55 zeolitic water molecules. The 3.1%
weight loss occurring in the second stage (between 120-250
o
C) can thus be quantitatively
attributed to the loss of 2.45 molecules of zeolitic water.
Figure 5.1 shows TGA curves for heat-treated ATT samples (HT500, HT600,
HT700, and HT800). The weight loss of the ATT heat-treated below 120
o
C was
attributed to evaporation of a small amount of adsorbed moisture. Between 120 and
550
o
C, the small amount of weight loss was caused by the loss of whole zeolitic water
and partial bound water. This observation can be clearly seen in HT500, HT600 and
HT700, an indication that after heat treatment at 500, 600, or 700
o
C, some zeolitic water
and bound water remained in the nano-rods. In the range of 550 to 800
o
C, the weight loss
was attributed to the loss of strongly bound water and the dehydroxylation of hydroxyl
groups from octahedral sheets. This scenario was not observed for the sample heat treated
at 800
o
C, which led to removal of all zeolitic water, bound water, and hydroxyl groups.

Table 5.1 Thermal analysis of attapulgite at various stages.

Step of dehydration and dehydroxylation
1 2 3 4
Range of temperature (
o
C) ＜120120-250 250-550 550-800
Weight loss (%) 10.1
Note: H
2
O
(z)
: zeolitic water; H
2
O
(b)
: bound water; OH: hydroxyl group
3.1 4.5 4.9
Water molecule number 2.45 H
2
O
(z)
1.55 H
2
O
(z)
+ 2.01 H
2
O
(b)
1.99 H
2
O
(b)
+ 2 OH

101

0 100 200 300 400 500 600 700 800
90
92
94
96
98
100
800
o
C
700
o
C
600
o
C
500
o
C
Weight Loss(%)
Temperature(
o
C)

Figure 5.1 TGA curves for the ATT nano-rods after heat treatments at various
temperatures.

102
103
5.3.2 FTIR characterization for heat-treated ATT
Because FTIR spectra are sensitive to the changes in OH vibrations, it has been
widely used to provide a characterization of clay minerals. For investigation of the
thermal behavior of ATT, FTIR spectra (Fig 5.2) of ATT can be divided into three major
regions. First, the high-frequency region (4000~3000 cm
-1
) is the stretch vibrations of
structural OH, OH of bound water, and hygroscopic water. Second, the mid-frequency
region (1700~1600 cm
-1
) is related to the bending vibration of hygroscopic, adsorbed,
bound water and structural OH. Third, the low-frequency region (1300~400 cm
-1
) is
assigned to the combination of stretch and bending vibrations of Si-O or Al-O framework
and octahedral M-OH (where M signifies a metallic element).
In the high-frequency region, the adsorption band at 3628 cm
-1
is assigned to Al
3+
-
OH. Nevertheless, the intensity of this absorption band decreased with increasing heat
treatment temperatures from 500 to 800
o
C. This finding is consistent with the TGA data
(Fig 5.1) in the range of 550 to 800
o
C. Within this temperature range, the weight loss
which resulted from the dehydroxylation of the octahedral sheet decreased with
increasing heat treatment temperature. In addition, one absorption band near 3420-3460
cm
-1
was observed which was invariant for the different heat-treated samples, and this
band was attributed to hydrogen bound to oxygen in Si-O-Si.
In the mid-frequency region, the FTIR spectra were almost identical for the four
heat-treated samples. The band that appeared at 1655 cm
-1
in untreated ATT shifted to
1639 cm
-1
after heat treatment. Mendelovici suggested that this shifting was related to the
energetically different hydrogen bondings [6]. Moreover, an adsorption band located at
1430-1426 cm
-1
for untreated ATT, HT500, HT600, and HT700
o
C was attributed to CO
3

104
stretching of calcite and dolomite [7]. However, this band shifted to 1393cm
-1
after heat
treatment at 800
o
C. This shift was attributed to the low degree of crystallinity of the
reformed calcite, which was recarbonated from the full decarbonation of calcite at 800
o
C
[8].
In the low-frequency region, a characteristic band that appeared near 1195 cm
-1
has
been assigned to the asymmetric stretch vibration of Si
U
-O-Si
D
(where U signifies that the
apical oxygen of the tetrahedron points upward in the direction normal to the sheet, and D
indicates that the apical oxygen points downward in the direction normal to the sheet) [9].
However, this peak disappeared after heat treatments. This finding can be associated with
changes of the Si-O-Si bond in the folded structure to perturb the Si-O vibration. In
addition, three absorptions at 985, 1030, and 1195 cm
-1
merged into one characteristic
band at 1030 cm
-1
for HT500, HT600, and HT700, although this band shifted to 1076 cm
-
1
for HT800. This band in the first three samples can be attributed to the perturbation of
the Si-O vibration in the folded structure after dehydration and dehydroxylation, while
the crystal structure was destroyed for HT800.

4000 3500 3000 2500 2000 1500 1000 500
1430
1393
1639
1655
e
d
c
b
a
Transmittance(%)
Wavenumber(cm
-1
)

Figure 5.2 FTIR curves for the ATT nano-rods (a) before, and after heat treatments at (b)
500
o
C, (c) 600
o
C, (d) 700
o
C, and (e) 800
o
C.

105
106
5.3.3 XRD analysis for heat-treated ATT
The WAXD patterns of untreated ATT and heat-treated ATT samples are shown in
Fig 5.3. The main peaks of untreated ATT appeared at 2 θ = 8.3°, 13.6°, 19.7°, and 29.4°.
Two peaks for quartz at 20.85° and 26.6°, and one peak for dolomite (MgCO
3
, CaCO
3
) at
30.9° were also present. After the heat treatment at 500
o
C for 1hr, the peaks for ATT at
2θ = 8.3° and 19.7° disappeared due to the irreversible folding of ATT (Fig 5.4). The
weak broad peak appearing at 2 θ = 19-21.5° was attributed to the destruction of the
internal structure. The intensity of this broad diffraction maximum appeared to become
weaker and more diffuse with increasing treatment temperature. The diffraction maxima
at 9.6° and 18.9° indicate the presence of ATT “anhydride” [10], but are not present for
HT500, HT600, HT700, and HT800. This suggests that the ATT “anhydride”
disappeared for heat treatments at temperatures greater than 500
o
C. Hayashi et al [11]
also reported the disappearance of the diffraction maxima at 9.6° and 18.9° and attributed
this to the destruction of the ATT “anhydride” phase. Although the dolomite has a
diffraction peak at 2 θ = 30.9°, the dolomite transformed to calcite as the temperature rose
to 600
o
C [12], and the peak completely disappeared at 800
o
C due to calcination of calcite.
This result is consistent with FTIR spectra, which show a peak at 1430 cm
-1
from CO
3

stretching of calcite and dolomite. Moreover, after heat treatment at 800
o
C for 1h, there
was no sharp peak appearing in the WAXD pattern (other than the two peaks for quartz at
20.85° and 26.6°). This finding indicates that the crystal structure underwent drastic
change and became amorphous for HT800.

107
40 020
020 40
Intensity

2 Theta(degree)
a
b
c
d
e

Figure 5.3 WAXD patterns for ATT (a) without heat treatment, with heat treatments at (b)
500
o
C, (c) 600
o
C, (d) 700
o
C, and (e) 800
o
C.

Figure 5.4 Morphological changes of ATT after heat treatment at high temperatures.

108
109
5.3.4 Morphology for acid-treated ATT
The morphologies of acid-treated ATT were revealed by AFM (Fig 5.5) and TEM
(Fig 5.6), respectively. The AFM image in Fig 5.5 shows sample of HT500 treated in 5M
hydrochloride acid for 1h (HT500AT1H). The AFM image (height data), shows that the
ATT nano-rod is partially degraded and exhibits a distinctly segmented morphology and
a rough surface. This scenario is due to the dissolution of octahedral layer of ATT after
acid treatment.
Figure 5.6 shows a TEM image from the sample of HT800 treated in 5M
hydrochloric acid for 8h (HT800AT8H). The external morphology of the acid-treated
ATT shows that the surface condition is preserved even after the structure becomes
amorphous phase. A rough surface with some pits in the ATT nano-rods was also
observed, consistent with the AFM image. The persistence of the fibrous morphology
indirectly shows that the acid treatment affected structure throughout the internal
microchannels.

Figure 5.5 AFM image for acid-treated sample HT500AT1H.

110

Figure 5.6 TEM image for acid-treated sample HT800AT8H.

111
112
5.3.5 FTIR characterization for acid-treated ATT
The changes in ATT functional groups after acid treatment in 5M HCl for 1h or 8h
were revealed by FTIR spectroscopy. For example, in Figure 5.7(a), the sample of
HT500AT1H reveals a spectrum similar to untreated ATT (Fig 5.2(a)) but unlike HT500
(Fig 5.2(b)). This finding can be attributed to the rehydration of HT500 in the acid media.
Most of the characteristic bands had been recovered from the acid treatment except the
peak at 1430 cm
-1
(CO
3
stretching of calcite and dolomite). This can be explained by the
reaction between calcite (CaCO
3
) and hydrochloric acid, and most of the CaCO
3
was
fully reacted. A similar effect of rehydration was also reported by Hayashi [11].
Hydrochloric acid can reportedly leach the majority of the octahedral cations within
ATT and produce silanol groups (Si-OH) (Fig 5.8) at 955-960 cm
-1
. The Si-OH group
also has been verified as a reactive site in ATT [6]. Moreover, Cai et al [13] found that
hydrochloric acid can diminish the masking effect in which OH bending vibrations
overlap with Si-O stretching bands in the range of 700-950 cm
-1
. The absorption band
near 798 cm
-1
was assigned as the symmetric vibration of Si
U
-O-Si
D
[14], and became
more intense after the acid treatment (compared to heat-treated samples). The intensity of
the peak also increased significantly when the acid treatment was increased from 1h to 8h.
The merging of the bands occurred not only in heat-treated samples but also in acid-
treated samples. After 8h of acid treatment in HT500 and HT800 (Fig 5.5(b) & (d)) or 1h
of acid treatment in HT800 (Fig 5.5(c)), the three absorption wells at 987, 1035, and 1194
cm
-1
were merged into one characteristic band at 1084 cm
-1
. This observation signifies
that the ATT crystal structure had been destroyed and replaced by an amorphous phase
[13]. In the sample HT500, the characteristic band appeared at 1030 cm
-1
but shifted to
113
1084 cm
-1
after 8h acid treatment. However, the sample of HT800 showed a small shift
from 1076 cm
-1
to 1088 cm
-1
,indicating that the crystal structure of HT800 was destroyed
prior to acid treatment. From these observations, the merging of three absorption bands
into a single band and shifting from 1030 cm
-1
to 1076-1088 cm
-1
can be an indicator of
structural change in ATT.

4000 3500 3000 2500 2000 1500 1000 500
d
c
b
a
Transmittance(%)
Wavenumber(cm
-1
)

Figure 5.7 FTIR spectra for ATT nano-rods after heat and acid treatments: (a)
HT500AT1H (b) HT500AT8H (c) HT800AT1H (d) HT800AT8H.

114

Figure 5.8 The mechanism of acid-activation for Attapulgite.

115
116
5.3.6 XRD analysis for acid-treated ATT
The WAXD patterns for HT500 and HT800 treated in 5M HCl for 1h or 8h are
shown in Fig 5.9. After the acid treatment on sample HT500, the peak at 2 θ = 29.4° (Fig
5.5(b)) disappeared, but a small broad maximum (2 θ = 29.2-31.2°) was observed (Fig
5.9(a) & (b)). The broad peak can be ascribed to the dissolution of the octahedral layer.
The intensity of the broad peak gradually became weaker and more diffuse after acid
treatment for 8h. However, the sample HT500AT1H (Fig 5.9(a)) revealed a small sharp
peak at 2 θ = 19.7° which belongs to main peak of ATT but does not appear in HT500
(Fig 5.3(b)). This finding can be attributed to the rehydration of HT500 in the acid media
for 1h, and this assertion was verified by FTIR spectra.
An additional diffraction peak arising from dolomite (MgCO
3
, CaCO
3
) at 2 θ =
30.9° also vanished in the acid-treated samples. This finding was consistent with FTIR
observations, which indicated that the dolomite was completely reacted by hydrochloric
acid. Comparing the heat-treated 800 sample (HT800) and the acid-treated 800 samples
(HT800AT1H and HT800AT8H), the WAXD patterns are quite similar. However,
stronger diffraction peaks corresponding to quartz (at 20.85° and 26.6°) were observed in
the acid-treated 800 samples. This is due to the increase of insoluble impurities in ATT
after acid treatment. The disappearance of the peak at 2 θ = 29.4°indicates that the crystal
structure underwent a gradual vitrification transformation to an amorphous phase during
extended acid treatments.

010 20 30
010 20 30
a
Intensity
2 Theta(degree)
b
c
d

Figure 5.9 WAXD patterns for (a) HT500AT1H (b) HT500AT8H (c) HT800AT1H (d)
HT800AT8H.

117
118
5.4 Conclusions
In this work, attapulgite was heat-treated at 500-800°C, followed by refluxing in
HCl solutions to determine changes in surface structure, crystal structure, and
composition. Thermal analysis of the untreated attapulgite revealed dehydration and
dehydroxylation, and clarified the controversy surrounding the specific composition of
water molecules (e.g., bound water versus zeolitic water) released during successive
stages of heating. The dehydration and dehydroxylation during heat treatment of
attapulgite led to structural changes, as demonstrated by FTIR spectra and WAXD
patterns. FTIR spectra also revealed that acid treatment hydroxylated the heat-treated
attapulgite. AFM and TEM images demonstrated that acid treatment caused surface
roughening and pitting, although the fibrous structure was retained. This work provides
insight into physicochemical changes resulting from thermal and acid treatments. This
understanding is necessary to extend applications of attapulgite to specific polymer
systems and to heterogeneous catalytic reactions.

119
Chapter 5 References
1. Wang Chia-Hao., Shieh Yeong-Tarng., Nutt Steven., Appl Clay Sci (submitted).

2. Caillere S., Henin S., X-ray Identification and Crystal Structures of Clay Minerals
1961, 343.

3. Prost R., Etude de l’ hydratation des argiles: interactions eau-mineral et mécanisme
de la rétention de l’ eau. Ph.D. Thesis, Université Pierre et Marie Curie, Paris VI 1975.

4. Purcell Jr, R.J., Parker D.C., US patent 2002, 6, 444, 601.
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6. Mendelovici E., Clays Clay Miner 1973, 21, 115.
7. Bukka K., Miller J.D., Shabtai J., Clays Clay Miner 1992, 40, 92.
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10. Preisinger A., Clays Clay Miner 1963, 10, 365.
11. Hayashi H., Otsuka R., Imai N., Am Miner 1969, 53, 1613.
12. Bayram H., Önal M., Üstünı şık G., Sarıkaya Y., J. Therm. Anal. Cal 2007, 89, 169.
13. Yuanfeng Cai., Jiyue Xue., Polya D.A., Spectrochimica Acta Part A 2007, 66, 282.
14. Song G.B., Liu F.S., Cao Y.G., Peng T.J., Dong F.Q., Wan P., Acta Petrologica
Sinica 1999, 15, 469.

120
Chapter 6 Conclusions and Suggestions for Future Works

6.1 Conclusions
This work has demonstrated that the surface chemistry of a nano-clay
significantly influences the properties of polyurethane based nanocomposites. The factors
which contribute to the mechanical and thermal behavior of polyurethane
nanocomposites containing montmorillonite (Cloisite 30B and Nanomer I.30E) or self-
modified attapulgite (ATT-MDI) were described, including the preferential modifier
between organophilic nano-clay and polyurethane, different soft segment molecular
weight in PU/MMT nanocomposites, and the functionalization of attapulgite, different
grafting amounts in PU/ATT nanocomposites.
The organophilic nano-clay showed uniform nano-scale dispersion for both
PU/MMT and PU/ATT nanocomposites. The results confirmed that regardless of the
nano-clay or the PU matrix, the nature of the filler surface chemistry dominated the
properties that were achieved. In sum, understanding the chemistry, the interfacial phase
between filler and matrix, and the physicochemical change of nano-clay is essential for
developing and designing nanocomposites with specific combinations of properties.

6.2 Suggestions for Future Works
As a new family of materials, nanocomposite is a good candidate for both
fundamental research and practical applications. In practice, the existence of layered
silicates in nanocomposites is more complex than initially believed. Much work remains
121
to be carried out to fully understand factors such as exfoliation versus intercalation, the
driving forces that yield different nanocomposite structures, and nonequilibrium
phenomena such as irreversible aggregation and percolation. In addition, a better
understanding of the structure-property relationships must be fulfilled in areas such as fire
retardancy and physico-mechanical properties.
Apart from conventional phenomenal observations and theoretical analyses, the
experimental techniques also can be improved. For example, microwave-exfoliation is a novel
processing technique that involves the use of both microwave radiation and polar agents (e.g.
water, 2-propanol). In this technique, the layered silicates are pretreated by a polar agent, then
mixed with polymer, followed by the microwave treatment. The polar molecules which penetrate
into the clay galleries are intensely heated and quickly evaporated by the microwave radiation. As
a result, a high internal pressure is created, and the repulsive force generated by the “popcorn”
process overcomes the attractive force between silicate layers and facilitates exfoliation of the
clay. The basic idea and mechanism of this technique is shown in Fig 6.1. Although this
innovative technique points out a new way to achieve exfoliation in polymer-layered silicate
nanocomposites, more work (e.g. optimizing the microwave treatment, choosing of clay, polymer
and polar agent) must be carried out to develop and refine the process before it can be a viable
process in practice.

(a)
Polymer

(b)

(c)
Fig. 6.1 Scheme of the microwave-exfoliation nanocomposite (a) Clay with polar
molecules in the galleries (b) Popcorn effect of polar molecules under microwave
treatment (c) Exfoliated silicate layers in nanocomposite.
122
123
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Abstract (if available)

Abstract Polyurethane-layered silicate nanocomposites with different layered silicates, montmorillonite (MMT) and attapulgite (ATT), were investigated to gain fundamental understanding of the role of nanofillers, the chemistry of the modifiers, and the physics of polymer nanocomposites.

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Asset Metadata

Creator Wang, Chia-Hao (author)

Core Title Nanoclay-reinforced polyurethane

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Materials Science

Degree Conferral Date 2009-12

Publication Date 10/06/2009

Defense Date 09/02/2009

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag nanocomposites,OAI-PMH Harvest,polyurethane

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Nutt, Steven R. (committee chair), Sammis, Charles G. (committee member), Shing, Katherine S. (committee member)

Creator Email chiahao_wang@yahoo.com.tw,chiahaow@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-m2646

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