Polyurethane/Polybenzoxazine Based Shape Memory Polymers
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POLYURETHANE-POLYBENZOXAZINE BASED SHAPE MEMORY POLYMERS A Thesis Presented to The Graduate Faculty of The University of Akron In Partial Fulfillment of the Requirements for the Degree Master of Science Numan Erden December, 2009
description
A thesis work on the synthesis and chracterization of polyurethan/polybenzoxazine polymers.
Transcript of Polyurethane/Polybenzoxazine Based Shape Memory Polymers
POLYURETHANE-POLYBENZOXAZINE BASED
SHAPE MEMORY POLYMERS
A Thesis
Presented to
The Graduate Faculty of The University of Akron
In Partial Fulfillment
of the Requirements for the Degree
Master of Science
Numan Erden
December, 2009
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POLYURETHANE-POLYBENZOXAZINE BASED
SHAPE MEMORY POLYMERS
Numan Erden
Thesis
Approved:
_______________________________
Advisor
Dr. Sadhan C. Jana
_______________________________
Faculty Reader
Dr. Kevin Cavicchi
_______________________________
Faculty Reader
Dr. Robert A. Weiss
Accepted:
_______________________________
Department Chair
Dr. Sadhan C. Jana
_______________________________
Dean of the College
Dr. Stephen Z.D. Cheng
_______________________________
Dean of the Graduate School
Dr. George R. Newkome
_______________________________
Date
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ABSTRACT
Shape memory polyurethanes (SMPUs) have attracted much attention from academic and industrial
researchers due to strong potential in biomedical and consumer applications. Some of the limiting factors of
these materials are low recovery stress (RS) and shape recovery (SR). Fundamental studies have focused on
the improvement of RS and SR values using primarily two approaches. The first utilizes the nanocomposite
route by which a few weight percentages of nanofillers are added to SMPU in order to increase the
modulus and consequently to obtain enhancement in recovery stress. Although successful in the case of
SMPU with amorphous soft segments, the nanofillers caused reduction in crystallinity of crystalline soft
segment leading to deterioration of shape memory properties of SMPUs. In the second approach, chemical
additives are added which either chemically bond with SMPU chains or form a separate phase and offer
much stronger modulus than the soft and hard segments of SMPU. This second approach was followed in
the current study.
Polybenzoxazine (PB-a) was incorporated into a thermoplastic polyurethane (PU) formulation,
anticipating that it would play a similar role to hard segment and improve the shape memory properties. It
was found that benzoxazine monomer formed miscible blends with the prepolymer derived from 4,4'-
methylenebis (phenyl isocyanate) (MDI) and poly (tetramethylene) glycol (PTMG) with average molecular
weight of 650 g/mol. This allowed chain extension of prepolymer using 1,4-butanediol (BD) as in the
synthesis of regular polyurethanes. The benzoxazine was later polymerized into polybenzoxazine (PB-a) by
thermal curing at 180 °C in 3 hrs.
The results of this study showed that both RS and SR increased with the addition of benzoxazine. A
specimen with 17 wt. % benzoxazine produced the best RS and SR values with 13 MPa and 93%,
respectively compared to RS of 6.8 MPa and SR of 72% for polyurethane. The deformation conditions
were also found to exert significant influence on RS and SR values. Both stretching rate and stretching
temperature increased the RS values. However, higher heating rates caused a reduction of the values of RS.
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The stress relaxation experiments were carried out to establish a correlation between the deformation
conditions and the values of RS. It was found that specimens with 9 wt. % and 17 wt. % benzoxazine
experienced high degrees of stress relaxation. Consequently, the RS values of these specimens, although
higher than polyurethanes, were somewhat compromised. Furthermore, an investigation on surface
morphology revealed that the specimens had different levels of hard and soft segment phase separation.
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DEDICATION
Dedicated to my parents; Casim and Şükriye Erden for their everlasting affection and love...
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ACKNOWLEDGEMENTS
I would like to acknowledge all my friends from both my work and social surroundings for their
motivating words and sincere wishes. Research fellows of my group and the small Turkish community in
Akron are those whom I am deeply grateful.
I would like to extend my thanks to the committee members, Dr. Kevin Cavicchi and Dr. Robert
Weiss, due to their kind efforts for reviewing this thesis.
Finally, I would like to express my sincere thanks for my advisor, Dr. Sadhan C. Jana. This study
could not have been a complete work without his professional support.
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TABLE OF CONTENTS
Page
LIST OF TABLES .........................................................................................................................................ix
LIST OF FIGURES ......................................................................................................................................... x
CHAPTER
I INTRODUCTION ....................................................................................................................................... 1
II LITERATURE SURVEY ........................................................................................................................... 7
2.1 Shape memory effect (SME) ................................................................................................................ 7
2.1.1 Parameters for characterization of SMPs ...................................................................................... 9
2.1.2 Shape memory effect in alloys .................................................................................................... 11
2.1.3 Shape memory effect in ceramics ............................................................................................... 13
2.1.4 Shape memory effect in polymers .............................................................................................. 15
2.2 Classification of SME in polymers ..................................................................................................... 19
2.2.1 Thermal activation ...................................................................................................................... 19
2.2.2 Activation by light ...................................................................................................................... 20
2.2.3 Activation by electric/magnetic field .......................................................................................... 21
2.3 Shape memory polyurethanes (SMPUs) ............................................................................................. 22
2.3.1 Basic information on polyurethanes ........................................................................................... 22
2.3.2 SMPU blends .............................................................................................................................. 25
2.3.3 SMPUs with different co-monomers .......................................................................................... 26
2.3.4 Chemically crosslinked SMPUs.................................................................................................. 27
2.3.5 SMPU nanocomposites ............................................................................................................... 27
2.3.6 Shape memory alloy (SMA)/shape memory polyurethane (SMPU) composite ......................... 29
2.4 Polybenzoxazine ................................................................................................................................. 30
2.4.1 Synthesis of the benzoxazine ...................................................................................................... 31
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2.4.2 Polymerization ............................................................................................................................ 32
2.4.3 Distinguishing properties of polybenzoxazines .......................................................................... 33
2.5 Earlier polybenzoxazine/polyurethane work ...................................................................................... 34
III EXPERIMENTAL .................................................................................................................................. 38
3.1 Materials ............................................................................................................................................. 38
3.1.1 Raw materials for synthesis of polyurethane .............................................................................. 38
3.1.2 Raw materials for benzoxazine ................................................................................................... 39
3.1.3 Preparation of PU/PB-a systems ................................................................................................. 40
3.2 Characterization methods ................................................................................................................... 44
3.2.1 Thermal characterization ............................................................................................................ 45
3.2.2 Characterization of mechanical properties .................................................................................. 45
3.2.3 Spectroscopic analysis ................................................................................................................ 46
3.2.4 Analysis of morphology .............................................................................................................. 46
3.2.5 Characterization of shape memory properties ............................................................................. 46
IV RESULTS AND DISCUSSIONS ........................................................................................................... 48
4.1 Thermal properties .............................................................................................................................. 48
4.2 Thermomechanical properties ............................................................................................................ 55
4.2 Spectroscopic analysis ........................................................................................................................ 60
4.3 Mechanical properties ......................................................................................................................... 64
4.4 Shape memory properties ................................................................................................................... 66
4.4.1 Recovery stress and shape recovery ratio ................................................................................... 67
4.4.2 Effects of deformation conditions and heating rate .................................................................... 69
4.4.3 Shape fixity ................................................................................................................................. 78
4.4.4 Stress relaxation behavior ........................................................................................................... 79
4.5 Morphological properties.................................................................................................................... 83
V CONCLUSIONS ...................................................................................................................................... 88
REFERENCES .............................................................................................................................................. 90
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LIST OF TABLES
Table Page
2- 1 Frequently used reactants for polyurethane synthesis. ........................................................................... 23
3- 1 Corresponding molar ratios of raw materials ......................................................................................... 43
3- 2 Testing temperatures for evaluation of shape memory properties ......................................................... 47
3- 3 DMA set parameters for application recovery stress measurements. ..................................................... 47
4- 1 Glass transition onset values determined through DSC with higher heating rates. ................................ 54
4- 2 Viscoelastic properties of the samples ................................................................................................... 58
4- 3 TGA analysis results. Scan rate was 20 °C/min. .................................................................................... 59
4- 4 Hydrogen bonding % of urethane domains. AFCO refers to 1730 cm-1
and ANCO; 1700 cm-1
................. 64
4- 5 Tensile properties of the samples. .......................................................................................................... 64
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LIST OF FIGURES
Figure Page
1- 1 A simplified schematic for a single SMP cycle adapted from Ref.16. ..................................................... 2
2- 1 A typical SME cycle for thermally induced SMPs adapted from Ref.21. .............................................. 10
2- 2 Shape memory behavior characterization with bending test adapted from Ref. 59, 60. ........................ 11
2- 3 A modified schematic of SME mechanism in a typical SMA adapted from Ref.1. ............................... 12
2- 4 Polarization states in different phases of SMC adapted from Ref.71. .................................................... 14
2- 5 SME mechanism of shape memory ceramics (SMC) adapted from Ref.1. ............................................ 14
2- 6 Viscoelasticity of crosslinked polymers with respect to temperature. ................................................... 16
2- 7 SME activation of SMPs by light stimulus, adapted from Ref. 102. ..................................................... 20
2- 8 Two step polymerization of polyurethane adapted from Ref.16. ........................................................... 25
2- 9 General chemical structure of monofunctional 3,4 dehydro-2H-1,3 benzoxazines. .............................. 31
2- 10 The synthesis of mono functional 3,4-dehydro,2H,1,3-benzoxazine adapted from Ref. 38. ............... 31
2- 11 A schematic of synthesis and polymerization of the benzoxazine adapted from Ref. 38, 163............. 33
3- 1 Chemical formula of MDI. ..................................................................................................................... 38
3- 2 Chemical formula of PTMG. ................................................................................................................. 39
3- 3 Chemical formula of 1,4-butanediol. ..................................................................................................... 39
3- 4 General chemical formula for DABCO T120. ....................................................................................... 39
3- 5 Chemical formula of bisphenol-A. ......................................................................................................... 40
3- 6 Chemical formula of paraformaldehyde. ............................................................................................... 40
3- 7 Chemical formula of aniline. .................................................................................................................. 40
3- 8 Timeline for sample preparation in brabender. Step I-IV shows the sequence of addition of the
ingredients. ........................................................................................................................................... 43
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3- 9 The time-temperature plot showing duration and temperature of mixing, compression molding and
vacuum oven. ....................................................................................................................................... 44
4- 1 DSC scans of B-a monomer and precursors with 13% curing. Scan rate, 10 °C/min . .......................... 49
4- 2 DSC scans of Sample I. Scan rate, 10 °C /min. ..................................................................................... 49
4- 3 DSC scans of Sample II. ........................................................................................................................ 50
4- 4 DSC scans of Sample III. ....................................................................................................................... 53
4- 5 WAXD and SAXD measurement results of Sample III, respectively the first and the second images. . 54
4- 6 Loss tangent (tanδ) as a function of temperature for all samples. Heating rate, 4 °C/min; frequency, 1
Hz.; scan rate between, -50 °C and 150 °C. ......................................................................................... 55
4- 7 Storage modulus (E') as a function of temperature for all samples. Heating rate, 4 °C/min; frequency, 1
Hz.; scan rate between, -50 °C and 150 °C. ......................................................................................... 56
4- 8 Loss modulus (E'') as a function of temperature for each sample. Heating rate, 4 °C/min; frequency, 1
Hz.; scan rate between, -50 °C and 150 °C. ......................................................................................... 57
4- 9 The effect of curing time on Tg and E' of Sample II and III. .................................................................. 58
4- 10 TGA curves of the samples. Scan rate was 20 °C/min. ........................................................................ 59
4- 11 FT-IR spectra of benzoxazine monomer. ............................................................................................. 60
4- 12 1H-NMR analysis of benzoxazine monomer (bis (3-phenyl-3,4-dehydro-2H-1,3-benzoxazinyl)-
isopropane)........................................................................................................................................... 61
4- 13 FTIR spectra of polyurethane prepolymer. .......................................................................................... 62
4- 14 ATR spectra of the samples. ................................................................................................................ 63
4- 15 Curve fitting at 1733 cm-1
and 1703 cm-1
for Sample I, II, and III, respectively. ................................ 64
4- 16 Stress-strain diagram of the samples at room temperature. .................................................................. 65
4- 17 Stress-strain diagram of the samples at deformation temperatures. Sample I, 71 °C; Sample II, 85 °C;
Sample III, 110 °C. Max strain was 150%. .......................................................................................... 66
4- 18 Recovery stress behaviors of 100% strained samples. Heating rate was 4 °C/min and stretching rate
was 50 mm/min. ................................................................................................................................... 67
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4- 19 Shape recovery ratio of 100% strained samples. Heating rate was 4 °C/min and stretching rate was 50
mm/min. ............................................................................................................................................... 68
4- 20 Strain effect on recovery stress trend of Sample I. Stretching rate was 50 mm/min and stretching
temperature was 71 °C. ........................................................................................................................ 70
4- 21 Strain effect on recovery stress trend of Sample II. Stretching rate was 50 mm/min and stretching
temperature was 85 °C. ........................................................................................................................ 71
4- 22 Strain effect on recovery stress trend of Sample III. Stretching rate was 50 mm/min and stretching
temperature was 110 °C. ...................................................................................................................... 72
4- 23 Strain effect on shape recovery behavior of Sample I. Stretching rate was 50 mm/min and stretching
temperature was 71 °C. Heating rate was 4 °C/min. ............................................................................ 73
4- 24 Strain effect on shape recovery behavior of Sample II. Stretching rate was 50 mm/min and stretching
temperature was 85 °C. Heating rate was 4 °C /min. ........................................................................... 74
4- 25 Strain effect on shape recovery behavior of Sample III. Stretching rate was 50 mm/min and stretching
temperature was 110 °C. Heating rate was 4 °C/min. .......................................................................... 74
4- 26 The effect of stretching temperature on recovery stress of Sample II.. ................................................ 75
4- 27 Stretch temperature effect on recovery stress of Sample II. ................................................................. 77
4- 28 The influence of heating rate on recovery stress of 50% strained Sample II specimens. ..................... 78
4- 29 Shape fixity of the samples. ................................................................................................................. 79
4- 30 The influence of strain on relaxation behavior of Sample II. ............................................................... 80
4- 31 The influence of stretch rate on relaxation behavior of Sample II. ...................................................... 81
4- 32 Strain effect on the relaxation behavior of Sample III. ........................................................................ 82
4- 33 The influence of stretching rate on the relaxation attitude of Sample III. ............................................ 82
4- 34 AFM phase (a) and height (b) images of sample I, scan size; 5um. ..................................................... 84
4- 35 AFM phase (a) and height (b) images of sample II, scan size; 5um. ................................................... 85
4- 36 AFM phase (a) and height (b) images of sample III, scan size; 5um. .................................................. 85
4- 37 Simple sketches for the morphological features of the samples. .......................................................... 87
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CHAPTER I
INTRODUCTION
Shape memory polymers (SMP) have gained a great deal of attention in the field of smart materials
research during the last two decades [1-3]. In some applications they have been used to substitute shape
memory alloys [4]. Through understanding of material properties, they also find use in many important
applications [5]. With the advent and excellent prospect of improvement, this class of materials is slated to
have more widespread applications. Current [6,7] and proposed [8-10] applications cover a wide range of
us spanning from biocompatible implants to vehicle parts.
To qualify as SMPs, polymers must be able to perform desired functions under cyclic conditions of
loading/unloading and thermal changes. The effects of thermal swing and mechanical loading/unloading
are defined as shape memory effect (SME). In a typical shape memory cycle, SMPs must adopt a new
temporary shape and revert back to original shape under the influence of an external stimulus. In these
cases, the shape memory performance not only depends on the molecular structure, but also on the mode of
deformation and the programming of the stimulus application process [11]. Obtaining a new shape is
dependent on both the nature of stimulus and the mode of deformation. On the other hand, regaining the
original shape is always dictated by how the stimulus is applied. The stimulus here refers to heat, light,
chemical reaction, electricity, and magnetism. This may be applied independently, e.g. only heat or light, or
in complex combinations [12]. This is depicted in Figure 1- 1.
The process leading to change of shape or preservation of permanent shape from the temporary shape
requires that the materials should have at least two independent or mutually/synergistically working phases
[13]. One phase or domain should be reversible while the second phase should be in the active range of the
stimulus. This way, SME would be realized in the transition range as the reversible phase undergoes
changes from temporary to permanent shape. During this transition, polymer chains can take a desired form
upon deformation. Once the new temporary shape is attained, the stimulus is taken away swiftly and the
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polymer assumes the form. However, it should be noted that if the field that produced the stimulus takes
longer time to return to a reference state, polymer chains can relax leading to poorer shape memory
performance [14]. An example may be found in the case of thermal gradients, e.g. low thermal conductivity
of polymers lead to slower cooling & heating. The role of elastically stored energy [12] on shape memory
performance should be discussed at this point. Polymer chains store an important part of the energy applied
to deform them. This energy is used in the recovery of the original shape upon the application of the
stimulus. The heat dissipation during both deformation and recovery in a single cycle is also important and
accounts for a loss in shape. The extent of such a loss however depends on the rheological properties of the
materials [15].
Figure 1- 1 A simplified schematic for a single SMP cycle adapted from Ref.16.
In a single shape memory cycle, the most critical conditions are temperature at which deformation
takes place, e.g., stretching temperature, and the extent of deformation, e.g., strain level. These deformation
conditions along with the rheological properties of the polymers determine the recovery force and the
extent of shape recovery. This knowledge is useful in deciding suitable applications of the shape memory
polymers. As an example, for a shape memory alloy (SMA) to be replaced by an SMP, one should carefully
evaluate the magnitude of recovery stress the polymer develops. It is known in literature that SMPs offer
very low value of recovery stress. The thermal transition temperature of SMPs is usually low compared to
SMAs. Even SMPs with very high transition temperature may not work due to thermal degradation.
A great variety of polymers exhibit shape memory functions [17]. These include thermoplastics based
on consumer polymers such as polyethylene and polystyrene copolymers, to thermosets based on
polyurethanes and epoxides. Among these, polyurethanes have gained the greatest attention due to the very
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unique properties [18]. Some attributes are biocompatibility, biodegradability, vast possibilities for
synthesis and productions from readily available commercial raw materials, and affordable cost.
The shape memory polyurethanes (SMPU) exhibit remarkable shape memory effects (SME), due to
microphase segregation [19]. This is very much interrelated to the chemical nature and the production
method of polyurethanes. Typical polyurethanes are produced via insertion polymerization by which the
polymerization results in thermodynamically incompatible blocks of hard and soft segments. Thus, this
incompatibility creates a system in which different blocks constitute separate phases and form
microdomains. This inherent structure of polyurethanes facilitates thermal activation of shape memory, as
very different thermal transitions are associated with the soft and hard segments.
The usage of polyurethanes as SMPs suffers from several challenges. The materials experience
thermal degradation, hydrolytic degradation, and a dimensional instability after processing, thus offer
insufficient recovery stress, long recovery time and in some cases low value of shape fixity [20,21]. Many
of these challenges, however, are also seen in other SMPs [22].
Recovery stress is defined as the stress needed to hold a restrained specimen at fixed dimensions
while the specimen attempts to undergo shape recovery upon application of the stimulus [11]. This is an
important qualifying parameter to project particular applications of SMPs. For example, recovery stress in
the range of 150-300 MPa are produced by SMAs, while 1-3 MPa by many SMPs [16, 23].
Shape recovery time is another significant property of SMPs which refers to the time that an SMP
takes to recover from a strained state (temporary shape) to the original state (permanent shape) [13]. This is
totally dependent on the rheological properties of SMPs. Shape recovery time should be as sharp as
possible, so that the SMP article undergoes the fastest recovery of its temporary shape upon application of
the stimulus. A typical recovery time for SMPs may be around a few minutes, while it is less than one
second for SMAs. Attempts to shorten the recovery time have not met with much success.
Shape or strain fixity is another shape memory property that simply refers to the ratio of deformation
(e.g. elongation) before and after the temporary shape is formed upon the completion of deformation [21].
In the case of thermal activation, e.g., heat as stimulus, the temporary shape is obtained by heating up the
polymer above the activation temperature, stretching it, and then rapidly cooling it below the activation
temperature with the load in place. After the specimen is cooled, the stretching force is removed. Because
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of the inevitable relaxation in some cases, the length of the stretched specimen at the end of deformation
and after cooling is not the same. Consequently, the value of shape fixity is less than unity.
Many research groups have paid special attention to these crucial points on shape memory properties
[15,24-26]. First, it has been thought that the presence of another polymer as a blend component can help
[27,28]. Second, copolymerization of unusual monomer units have been used as alternative [29,30]. Third,
various fillers have been compounded with SMPs; carbon black [31,37], carbon fiber[31,32], functional
single or multiwall carbon nanotubes [33-35]. Lastly, composites of wires of SMAs with SMP matrices
have been attempted for the applications of bending actuation [36,37]. Nevertheless, only particular
improvements were achieved.
In this study, the focus was centered upon a partially reactive system in which polyurethane (PU) and
polybenzoxazine (PB-a) were allowed to react to an extent. Important parameters, especially recovery
stress, shape recovery, and the programming effects, e.g., the effects of deformation conditions, were
investigated. Thermal, mechanical, structural, and morphological test results are presented and correlated in
this thesis.
Polybenzoxazine (PB-a) is a thermosetting polymer that has gained surprising attention in the
polymer industry in the last fifteen years [38], even if it had been known since 1930s [39-40]. The level of
attention increased after a patent was granted to Ishida et al. [41] whereby the production method of a
particular class of this resin was enhanced [42]. Before this invention, PB-a was obtained by traditional
novolac resin production methods [43,44].
A typical benzoxazine is synthesized from three main reactants: an amine, an alcohol, and preferably
paraformaldehyde. The product mostly contains oligomers along with the monomer mixed in it. In view of
this, the mixture is simply referred as the precursor. The precursor undergoes thermally induced ring
opening polymerization through cationic bond cleavage at elevated temperatures. This is due to the strain
generated on the six member heterocyclic ring bearing O and N atoms. The strong basicity of N and O is
also considered to play an important role on this mechanism. This mechanism is thoroughly explained
elsewhere [38].
The thermosetting polymer, polybenzoxazine and its derivatives offer several significant features. As
an example, these materials demonstrate near zero shrinkage or slight expansion due to the hydrogen
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bonding network and established crosslinking, and void free products. Other useful attributes are
inexpensive raw materials, no need for strong acidic or basic catalysts, short curing periods, very low water
absorption, and very high glass transition [45]. On the other hand, an important disadvantage of these
materials is brittleness, which limits many industrial applications of this class. Some applications of PB-a
include adhesives, high temperature resistant composites, and highly flame retardant materials [46]. Some
major applications are found in electronics and electric insulations.
From a few earlier studies upon PU/PB-a systems [47-48], it is known that the presence of PU in PB-
a gives toughness to PB-a and eliminate the most important disadvantage arising from brittleness of PB-a.
The presence of PU also offers an increase in modulus, synergistic Tg enhancement, and elasticity
improvement [49].
The following points were considered as motivations for the development of shape memory PU/PB-a
systems. First, the network system that PB-a forms during curing and after interactions with PU domains
may act as the fixed phase of the potential shape memory. The network also provides the polymer with high
modulus, high Tg and mechanical integrity [50]. This network has been studied and found to coexist with
PU hard and soft segment, whereby it interacts both with hard and soft segments and additionally establish
its own domain depending on the ratio of functionality [51]. This network can be considered as a
supplementary fixed network in PU system. Note that a majority of shape memory polymers contain at
least two phases or domains to produce reversible and permanent shapes. In this context, PU/PB-a systems
offer an additional fixed phase, which can add to the ability to have more complex SMEs [52]. Second, the
PB-a networks, e.g., hydrogen and crosslinking, were expected to enhance the shape recovery force, due to
their contributions to high glass transition temperature. As stated previously, recovery force is a direct
result of entropic relations and based on elastically stored energy during deformation. To maximize the
storage of elastic energy, a matching combination of viscoelasticity is needed. It was expected that by
adding partially reactive PB-a into PU system, the viscoelasticity of the system would be dramatically
improved. Third, in addition to the PB-a forming its own network, a part of it can be chemically connected
to PU hard segment derived from 4,4’ bis (phenyl diisocyanate) methylene (MDI).
It was anticipated that polyurethane/polybenzoxazine system as shape memory materials should offer
unique properties. This thesis covers development of such materials by chemical reactions, processing, and
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characterization of shape memory performance. Particular attention was paid upon the effects of
deformation rate, extent and temperature, and heating rate on shape memory properties. In addition,
mechanical, thermal, and morphological properties were investigated.
The rest of this thesis is organized as follows. Chapter II contains a survey of literature. Foundations
and different mechanisms of SME in various materials are presented. In addition, an overview of the earlier
studies on SMPUs and PB-a are presented. Chapter III consists of various materials, experimental
techniques, and the details on PU/PB-a sample preparation. Chapter IV presents the results of experimental
work along with the discussions. Some concluding remarks are presented in Chapter V.
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CHAPTER II
LITERATURE SURVEY
The field of research on smart materials has grown considerably in past decades. Shape memory
polymers and alloys are good representatives of smart materials. These materials have become very popular
in the last two decades as they find applications in biomedical, textiles, electronics, and packaging
industries. The number of patent applications on shape memory materials reflects this growth [1-2].
The literature survey in this chapter is intended to give a short overview of shape memory effects in
various materials and particularly in polymers. In addition, classification of shape memory polymers,
earlier work on shape memory polyurethanes and their nanocomposites are presented. The basic chemistry
of benzoxazine and earlier work on polyurethane/polybenzoxazine systems are also highlighted.
2.1 Shape memory effect (SME)
The shape memory materials produce shape memory effects by a set of very different mechanisms.
The underlying principles vary greatly from polymers to alloys or ceramics. Accordingly, a thorough
understanding of the principles is needed to achieve the best shape memory properties from a polymer
compound.
In polymers, SME is produced when an initially deformed polymer article returns to its original
shape with the application of an external stimulus which may be heat, light, chemical reagent, humidity, or
radiation [26]. In most studies to date, SMEs have been produced by the application of heat. Three
important factors influence SME in polymer: (1) enabling molecular structure, (2) proper processing, and
(3) punctual programming [11]. Chemical structures of SMPs should be based on at least two phases such
that each phase produces very different responses to the stimulus. If the stimulus is heat, thermal transitions
of these two phases must be widely different such that while one phase undergoes a transition and becomes
easily deformable, the other phase preserves its form and offers durability in shape memory cycle. The
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phase responding to the stimulus is referred to as the reversible phase, while the other phase that preserves
its form is defined as the permanent phase. Processing is another factor that influences SME action. This
includes how deformation is produced, e.g., mode, rate, and temperature. Stretching, bending, and shear are
the most used deformation modes [53]. Of these, uniaxial stretching was found to be the dominant
deformation mechanism on account of its simplicity and sharp impact on crystallization, especially when
melting of crystallites in the reversible phase is used to activate SME. Programming is the combination of
sequence and duration of deformation as well as the intensity and duration of stimulus [54]. In the case of
heat as stimulus, a sample specimen is subjected to heat and its temperature is raised up to the transition
range of the reversible phase. The desired deformation must be produced in a limited time frame to avoid
detrimental effect of stress relaxation. Once deformation is completed, without removing the force, the
specimen is cooled. After complete cooling of the specimen, the force is removed. if the sequence and the
duration are right, the programming steps is successful and the sample is called “trained” [55]. The
actuation of temporary shape and its recovery to the original shape is easy and smooth for the trained
samples.
SME in alloys is completely different from those for shape memory polymers. A structural change
in the atomic order of alloys is responsible for the effect [56]. These alloys are easily deformed at low
temperatures and fixed into new shapes upon removal of the deforming force. Above a specific
temperature, the original shape can be recovered. For all alloys having shape memory effect, a phase
transformation from low temperature phase, e.g., martensitic phase to high temperature phase, e.g.,
austenitic phase is observed, during which deformation that took place earlier in martensitic phase is
recovered. This effect in SMA may be matched with superelasticity or pseudoelasticity, even though it
covers both. They actually refer to high temperature deformations where material is able to change its
shape reversibly upon deformation without fixation. The nuance is that superelasticity or pseudoelasticity
occurs only when the material is in austenitic phase (high temperature phase) and large strains are
recoverable.
Some ceramics are known to demonstrate dimensional change accompanied by crystal structure
alterations under electricity, magnetism, and pressure [57]. Ferroelectricity and piezoelectricity are
produced due to such effects. In ceramics, just like the martensitic phase transformation in alloys, below
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and above a critical temperature, there is a phase change called the ferroelastic phase transition. This
transition can happen in two directions from paraelectric to ferroelectric or from antiferroelectric to
ferroelectric. The induction of transition occurs by heat in the former while it occurs by electric field in the
latter [58].
2.1.1 Parameters for characterization of SMPs
Many simple definitions have been suggested to characterize shape memory materials. The most
frequently used terms are shape or strain fixity, recovery stress, strain or shape recovery, and recovery time.
The definitions of these terms may vary from author to author. In view of this, each term is explained
below, before presenting the definitions adopted in this work.
Recovery stress can be defined as the stress to keep a predeformed, e.g., stretched, specimen at a
fixed shape, undergoing shape recovery during the recovery step as an external stimulus is applied to the
material so as to return to its original shape [11].
In shape memory cycle, the specimen is deformed to a predetermined extent, after it is heated to a
prescribed temperature above the transition temperature of the reversible phase. The extent of
predetermined deformation extent is also the maximum value of deformation. In the case of uniaxial
stretching of a film specimen as depicted in Figure 2- 1 and also in Equation I, the maximum value of
deformation coincides with LS. After deformation, the specimen is rapidly cooled down to below the
transition temperature with deformation force still in place. After cooling, the active deformation force is
removed for specimen to have its final temporary shape that is designated as LD in Figure 2- 1. As the force
is removed, the specimen undergoes an instantaneous relaxation or unconstrained recovery and shrinks to
length LD. This shrinkage brings about a difference in the values of specimen length before and after
unloading of the deforming force. Thus LS is greater than LD. In view of this, shape fixity is defined as the
ratio of the strain before and after deformation force is removed [21], which is given in Equation I. L0 is the
original length as depicted in Figure 2- 1. Shape fixity is also a measure of how well a particular shape
memory material can preserve its own temporary shape in shape memory cycles.
I
10
Recovery ratio or shape recovery ratio of a shape memory material is a measure of the original
shape that can be recovered after the shape memory cycle is completed. Taking the symbols of Figure 2- 1
as reference, equation II defines recovery ratio, where LF is the final length after shape recovery cycle.
II
Accordingly the limiting values of SF and RR are zero and unity. Recovery time refers to the time
period in which an SMP recovers from a temporary shape to its permanent shape [13].
Figure 2- 1 A typical SME cycle for thermally induced SMPs adapted from Ref.21.
A similar graphical explanation was produced by Ratna et al. [13]. Additionally, it has been
suggested that the ratio of glassy plateau modulus to rubbery plateau modulus can be used to represent
shape fixity, since reversibility depends very much on the ratio of these two important properties. Similarly,
the same author has suggested the use of viscous flow strain, e.g., fIR, and strain, e.g., fα, defined below for
the characterization of shape recovery.
⁄ III
⁄
IV
In equations III and IV, Eg is glassy modulus, Er is rubbery modulus, fIR, viscous flow strain, fα is
strain at time is much longer than a reference time t>>tref.
Another concept exists for SMP characterization if specimens are prepared by bending. This is
illustrated in Figure 2- 2. According to this concept, SMP specimen is heated above the activation
temperature and the specimen is deformed with an angular strain θmax, which is the maximum deformation
11
angle. This is followed by rapid cooling below the activation temperature. After cooling, the bending force
is removed. During the removal, there may be an unconstrained relaxation causing the angle to change to
θfixed, which is the angle at temporary shape. In this case, shape fixity is represented as the ratio of (θfixed-
θfinal) to (θmax- θfinal). When the stimulus is applied again for shape recovery to reach the activation
temperature, the sample unwinds to an angle θfinal, which is the final recovery angle. Recovery ratio in this
case can be expressed as the ratio of (θmax-θfinal) to θfinal. Equation V and VI define recovery ratio and shape
fixity respectively.
V
VI
Figure 2- 2 Shape memory behavior characterization with bending test adapted from Ref. 59, 60.
2.1.2 Shape memory effect in alloys
Shape memory effect in alloys is originally based on a diffusionless phase transformation from low
temperature phase (martensite) to high temperature phase (austenite) [56]. This phenomenon occurs due to
the atomic rearrangements. Shape memory alloys like metals are malleable at low temperature and easily
deformed with an aptitude of preserving the temporary shape. At high temperatures, this mechanical
deformation is recoverable almost entirely.
The discovery of shape memory alloys dates back to 1930s. The first alloy that had this property
was Au-Cd obtained by Cheng and Read et al. [61]. But more significant interest appeared after the
discovery of nickel-titanium (Ni-Ti) alloys also called “nitinol” [62,63]. Currently, many other alloys are in
use including In-Ti, Cu-Zn, and ternary alloys, Ni-Ti-Nb, and Ni-Mn-Ga [4,64]. The remarkable properties
12
of these alloys are the precision of transition temperature and the high level of stress that is produced
during the constrained recovery.
Figure 2- 3 A modified schematic of SME mechanism in a typical SMA adapted from Ref.1.
Figure 2- 3 illustrates a single cycle of SMA having shape change due to the application of heat and
deformation force. At low temperature, SMA specimen is in martensitic phase and is precisely deformed.
This deformation is compensated for by heating the specimen above the austenitic phase temperature. Once
the specimen is cooled down, the original martensitic phase is recovered.
The SMAs have found applications in many industries such as biomedical, electronics, and
automobile industries. The SMA product range is very diverse and covers many inventions such as pace
makers, bone fixing staples, guide wires, antenna, micro actuators, and sensors, anti-choking systems
[64,65].
Some advantages of SMA can be listed as follows: large shape recovery, e.g., ~99.9%, very high
elastic modulus, e.g., 80 GPa, high temperature resistance, easy deformation at application temperature,
very high recovery stresses, e.g., 150-400 MPa and very short recovery time, e.g., less than one second
[14]. On the other hand, SMAs suffer from significant drawbacks such as high manufacturing cost and very
low recoverable strains, e.g., usually around 1-2%, and 10% at most [66]. These are valuable incentives for
development of shape memory polymers for those applications requiring shape memory effect but not
cannot use SMAs or ceramics.
13
2.1.3 Shape memory effect in ceramics
Shape memory ceramics (SMCs) are stimuli responsive materials and their SME mechanism
depends on activation through application of electric field. The term for SME in ceramics is referred to as
piezoelectricity or ferroelectricity in which a two directional electricity-strain relationship exists [67].
SMCs offer a transition in mechanical behavior above and below a critical temperature called Curie
temperature, e.g., TC [68]. Three different phases are used to characterize SME in ceramics: paraelectric,
ferroelectric, and antiferroelectric phases [69]. Above TC, SMCs have only one phase which is paraelectric
phase. In this phase, there is no overall directional polarization that can induce a stable change in
mechanical properties, even though reversible spontaneous changes are allowed [70]. In view of this,
paraelectric phase has one way nonlinear relation in electricity-polarity diagram. For a pronounced electric
field induced SME, the ceramics must exhibit ferroelectricity where electricity induced polarization has
different nonlinear routes similar to hysteresis illustrated in Figure 2- 4. The polarization emerging from
electrical current causes charged particles inside the unit crystal to change their locations and to form
charged poles, which in turn create stable but reversible strain on a macro scale.
There are two possibilities that can take place in a stimuli-responsive action of SMCs: a transition
from antiferroelectric to ferroelectric phases or from paraelectric to antiferroelectric phases. The first
transition is an electric field driven transition. However, the second transition generates a recovery from
strained state to the original state, which is thermally driven. This is depicted in Figure 2- 5 where heating
from ferroelectric to paraelectric phase is accompanied by strain recovery and an electric field application
acts reversibly between antiferroelectric and ferroelectric phases. Besides this reversible movement, the
electrical behavior changes from paraelectric to antiferroelectric upon cooling, thus leading SMC to return
to the original shape.
14
Figure 2- 4 Polarization states in different phases of SMC adapted from Ref.71.
Figure 2- 5 SME mechanism of shape memory ceramics (SMC) adapted from Ref.1.
Some of the well-known ceramics demonstrating SME are barium titanate (BaTiO3), which is also
the first piezoelectric ceramic discovered, lead titanate (PbTiO3), lead zirconate titanate (Pb[ZrxTi1−x]O3)
(0<x<1) ( known as PZT and the most frequently used commercially available piezoelectric ceramic),
potassium niobate (KNbO3), sodium tungstate (Na2WO3) [72].
The applications of SMCs are very much diverse and mostly concentrated in the areas of actuators,
sensors, and energy generation from motion. The most famous applications include sonar detectors in
submarines, lighters, and DARPA (defense advanced research projects agency) project, which have focused
on energy generation from motion, transducers, and vibration detectors [73,74].
15
The advantages of SMCs over SMAs and SMPs are that these materials can work using incredibly
small amount of energy compared to SMAs or SMPs and that SMCs are very efficient in small scale, e.g.,
nanometer scale. Additionally, SMCs are built upon relatively inexpensive raw materials and are very
resistant to chemicals and thermal lag effects.
2.1.4 Shape memory effect in polymers
Shape memory effect (SME) in polymers is greatly dependent upon chemical structure of polymers.
However, processing influence and prompt application of processing parameters are as significant as
chemical structure and offer the prospect of outstanding improvements [75]. Molecular design of polymer
chains is decisive such that modulus, elasticity, resistance to chemical environment and thermal lags are
affected by the chemical architecture. On the other hand, thermomechanical history of SMPs is another
factor that is ultimately related to processing and very much interrelated with rheological properties of
SMPs [76]. Still, another aspect is the timing and duration of steps to produce SME. Besides these three
important factors, viscoelasticity of SMPs plays an important role on choosing a correct processing method
and deformation mode [77]. Hence, a review of these factors on SME is considered to be an important task
before presented further detail on SMPs.
2.1.4.1 Viscoelasticity
The role of viscoelasticity in SMPs has been investigated by various research groups [78]. Entropy
elasticity [79], free volume [80], dissipated heat, recovery and relaxation under various conditions, e.g.,
strained or unstrained [81] are major concepts that influence viscoelasticity of polymers and are very
critical for performance of shape memory polymers.
Polymer structure undergoes large mechanical deformation during an ordinary cycle of SME
process. This is depicted in Figure 2- 6. Below the activation temperature, polymer chains are relatively
immobile. With an increase of temperature, especially beyond the transition temperature, the polymer
chains move more freely and macroscopic deformation can be easily produced [11,12]. Since recovery
stress and shape recovery are related to the entropy elasticity, their relation with viscoelasticity is important
to consider at this point.
16
Figure 2- 6 Viscoelasticity of crosslinked polymers with respect to temperature.
In the absence of external forces, polymer chains intertwined with each other without crystallization
or crosslinking will undergo incessant conformational changes at molecular level. Once stress is applied,
strains are produced. Consequently, carbon-carbon bonds rearrange to assume another possible and
favorable conformation with lower entropy. In glassy state, at small strains, polymer specimen can retract
to original state easily, as no plastic deformation occurs. At small strains polymer chains experience a strain
that is reversible and so small as not to harm the conformational arrangement permanently [82]. However,
at large strains, this polymer chains undergo displacement and assume entirely new conformations with
permanent changes at macro level, e.g., plastic deformation. This is especially the case when shape
memory cycles are carried out at longer time scales. Some heat is also generated due to shearing, defined as
viscous dissipation, which can be inferred from the value of loss modulus [83]. The extent of heat
dissipation, if significant, may cause serious loss of elasticity in some cases [84].
2.1.4.2 Molecular architecture
The design of chemical structure is one of the three major factors to attain desirable SME properties.
As mentioned earlier, the presence of at least two independent or related phases is preferred to establish
SME steps properly within a desired thermal activation range. Since SME process relies mostly on
reversible phase to adapt a temporary shape, mechanical and thermal properties of reversible phase within
the desired thermal activation range are of great significance. Unlike temporary shape, permanent shape is
preserved to a high degree by mainly the fixed phase. As anticipated, the thermal and mechanical properties
17
of the fixed phase do not undergo large changes within the thermal treatment range of SME process. Many
blends with/without interaction among the polymers, copolymers, polymers with various fillers, and
complex combinations of these have been successfully studied for their potential shape memory properties
[86-89].
2.1.4.3 Programming: deformation modes and sequence
The programming establishes the most critical conditions of SME process. During shape memory
cycles, many diverse effects influence morphological, thermal, and mechanical properties. These effects are
related to the nature of the polymers. In addition, processing parameters, e.g., deformation conditions,
heating rate, are known to influence the performance. Among these conditions, deformation mode,
deformation rate, temperature and cooling/heating rates are more important. So far, the effects of such
variables were studied only in the context of reversible and fixed phase morphologies, e.g., size and
distribution of domains as well as crystallinity and phase separation [25,90].
In the field of SMP research, only a few deformation modes have been considered, e.g., uniaxial
stretching, bending, compression, and shear [91-93]. These different modes also affect polymer
morphology and accordingly the shape memory properties [94].
In the case of bending as the mode of deformation, only a part of SMP specimen in the vicinity of
bending location experiences critical changes, while the rest of the specimen does not go through the same
deformation. Thus, in bending deformation, the SMP morphology is not affected as much as it is in the
cases of deformation by uniaxial stretching or shear.
Uniaxial elongation has been reported as the most frequently employed deformation mode in the
field of SMP research. First, uniaxial elongation is relatively easy to produce. Second, this mode of
deformation has strong influence on polymer morphology, e.g., reversible phase crystallinity. Many SMPs
have at least one phase that contains some level of crystallinity. This kind of SMPs work well with uniaxial
stretching, since uniaxial stretching produces better strain induced orientation, leading to enhanced
crystallinity. However, the extent of strain induced crystallization may be eliminated or reduced by other
effects such as those generated due to low heating/cooling rates especially as polymer are poor thermal
conductors [21,83,94].
18
2.1.4.4 Influences of processing conditions
The effects of processing conditions on properties of SMPs have been investigated by many
researchers. The conditions that have been investigated so far include heating/cooling rates, deformation
level, e.g., strain, deformation temperature, and deformation rate. The parameters of SME process affected
by processing conditions are recovery stress, shape recovery, and shape fixity [53,95]. A short overview of
these deformation conditions and their influences on SME parameters is presented below.
Heating rate is of great significance due to its influence on relaxation of polymer chains, which
affects the values of maximum recovery stress and shape recovery ratio. It has been reported by a study of
Lagoudas et al. [95]. High heating rate resulted in lower recovery stress maximum this was due to faster
relaxation of polymer chains and lower viscosity at higher temperature. However, an opposite finding was
reported by Cao [53]. In the study of Cao [53], higher heating rate led to higher recovery stress values. It
was attributed to longer relaxation times of the particular polyurethane SMPs used.
Another processing condition is cooling rate of deformed samples. At low cooling rate, sample
needs more time to cool down below the transition temperatures. This provides more time for stress
relaxation as polymer chains experience temperatures higher than the transition temperatures over a longer
period of time. This in turn reduces the recovery stress and shape recovery ratios. Cooling rate also
influences the crystallinity of SMP reversible phase and the onset temperature of the rubbery to glassy
transition. The values of Tg or Tm of the reversible phase are known to depend the cooling rate [26,95,99].
The effects of strain and stretching temperature are remarkable for many SMPs. Strain is accepted as
an important influence on molecular orientation, given that the SMP preserves the strain-induced
orientation produced during deformation. The primary effect of strain is found on recovery stress through
entropic relations, since additional energy storage due to orientation is attained. Stretching temperature is
yet another deformation condition. Generally, the deformation temperature is set at 15 °C or 20 °C above
the transition temperature.
19
2.2 Classification of SME in polymers
The variety of activations by which SME is obtained has led to a set of classifications. Due to large
number of studies on thermally activated SMPs, there has been a preference of the classification based on
thermal activation [4,11-13]. However, a broader view on the classification of SMPs is followed here.
It has been mentioned that SME in polymers can be activated in many ways depending on the nature
of polymers and processing methods. This activation or triggering mechanism was taken as the basis of
further classification below. Consequently, SME was divided into different categories such as light
activated, thermally activated, and electrically/magnetically activated. By this approach, it was aimed at
establishing a classification based on SMEs rather than SMPs. In this context, only the work on shape
memory polyurethanes (SMPU) was taken into consideration, as SMPUs were considered in this thesis.
2.2.1 Thermal activation
The mechanisms of thermal activation essentially depend on the reversible phase in SMPs. The
reversible phase can be amorphous or crystalline. Accordingly, glass transition temperature of amorphous
reversible phase or the crystalline melting temperature of crystalline reversible phase can be used as the
activation temperature [75,82]. Predetermined temperatures above the activation temperatures are set as the
triggering temperature. Any deformation for a temporary shape is obtained at this temperature. Thus, to
initiate activation, the transitions of these phases, e.g., glass transition or melting temperature, must be
known precisely. Crystalline reversible phase may be preferred over amorphous reversible phase, if the
triggering mechanism is desired to work within a narrower range of temperature. However, the influence of
cooling and heating rates [94,99,100] as well as the effect of pressure [100] on crystallization should be
noted, because these deformation conditions can enhance or deteriorate crystallinity. These changes in turn
may lead to contingencies in activation range. If a sharp activation is not a requirement for a specific
application, then an amorphous reversible phase is a promising option. This is due to the fact that the
susceptibility of amorphous reversible phases to cooling/heating rates is not at the same level compared to
crystalline reversible phase. However, an amorphous reversible phase is not entirely independent of the
effects of heating and cooling rates either [93,101-103].
20
The most often used crystalline reversible phases in shape memory polyurethanes are
polycaprolactone diol (PCL) [76,103, 104] and polyethylene glycol (PEG) [29,52,105-107] with the
reversible phase accounted for around 70% of the material. The most preferred amorphous reversible phase
is poly(tetramethylene)glycol (PTMG) [26,96].
2.2.2 Activation by light
Activation of shape memory action by light was reported recently [108,109]. The mechanism of
light activation depends on polymer chains reversibly switching between a crosslinked state and
uncrosslinked state by particular wavelengths of light [109,110]. The application and usage are dependent
on the variety of chemical structures and the methods how the activation by light can bring about a
significant change of modulus [108]. There are two routes to obtain shape memory effect (SME) in light
activated polymers. In the first one, deformation and subsequent crosslinking are needed to retain the
temporary shape [109]. In the second route, deformation accompanies crosslinking [111-113].
Figure 2- 7 SME activation of SMPs by light stimulus, adapted from Ref. 102.
21
Figure 2- 7 illustrates a modified schematic for light activation of SME by crosslinking polymer
chains. Step one depicts the original shape. This is deformed with the application of loading. In step two,
light is introduced to initiate photocrosslinkable points at a particular wave length. After crosslinking and
removal of deformation load, the temporary shape is obtained. The last step consists of breaking the
crosslinks to have the original shape, which is attained by using light source of another wavelength.
The use of light activation to obtain SME offers a very important advantage over thermal activation.
SME by light activation is energy saving, since the method employs a particular wavelength of light rather
than heat. On the other hand, it should be noted that light activation method has important limitations as
well. Since the activation is needed to occur using light, SMP specimen must be both transparent and thin
enough to let a complete transfer of the light before and after attaining the temporary shape. The studies
mentioned above have reported that micro level thickness is very suitable for light activation.
2.2.3 Activation by electric/magnetic field
Polymers are known to be insulator and also non-magnetic [114]. However, with the incorporation
of fillers with magnetism or electric conductivity, these compounds can be rendered partially conductive or
magnetic [115].
Recently, electric field induced SMPs were developed by several research groups [23, 116,117].
Various carbon materials, e.g., carbon black, carbon nanotubes, and carbon nanofiber, were dispersed or
chemically attached to the polymer chains, giving an additional help for activation through the enhanced
properties of fillers [116-118]. Similarly, SMPs filled with micrometer or nanometer level magnetic
particles have been demonstrated to function efficiently [34,116,118]. Even though mechanical and
electrical properties of polymers are enhanced through these fillers, SME enhancement has been the true
goal for SMPs. In this respect, fillers have been used in SMPs that already have an established mechanism
of SME, e.g., thermal activation. For example, SMPs activated thermally are used as composite matrices
and conductive/magnetic fillers constitute dispersed phase in some applications. SME in this case is
improved by providing heat through conductive or magnetic media.
22
2.3 Shape memory polyurethanes (SMPUs)
Shape memory polyurethanes (SMPUs) have a very distinguished place in the field of SMPs
[26,75,82]. They have been used in various systems, e.g., polymer blends [28,123], fillers [89,96], and
composites [59,124]. The following presents an overview of the earlier work of SMPUs, including the
basics of polyurethane chemistry.
2.3.1 Basic information on polyurethanes
The diversity of chemical structures that can be used to obtain polyurethanes is quite large. The two
main raw materials, e.g., polyols and isocyanates, can offer extraordinary structures and functionalities
[19]. Furthermore, small molecule chain extenders, e.g., either an amine or alcohol, can be chemically
inserted between the urethane groups to raise molecular weight of the end product. Polyurethane synthesis
is always carried out in the presence of catalysts which enhance reaction rates of especially chain extension
reactions.
Polyols, relatively high molecular weight alcohol compounds, can have aromatic, e.g. polyester
polyols, or aliphatic, e.g., polyether polyols, chain backbones. The toughness and flexibility of
polyurethanes are mainly derived from polyols. Crystalline or amorphous state of SMPU reversible phase is
also dictated by the nature of polyols. The most used polyols are poly (caprolactone) diol (PCL), poly
(tetramethylene) glycol (PTMG), polysilicates, and various glycerol compounds such as castor oil. Among
these polyols, PCL and some glycerols are crystalline, while PTMG is amorphous. The isocyanates produce
urethane groups upon the reaction with polyols.
The isocyanate compounds may be either aromatic or aliphatic in nature. The urethane linkages
made of isocyanate reactions are capable of establishing hydrogen bonded networks, with positive effects
on mechanical properties and phase separation. In addition, the chemical stability in severe solvents is due
to the presence of isocyanates. The most frequently used aromatic isocyanates are methylene bis (p-phenyl
isocyanates) (MDI) and toluene diisocyanate (TDI). In some cases, aliphatic analogs of these isocyanates
are preferred such as 1,6 hexane diisocyanate (HDI), isophorone diisocyanate (IPDI) and methylene bis(p-
cyclohexyl isocyanate) (H12MDI). Table 2- 1 illustrates chemical structures of common isocyanate and
polyol compounds [128].
23
Table 2- 1 Frequently used reactants for polyurethane synthesis.
N N
C
O
C
O
O
O
OOH
O
O
O
OH
OH
N
C
O
N
C
O
H
O
O
OO
O
O
OH
nn
N
N
C
C
O
O
H
O
O
H
n
C
N
N
C
O
O
H
O
O
O
O
OH
O
n
N
C
N
C
O
O
Si OHO H
n
Many polyurethane products provide good thermal durability for long periods of time at around 70-
100 °C. The chemical and thermal stability of polyurethanes depend on the chemical nature of polyols, e.g.
polyester polyols undergo hydrolytic degradation, while polyether polyols are susceptible to thermal
Castor oil Methylene bis (p-cyclohexyl isocyanate) (H12MDI)
Methylene bis (p-phenyl isocyanate) (MDI)
Isophorone diisocyanates (IPDI)
Hexamethylene diisocyanate (HDI)
Toluene diisocyanate (TDI)
Polycaprolactone diol
Poly (tetramethylene) glycol (PTMG)
Poly (ethylene adipate)
Poly (dimethyl siloxane) PDMS (hydroxy butylated)
24
degradation. In addition the amount of residual catalyst creates problems during the melt processing [128].
For example, the active catalyst initiates reversible reactions under certain circumstances [19].
Polyurethanes undergo dissociation of urethane linkages at elevated temperatures [129].
The microphase separation of polyurethanes originates from the incompatibility between hard and
soft segment domains. These two incompatible blocks offer complementary properties, e.g., rigidity in
aromatic isocyanates and flexibility in polyether polyols. Conventionally, the soft phase constituted of
polyols creates the continuous phase whereas the hard segment domains remain dispersed and act as
physical crosslinking points between the soft segment domains [130].
Microphase separation brings about properties related to blood compatibility and biodegradability
[132,133]. For this reason, polyurethanes have been used in many biomedical applications such as catheters
[134], artificial organs [135], pace maker lead insulations [136], and wound bandages [137].
Although phase separation in polyurethanes offers desirable attitudes, it also produces stress-strain
hysteresis. This term refers to the inability of the polymer to follow the same stress-strain curves during
incessant mechanical cycles. Stress-strain hysteresis in polyurethanes originates from “disruption and
rearrangement of hard segment blocks” [131] and receives consideration in determining shape memory
performance under mechanical cycles.
A review of polyurethane synthesis methods can be found in literature [138]. Polyurethane synthesis
usually involves one of the two common synthesis routes, e.g., one step or two step polymerization [139].
In one step polymerization, polyols, isocyanates and chain extender with catalyst are mixed in the same
media at the same time. In two step polymerization, first a prepolymer which is a low molecular weight
product is obtained from the reaction of isocyanates and polyols. This is followed by the second step in
which the chain extension reaction in the presence of catalyst is carried out. The two step polymerization is
advantageous over one step method due to the additional control on reactions [140].
25
HO
OHn 2n+ x O C N R N C O
O C N R N C
H
O
O O C
O
N R N C O
H
n
x O C N R N C O
n+x HO R' OH
and organotin catalyst
Step II
Step I
O OC
O
N R NC
O
H H
O R' OC
N R N
O
H H
C
O
2n + x
polyol diisocyanate
Prepolymer
Polyurethane
low Mw diol
Figure 2- 8 Two step polymerization of polyurethane adapted from Ref.16.
Figure 2- 8 presents a general schematic of two step polymerization method for synthesis of
polyurethanes. In the first step, both polyol and isocyanate compounds are allowed to react. This is
followed by chain extension in the presence of catalysts.
2.3.2 SMPU blends
A limited number of studies have focused on blends of SMPs [28,123,141] where mostly heat
activated SMPUs are used.
Ebrahimi et al. [123] studied a blend of PCL/PU. The polyurethane was synthesized from poly(ε-
caprolactone) diol soft segment with a number average molecular weight of 2000 g/mol and 4,4’ methylene
bis (phenyl diisocyanate) (MDI) hard segment that was chain extended by 1,4 butanediol. PCL was blended
with molten polyurethane at various weight ratios. It was reported that a blend of PU/PCL (70/30) produced
a trigger temperature around the body temperature and found application as expanding stent implant. The
rationale for using PCL was derived from the fact that PU crystallinity could be easily adjusted using PCL.
26
Another study on SMPU blends was conducted by Jeong et al. [141]. This study considered using
polyurethane/PVC blends to minimize stress-strain hysteresis of PVC during mechanical testing cycles.
The specific polyurethane was synthesized from poly (ε-caprolactone) diol and hexamethylene
diisocyanate. An aromatic chain extender was also used to obtain higher molecular weight PU chains. It
was found that amorphous PVC in the polyurethane led to phase separation in polyurethane. PVC and PCL
domains were found miscible to such an extent that the miscibility influenced the glass transition of PVC
and the melting range of PCL. It was determined that a weight ratio of 8/2 PU/PCL can produce a system
with less strain-strain hysteresis in thermomechanical cycles of SME.
2.3.3 SMPUs with different co-monomers
Although most SMPUs are produced through well known co-monomers, e.g., polyols and
isocyanates, co-monomers such as transisoprene and poly (D,L-lactide)-co-poly(e-caprolactone) diol
(PDLA-co-PCL) have also been used in synthesis of shape memory materials. These are attractive due to
high flexibility and biodegradability.
Ni et al. has studied shape memory properties of transisoprene-urethane copolymer [142]. SMPU was
synthesized from hydroxyl terminated transisoprene which was copolymerized with toluene diisocyanate
(TDI) and chain extended by butanediol. It was determined that the urethane domains formed extensive
hydrogen bonding. Shape fixity of films was found to be almost 100%. However, shape recovery did not
exceed 85%.
In another study, Mather et al. [143] used POSS hybrid monomers as chain extender. The PU
system consisted of MDI based hard segment and poly(D,L-lactide) co poly(ε-caprolactone) diol (PDLA-
co-PCL) based soft segment. It was reported that POSS units segregated and crystallized because of
incompatibility with the soft segment domains. The presence of POSS in PU led to extra crosslinking and
phase separation and produced enhancements in recovery stress and shape recovery ratio. In addition, the
use of PDLA-co-PCL polyols provided a biodegradable soft segment. Note that POSS containing SMPUs
are not beneficial for SMPUs for human body as implant. However, Mather’s study [143] revealed that this
particular SMPU system was biodegradable and offered an SMPU reversible phase with activation
temperatures between 31°C and 45.5 °C.
27
2.3.4 Chemically crosslinked SMPUs
In chemically crosslinked SMPUs, the hard segment domains possess very strong chemical
interactions in the form of permanent crosslinks. A variety of chemically crosslinked SMPUs has been
achieved by using a set of diverse functional crosslinkers [144-146]. The main advantage of chemically
crosslinked SMPUs are higher activation temperatures and higher mechanical strengths. Besides, higher
shape retention and shape recovery are not uncommon.
Hu et al. [147] used a glycerin compound, e.g., dimethylol propionic acid (DMPA), to obtain
chemically crosslinked SMPU with some degree of crystallinity. It was reported that high degree of
crosslinking also led to a reduction in soft segment crystallinity. As a result, the heat of melting and melting
temperature of the soft phase increased in addition to enhancements in mechanical properties. The
transition temperature in this study was around 45-55 ⁰C, which was the crystalline melting range of the
soft phase derived from PCL with a molecular weight of 4000 g/mol.
Another study carried out by Kim et al. [148] utilized surface functionalized silica fillers as the
crosslinking agent in an SMPU system with an amorphous soft segment and isophorone diisocyanate
(IPDI) based hard segment. The crosslinking was obtained through Si-O-Si bridges that occurred after the
sol-gel reactions between the epoxidized polyurethane and functionalized silica fillers. 3-aminopropyl
triethoxysilane (APTES) was used in chain extension from the prepolymers. The crosslinking points were
formed by using triethyl amine (TEA) catalyzed sol–gel reaction between the APTES extended PU hard
domains. The presence of silica was reported to act as both the filling material and crosslinking points,
which increased mechanical properties for example, toughness, strength, and modulus. The long relaxation
time was attributed to higher degree of crosslinking. The materials also provided high shape fixity and
shape recovery ratio.
2.3.5 SMPU nanocomposites
Nanofiller incorporation into polymers has brought a newer horizon to polymer research. Small
quantities of nanofillers have initially been used to increase the mechanical properties. The advances in
functionalization methods allowed several modifications such as surface modifications [149,150], tailored
28
molecule attachments [148], enhanced electrical conductivity and optical arrangements. Of particular
importance to shape memory polymer research are electro-responsive or magnetically active particles.
These properties not only enhance shape memory polymers but also expedite shape memory activation by
resistive heating or magnetic induction heating. These materials are suitable in applications where sensible
heating is not possible, e.g., as implant.
2.3.5.1 Shape memory activation by induction heating
The use of micrometer level magnetic particles embedded in polymer matrices is a well known
concept in the field of smart materials [151]. Many studies on SMA focused on magnetic induction heating
to actuate SME process [152,153]. Another attribute is directional deformation of specimen by application
of magnetic fields. However, several important points must be considered [118]. First, magnetic induction
heating must produce stable activation temperatures. Second, the dispersion of magnetic particle should be
such that the materials are uniform. Third, how the magnetism is produced should be considered, e.g., high-
low frequency or rotational magnetism.
The most critical advantages of shape memory polymers offering magnetic induction heating are
realized in surgical operations in human body, where polymer has to be heated inside the body above the
activation temperature of the reversible phase. Sensible heating by convection may lead to tissue paralysis
and considerably longer times. This becomes a difficult issue if higher activation temperatures than the
body temperature are used. The use of magnetic particles in SMPs offers less activation time and causes
much less damage to the surrounding tissue. Mohr et al. [155] studied magnetically heated SMPs based on
a cyclo-aliphatic biomedical polymer named Tecoflex® EG72D (TFX) filled with as much as 10 wt. %.
magnetic particles. The magnetic particles were iron (II), iron (III) oxides coated by silica.
2.3.5.2 Shape memory activation by resistive heating
Conductive fillers have been used in conjunction with SMPUs to obtain actuation by resistive
heating. In this case, heat is provided from resistive heating of polymer compounds of conductive fillers
[89,116,117]. Paik et al. [23] studied SMPUs filled with carbon nanotube (CNT) and carbon black (CB). It
was found that electric resistance showed positive temperature coefficient for CB between 20 ⁰C and 40
29
⁰C. A fluctuating resistance vs. temperature curve was obtained for CNT around 70 ⁰C. Furthermore,
electrical resistance increased with strain. This indicates that the filler concentration decreased with strain
to values below the percolation threshold, leading to poor conductivity. This finding posed a critical
question as regards to the stability of electrical conductivity especially at large values of strains [89,96].
The positive temperature coefficient (PTC) effect explains the relation between strain and resistivity
for filled systems undergoing large strains [156]. This effect simply refers to a dramatic change of
resistivity with increase of temperature, whereby polymer specimens change from being conductive to an
insulator. This can be interpreted using displacement of filler particles inside the polymer matrices [156-
158]. For shape memory polyurethanes, this issue was addressed by Gunes et al.[157]. In that study, an
SMPU filled with a variety of fillers such as carbon black (CB), carbon nanofiber (CNF) and oxidized
carbon nano fiber (ox-CNF) were investigated. It was found that the reversible phase morphology, e.g.,
crystallinity, was notably affected by the presence of CNF. As a result, this led to an unexpected reduction
in PTC that took place in the same temperature range as melting range of the reversible crystalline phase.
Hence, the PTC effects in SMPU are determined by morphology. Similar observations were made by other
researchers as well [157,159].
Huang et. al [116] studied SMPs filled with 0.5 vol. % Ni and hypothesized that Ni particles served
as conductive connection between CB particles or its agglomerates. These researchers prepared three
different samples. The first sample contained Ni powders aligned by magnetic field. The second sample
contained Ni powders dispersed, but not aligned. The third sample contained the SMP without magnetic
particles. All samples had the same weight percentage of CB. It was observed that electrical conductivity of
Sample I was the highest, followed by Samples II and III. However, repeated application of shape memory
cycles deteriorated the alignment of Ni particles and caused an increase of resistance.
2.3.6 Shape memory alloy (SMA)/shape memory polyurethane (SMPU) composite
There are only a few studies on SMA/SMP composites. The number of studies is low due to the
differences of mechanisms of shape change, e.g., atomic dislocations in SMA vs. thermal transitions of the
reversible phase in SMP. Despite this, Tobushi et al. [159, 160] reported that the differences can be
30
beneficial, if optimum processing conditions and viable applications are found. In those studies, the
composite of SMA wires inserted in SMPU matrix was proposed for an application where specimens bent
at different angles underwent recovery [159]. The transition temperatures of the materials were close, e.g.,
56 °C for SMP and 53 °C for SMA, leading to hardship in control of various steps in shape memory cycle.
It was found out that the maximum recovery stress occurred during actuation above the transition
temperature of SMPU. In addition, shape fixity and shape recovery tests produced very good results.
2.4 Polybenzoxazine
Polybenzoxazine (PB-a) is a developing class of phenolic resins [38]. Even though the resins were
discovered as byproducts of “Mannich” reaction in the last century [39,40], their usages have gained
popularity in the last two decades [41]. Many shortcomings of traditional phenolic resins were overcome,
with the development of polybenzoxazines. These included by-product formation, usage of strong acidic or
basic catalysts, and void formations [161]. Polybenzoxazines can be synthesized in shorter periods and
polymerized thermally employing recent techniques [42,162]. Polymerization of benzoxazine, in contrast to
novolac resins, is associated with near zero shrinkage or a little expansion. Some general attributes of
polybenzoxazines are flame retardancy, very low water absorbtion, and low melt viscosity [50,163]. In
addition, polybenzoxazine has a very extensive polymer network based on hydrogen bonding that leads to
superior mechanical properties such as high glass transition temperature and flexural modulus. Inexpensive
raw materials, short curing periods, high conversion make this class of materials very attractive for
industrial applications which have so far included electronic packaging, flame resistant materials,
adhesives, and electric insulators [46]. However, brittleness of polybenzoxazine is an important drawback
limiting further applications. Many research groups have attempted to eliminate this limitation by either
chemical or physical modifications [38, 42,163].
Below, a brief review of polybenzoxazine with special emphasis on monomer synthesis,
polymerization, and properties is presented.
31
2.4.1 Synthesis of the benzoxazine
In this thesis, benzoxazine refers to a particular member of benzoxazine class, which is based on
3,4-dihydro-3-substituted-2H,1,3-benzoxazines. The monomer unit of benzoxazine consists of a benzene
ring and a heterocyclic ring of oxygen and nitrogen, adjacent to the benzene ring, which is also called the
oxazine ring. In Figure 2- 9 below, the structure of a monofunctional benzoxazine monomer is presented
[38].
O N
R2
R1
Figure 2- 9 General chemical structure of monofunctional 3,4 dehydro-2H-1,3 benzoxazines.
Two frequently used methods of benzoxazine synthesis are given by Burke et al. [40] and Liu et al.
[41]. The former utilizes solvent, while the latter uses bulk method. With the use of dioxane as the solvent,
amine and formaldehyde are mixed together followed by the addition of phenol. This mixture is refluxed 2-
6 hours. The product is cooled and recrystallized from ethanol to obtain purified monomer. The second
method is depicted in Figure 2- 10, a variation of which was also patented by Ishida et al. [42]. In this
method, benzoxazine is synthesized by having a mixture of the following compounds in stoichiometric
ratio a primary amine, an aromatic alcohol and preferably paraformaldehyde. The reaction is carried out at
around 110 °C.
2 CH2O
NH2
NOHHO
OH
N
O
Figure 2- 10 The synthesis of mono functional 3,4-dehydro,2H,1,3-benzoxazine adapted from Ref. 38.
32
In both methods, a variety of phenols and amines can be used to generate an array of benzoxazines
compounds. The solventless method and the advantage of shorter reaction time made Ishida’s method
popular. This alternative method requires no solvent and is carried out in a shorter period with more control
over reaction parameters.
2.4.2 Polymerization
The thermal polymerization of benzoxazine occurs by cationic ring opening of the heterocyclic
oxazine ring which is promoted by high temperature [166,167]. The ring opening reaction of the
benzoxazine was first explained by Burke et al. [40] who reported that the presence of free ortho and para
positions on the phenolic component led to the preferential aminoalkylation at the free ortho position to
form a basic Mannich bridge. This was confirmed by Riess et al. by employing 2,4-di-tert-butylphenol as
catalyst [168]. Recently, Ishida et al. investigated many different chemical structures for their catalysis
effect on benzoxazine polymerization [169]. It was reported that the phenolic compounds having free ortho
positions and aliphatic dicarboxylic acid such as adipic acid can act as catalysts and enhance the rate of
polymerization degree, temperature, and time. The self catalyzing nature of benzoxazine with the phenolic
compounds having free ortho positions was capitalized in a research carried out by Ronda et al. [170]. In
this study, a benzoxazine structure with carboxylic acid functionality was added to usual benzoxazine
monomers to obtain catalytic effects. Results of thermal and kinetic studies revealed that such monomers
improved the degree of conversion, flame retardancy, and thermal stability.
Figure 2- 11 depicts the synthesis scheme of a bifunctional benzoxazine monomer followed by
thermal polymerization. The polymerization starts with thermal activation and C-O bond cleavage on the
heterocyclic ring. This changes the tri substituted ring into a tetra substituted Mannich base bridge through
which polymerization occurs. Subsequent unit addition during polymerization is closely related to the ortho
position of phenolic component and the Mannich bridge [171].
33
HO OH
NH2
CH2On
8<n<100
N
O O
N
2 4
OH
N
OH
N n
m
Figure 2- 11 A schematic of synthesis and polymerization of the benzoxazine adapted from Ref. 38, 163.
2.4.3 Distinguishing properties of polybenzoxazines
Traditional novolac resins and polybenzoxazine have similar polymeric structures, e.g., a highly
crosslinked network of polymer chains. However, polybenzoxazine significantly differs from these resins
with respect to a very extensive network of hydrogen bonding, which is assumed to be the source of better
mechanical properties [41,161,163].
Polybenzoxazine offers many advantages. First, a variety of reactants can be used to design desired
molecular structures [172]. Second, thermal polymerization is easy to carry out and is not dependent on the
use of a strong acidic or basic catalyst. Third, processing is easier compared to traditional resins because of
low viscosity. Fourth, benzoxazine polymerization does not produce any byproduct. Consequently,
benzoxazine can be easily processed to make void free products.
The polymer network structures are responsible for high modulus of polybenzoxazines based on
bisphenol A compounds. Earlier, it was assumed that high flexural modulus and near zero shrinkage (or
little expansion) upon polymerization are due to chemical crosslinking only. However, more recent studies
demonstrated that extensive hydrogen bonding networks also contribute substantially. The intra and
intermolecular hydrogen bonding were found to establish large networks, which in turn prevent shrinkage
and lead to higher modulus [173,174].
34
Thermal properties of polybenzoxazine are the most studied properties besides mechanical
properties. High glass transition temperature, flame retardancy, and high char yield are among the
significant features. Furthermore, it has been reported in detail that polybenzoxazine can be tailored
through two substitutions, preferably phenols and amines, to improve thermal stability. In this regard,
polybenzoxazines containing propargyl, allyl, and maleimide functionalities were developed and
investigated. The studies revealed that the modifications in chemical structure result in polymers stable in
the temperature range between 200 °C and 350 °C [175-177], offering very high glass transition
temperatures from 100 °C to 250 °C [38,163].
2.5 Earlier polybenzoxazine/polyurethane work
Although polyurethane and polybenzoxazine separately enjoy a lot of attention from researchers, the
number of studies on polyurethane/polybenzoxazine blend systems is rather low. These studies were
carried out by four research groups, to the best of our knowledge.
Rimdusit et al. [178-180] studied polyurethane/ polybenzoxazine copolymers and investigated their
thermal stability and mechanical properties. Takeichi et al. [181-183] mainly focused on obtaining films
with enhanced mechanical and thermal properties. Recently, Yeganeh et al. [184,185] investigated the
electrical properties and Wang et al. [186] examined the network structure by spectroscopic techniques.
Rimdusit et al. [178-180] investigated various aspects of polyurethane/polybenzoxazine systems in
three subsequent studies. The first study [178] was to understand the prospect of toughening brittle
polybenzoxazine and to reveal how the other mechanical and thermal properties are affected by the
incorporation of polyurethanes. A bifunctional polybenzoxazine (PB-a) based upon bisphenol A, aniline,
and paraformaldehyde was used in this study. Polyurethanes were prepared as prepolymer (only first step
of the two step synthesis) and consisted of toluene diisocyanate and polyether polyol (Mw~2000). Some
important conclusions are presented as follows. First, glass transition was synergistically improved passing
beyond 200 °C which is above the Tg values of the parent constituents. Second, flexural strength showed an
unexpected behavior for a compound of 90/10 PB-a/PU by weight that produced a high value of 138 MPa.
Third, no influence on thermal degradation temperature was found.
35
The second study [179] particularly concentrated upon the effects of polyol molecular weight used
in polyurethane prepolymers. The experiments were conducted by using the same polybenzoxazine and
similar polyurethanes based on toluene diisocyanate and poly (propylene glycol) with molecular weights
varying from 1000 to 5000 g/mol. The results of the investigation confirmed the synergistic effect on Tg,
while no notable change in the degradation temperature or in char yield amount was detected.
The third study [180] adopted a different approach on PB-a/PU systems. Various isocyanate
compounds including toluene diisocyanate (TDI), isophorone diisocyanate (IPDI), and methylene bis
phenyl diisocyanate (MDI) along with a polyether polyol Mw of 2000 were used in this research.
Furthermore, carbon fiber was added to investigate the prospect of PB-a/PU composites reinforced with
carbon fiber. Results of the characterization tests demonstrated that these composites had more flexural
strength as the PU percentage was raised from 10 wt. % to 40 wt. %. This change in flexural strength was
attributed to the enhanced surface adhesion of polymer matrix with carbon fibers. The surface adhesion was
assumed to originate from polar natures of polyether polyol and carbon fiber, which led to better fiber
wetting and structural integrity. However, no microscopic technique was used to obtain evidence of the
enhancement by examination of the microstructure of the composites.
Takeichi et al. [181-183] carried out successive studies on mechanical and thermal properties of
different PB-a/PU systems. These researches used mono and bifunctional benzoxazines blended with
polyurethane prepolymers based on toluene diisocyanate (TDI) and polyethylene adipate (PEA). The first
study [181] investigated the films of bifunctional benzoxazine/polyurethane blends. The second study [182]
had organophilic montmorillonite (OMMT) clays inside the PU/PB-a system to enhance tensile modulus.
The third study [183] included the efforts of one specific monofunctional benzoxazine with the same
polyurethane and examined the same mechanical and thermal properties.
The study on clay addition into PB-a/PU system was based on the following materials:
monofunctional benzoxazine, e.g., 3-phenyl-3,4-dihydro-2H-1,3-benzoxazine (designated as Pa),
polyurethane prepolymers based on toluene diisocyanate (TDI) and polyethylene adipate (1000), besides
organically modified montmorillonite clay. The clay, benzoxazine, and prepolymer were all consecutively
mixed in dimethyl acetamide (DMAc) and cured under different conditions. It was determined by X-ray
diffraction that only those samples having clay weight percentages less than 5wt. % showed exfoliation. In
36
addition to this, the presence of clay was found to lower the curing temperature of benzoxazine. This effect
in turn helped improve the degree of benzoxazine as well as that of the reaction between free-NCO groups
of prepolymer and OH groups of polybenzoxazine. The reduction in curing temperature was successfully
achieved by as much as 30 °C and was attributed to the catalytic effect of the acidic onium protons. As
expected, an increase in filler concentration caused a decrease in the percentage of elongation at break and
an increase in tensile modulus.
The second study was on the film properties of PB-a/PU systems based on both mono and
bifunctional benzoxazines reacted to the same polyurethane prepolymer [182]. The bifunctional
benzoxazine utilized in the research was bis (3-phenyl-3,4-dehydro-2H-1,3-benzoxazinyl)-isopropane (Ba)
and the mono functional monomer was 3,4-dihyro-3,6-dimethyl-2H-1,3-benzoxazine (Cm). The films were
obtained by solution casting method. Only a single glass transition temperature was observed, which
indicates the formation of uniform microstructure without phase separation. The thermal degradation
temperature of PB-a/PU specimens increased with an increase of Ba content.
Wang et al. [186] investigated the morphology and chemical properties of PB-a/PU systems by
spectroscopic methods. FTIR and SEM/TEM techniques were used to perform the tests. Polyurethane was
based on methylene bis diisocyanate (MDI) and polyethylene glycol polyol (Mw~1000 g/mol). The
prepolymer was chain extended by both butanediol and trimethylol propane. It was observed that
benzoxazine monomer could not be dispersed well in polyurethanes. This in turn led to deterioration of
interactions between PB-a and PU and loss of mechanical properties.
Yeganeh et al. focused on several properties of PU/PB-a systems [184,185]. In these studies, a
bifunctional benzoxazine based on bisphenol-A, methylene dianiline, and paraformaldehyde was used. In
addition, epoxy terminated polyurethane based on hexamethylene diisocyanate and polycaprolactone diol
(CAPA) in varying molecular weights was used. Polyurethane was prepared by end capping the prepolymer
with epoxy to increase crosslinking functionality. PB-a/PU mixtures were synthesized in chloroform.
The first study [184] was particularly designed to reveal the effects of soft phase molecular weight
on electrical properties. The second study [185] aimed at finding out what specific weight percentage of PU
in a PB-a/PU system could produce better properties. It was learnt that incorporation of polyurethanes in
benzoxazine led to systems with lower dielectric constant and dissipation factor compared to pristine
37
polyurethane and enhanced toughness with respect to polybenzoxazine. The values of dielectric constant
(DC) were intimately related to the polarizability of PB-a/PU. The polar groups in B-a, e.g., its polar
aromatic rings, were observed to increase the value of DC. Dissipation factor (DF) is an indication of how
much power is converted into heat as electric current is passed through the material. The DF values were on
the order of 0.1-0.8 x 10-3
. Dielectric strength (DS) values obtained ranged from 24 to 42 depending on the
structure.
38
CHAPTER III
EXPERIMENTAL
This chapter presents detailed information on synthesis of polybenzoxazine and polyurethane
polymers included in this study. In addition, the characterization techniques and sample preparation method
are elaborated.
3.1 Materials
The raw materials used in the study are described below. References to prior literature are presented
where needed to determine the end product and to highlight specific details on the chemical interactions
involved.
3.1.1 Raw materials for synthesis of polyurethane
Diisocyanate: 4,4'-methylene diphenyl diisocyanate (MDI) is commercially available (MDI,
Mondur M) and was obtained from Bayer Material Science (PA,USA) with a molecular weight 250 g/mol
and a melting point of 39 °C. The material was kept in refrigerator, at -10 °C with no contact with water or
light, in a sealed container. Figure 3- 1 illustrates the chemical formula of MDI.
N
C
O
N
C
O
Figure 3- 1 Chemical formula of MDI.
Polyol: Poly tetramethylene glycol (PTMG) (in the form of Terathane® 650) with a molecular
weight of 650 g/mol and melting temperature of nearly 16°C was obtained from Invista (Wichita, KS).
Figure 3- 2 presents general chemical formula of PTMG.
39
HO
O
Hn
Figure 3- 2 Chemical formula of PTMG.
Chain extender: 1, 4-butanediol (BD) with molecular weight of 90 g/mol was purchased from
Avocado (Heysham, Lancs, UK). These small molecular weight diols are used to increase the molecular
weight of polyurethane and to enhance phase separation. Detailed information on chain extenders and
catalysts can be found in Ref. 128-130. Figure 3- 3 demonstrates the chemical formula of 1,4-BD.
HO
OH
Figure 3- 3 Chemical formula of 1,4-butanediol.
Catalyst: Dibutlytin dilaurate (DABCO T120) was obtained in liquid form from by Air Products
Inc. Allentown, PA. The figure below illustrates a general formula of DABCO T120.
Sn
C
C
O
O
O
O
(CH2)3CH3
(CH2)3CH3
Figure 3- 4 General chemical formula for DABCO T120.
3.1.2 Raw materials for benzoxazine
Bisphenol-A: 4,4´-isopropylidenediphenol was purchased from Aldrich as white beads with a
molecular weight of 228.29 g mol−1
. The particular properties of bisphenol-A and other reactants for
benzoxazine synthesis are discussed in Ref.41 and 43.
40
HO OH
Figure 3- 5 Chemical formula of bisphenol-A.
Paraformaldehyde: Polyoxymethylene was obtained as white powder state with a monomer
molecular weight of 30.03 gmol-1
, density of 0.88 g/mL at 25 °C and melting point of 120-170 °C, provided
by Sigma Aldrich. No exact number of repeating unit was indicated.
C OHO
H
H
H
8-100
Figure 3- 6 Chemical formula of paraformaldehyde.
Aniline: Aminobenzene was obtained as brown liquid having a molecular weight of 93.13 gmol-1
and a boiling point of 184 °C, purchased from Sigma Aldrich.
NH2
Figure 3- 7 Chemical formula of aniline.
3.1.3 Preparation of PU/PB-a systems
A new preparation method was followed in this study for synthesis of PU/PB-a system compared to
the earlier methods of bulk polymerization. Polyurethane and benzoxazine monomer were synthesized
according to well established methods [129, 42]. Pristine polyurethane (SMPU) was synthesized by a two
step polymerization method and was used as the control sample, e.g., sample I. This was achieved by
utilizing a three neck reaction bottle for the synthesis of prepolymer and Brabender mixer for chain
41
extension reaction. The specimens, e.g., Sample II and III, contained certain amounts of benzoxazine
precursor and were synthesized by following a series of chemical reactions as explained below.
3.1.3.1 Synthesis of benzoxazine monomer
The benzoxazine monomer used in this study was a bifunctional benzoxazine compound with the
chemical formula; bis (3-phenyl-3,4-dehydro-2H-1,3-benzoxazinyl)-isopropane and is designated as B-a. In
this report, the synthesis of benzoxazine monomer was covered by a US patent [42]. According to this
patent, the monomer can be obtained from a mixture of bisphenol-A, aniline, and paraformaldehyde with
1:2:4 molar ratio. The ingredients were added in a preheated three neck bottle in the following order:
bisphenol-A, aniline, and paraformaldehyde, respectively. The mixture was mechanically stirred at 110 °C
for almost 20 minutes to obtain a yellowish viscous liquid that solidified as a brittle product at room
temperature. The product was pulverized and stored in a cool place.
In this study, B-a was intentionally polymerized up to 10-15% before being used in any other
chemical reaction. This was intended for promotion of the reactions between NCO functional groups of
urethane prepolymer and OH functional groups of polybenzoxazine during mixing of prepolymer and
benzoxazine. This also reduces the time required for thermal polymerization of benzoxazine in film
forming. In addition, the benzoxazine oligomers were found to produce significant catalytic effects on
polymerization of benzoxazine [169-171].
3.1.3.2 Polyurethane prepolymer preparation
Polyurethane was obtained via a two step polymerization method mentioned earlier. PTMG and 1,4-
BD were dried overnight in vacuum oven. The catalyst (DABCO T120) and MDI were used without further
purification or drying. The prepolymer was synthesized in a three neck reaction bottle (500 ml). The bottle
was equipped with a nitrogen purge system and a mechanical stirrer. Nitrogen purge provided dry condition
in the reaction media. Heating was provided through the use of an oil bath in which the reaction bottle was
carefully placed to cover at least three fourth of its body inside oil. Before adding the ingredients, the
42
reaction bottle was heated to 75 °C and the purge gas was turned on. PTMG was poured fist followed by
the addition of MDI. Molar ratio of MDI/ PTMG was 5:1. The reaction was carried out for 2.5 hrs at 75 °C
under nitrogen atmosphere. The prepolymer was kept in the refrigerator for further use.
3.1.3.3 Bulk polymerization of PU and the addition of B-a
Pristine polyurethane, henceforth designated as Sample I, was obtained from chain extension
reaction of prepolymer by 1,4-BD in the presence of the catalyst. This formed the second step of the two
step polymerization process described earlier. The second step was carried out in a traditional mixer, e.g.,
Brabender mixer, C.W. Brabender, Model EPL 7752. The materials were put into the mixing chamber in
the following order: prepolymer, 1,4-BD with catalyst. The catalyst amount was 2-3 drops for 65 cm3 of the
materials, which was around 80 volume % of the mixing chamber. In addition, the mixing torque and
temperature of individual heating blocks in a three block heating system were monitored by a computer.
The values of torque were used to evaluate the mixing quality. The best mixing was achieved after reaching
a plateau in torque vs. time plot. The rotation speed of the rotors was set at 100 rpm.
PU/PB-a systems were synthesized with different ratios of benzoxazine to polyurethane.
Benzoxazine was added to the mixing chamber after prepolymer. A mixture of BD and catalyst was added
after benzoxazine to ensure proper mixing. B-a precursors were inactive at 80 °C in terms of both ring
opening polymerization and the reaction with NCO groups of prepolymer. Therefore, the sequence of
addition was presumed to retain the desired structures of the product. The products, PU/PB-a systems, were
obtained in the form of bulk solid with uniform color and consistency. Figure 3- 8 depicts an example time
line in the Brabender mixing step of the film making process. Note that for the pristine PU, the mixing took
nearly three minutes.
43
0 30 90 180-240Time (sec)
I II III IV
Prepolymeradded Ba 1,4-BD
Solid product taken out
Figure 3- 8 Timeline for sample preparation in Brabender. Step I-IV shows the sequence of addition of the
ingredients.
Table 3- 1 Corresponding molar ratios of raw materials
Sample MDI PTMG BD B-a HS %
I 5 1 4.0 0 71.2
II 5 1 3.5 0.5 73.4
II 5 1 3.0 1.0 75.3
Table 3- 1 presents the molar ratios of the ingredients used in this study. Additionally, all samples
had small amounts of catalyst DABCO T120 to improve and fasten chain extension reaction. In Table 3- 1,
HS % stands for the hard segment percentage in the corresponding sample. In this calculation, the amounts
of MDI, BD and benzoxazine were considered as the hard segment in PU/PB-a system.
3.1.3.4 Film preparation
Bulk samples were dried under vacuum at 40 °C to prevent moisture related side reactions during
compression molding. Films were prepared by using a temperature controlled Wabash Hydraulic
compression molding machine a pressure between 3000 -4000 psi. For the best film properties, the
temperature range and duration of molding employed by earlier researches were examined. Consequently,
the optimum temperature and time values were set to 210 °C and 10 min, respectively [53]. Half a
millimeter thick polyimide sheets kept between 1 mm thick teflon sheets and 2 mm thick steel plates were
used for compression molding of polymer specimen. The mold was preheated to ensure efficient heat
transfer. Bulk sample as chunks was placed between the preheated polyimide films. In the first two minutes
of compression, the pressure was released several times to discharge any gas inside the sample and to
provide a better consistency of the melt. After compression, the assembly was taken to another compression
44
molding machine kept at room temperature. The sample was rapidly quenched by flow of tap water for 10
minutes. The films were found to be homogeneous in color and texture. The thickness was consistent with a
deviation of ± 0.05 mm across film as 30 x 30 cm. Afterward, the films were cured in vacuum oven at 180
°C for 3 hours. A color change was observed from light yellow to darker yellow. In both cases, the
specimens preserved their transparency. These films were stored at room temperature for the
characterization tests.
In Figure 3- 9 a time-temperature line is presented to explain the film making process. The time and
temperature information from each particular step of the process were used to prepare Figure 3- 9. In
Brabender mix up step, raw materials put inside the mixer were heated up to 80 °C to carry out chain
extension reaction. Compression molding was carried out at 210 °C in 10 minutes to covert chunks of
mixed materials into film. The curing in vacuum oven was performed at 180 °C in 3 hours.
Temperature (oC)
Time (min)
210
180
80
25
0 4 0 10 0 180
Brabender
Compression MoldVacuum Oven
Figure 3- 9 The time-temperature plot showing duration and temperature of mixing, compression molding
and vacuum oven.
3.2 Characterization methods
The experimental characterization methods utilized in this study can be divided into five
subcategories; thermal, e.g., differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA),
thermo gravimetric analysis (TGA); mechanical, e.g., tensile testing; structural, e.g., Fourier transform
infrared (FT-IR) and attenuated total reflectance (ATR spectroscopy); morphological, e.g., atomic force
microscopy (AFM), and shape memory property tests that require the use of particular DMA modes and
tensile testing schemes. The following is an overview of these methods.
45
3.2.1 Thermal characterization
Three major thermal characterization methods were used to determine the thermal properties in this
research. Glass transition temperature, crystalline melting temperature, curing temperature and heats of
curing and crystallization were determined by DSC. As thermal transitions occur in a broad range of
temperature, the onset temperature, end temperature and mean temperature were determined for thermal
changes. The thermal scans were carried out under nitrogen atmosphere from room temperature to 275 °C
for pristine polyurethane and to 300 °C for other compounds. Heating rate was 10 °C/min. Aluminum
hermetic pans with an average of 6 mg sample loading were used. A TA Instruments, DSC (Model:
Modulated DSC-2950) was used for all measurements.
Thermal stability of specimens was determined by TA Instruments thermogravimetric analyzer
(TGA) (model 2950) with a heating rate of 20 °C/min and temperature sweep from 25 °C to 800 °C.
Samples weighing around 10 mg within a platinum sample holder were used. Tests were conducted under
air flow and the balance was preserved by constant nitrogen gas flow. The residue and temperatures at 5wt.
% loss were detected through thermal analysis software.
Dynamic Mechanical Analyzer (DMA) was utilized to monitor viscoelastic parameters such as loss
and storage modulus as well as shift angle of the specimens as a function of temperature and time. DMA of
Perkin Elmer Instruments, Pyris Diamond DMA, operated in tensile mode at a frequency of 1 Hz and a
heating rate of 4 °C/min from – 50 °C to 150 °C, was used. Besides DMA in tensile mode, DMA in SS
(stress-strain) mode, e.g., force control and length control, was used to determine shape recovery
percentage and recovery stress. The temperature range was set between 20 °C and 150 °C with a heating
rate of 4 °C/min. In some cases, other heating rates such as 8 °C/min and 16 °C/min were also used.
3.2.2 Characterization of mechanical properties
Mechanical properties including modulus of elasticity, maximum strength, elongation at break, and
toughness were calculated from load vs. displacement data obtained from tensile testing machine. Instron
tensile testing machine (model 5567) having a 1 kN load cell was used for this purpose. The tests were
executed at room temperature with a crosshead speed of 50 mm/min. The operating conditions were as per
ASTM D 882 standard method. Note that this standard is specially designed for films having thicknesses
46
less than 0.4 mm and a supplementary standard of ASTM D 638M. For each compound, five representative
specimens were tested. The results presented in this report include the mean and standard deviation values.
3.2.3 Spectroscopic analysis
Perkin Elmer attenuated total reflectance (ATR) and Fourier transform infrared (FT-IR; Model
16PC) spectroscopy method with a resolution of 4 cm-1
in the range 400−4000 cm-1
were used to carry out
spectral analyses. Functional groups, particularly hydrogen bonded and free urethane carbonyl groups were
followed to investigate hydrogen bonding within urethane domains. Some important peaks assigned to
NCO, C=O (hydrogen bonded or free), and tri-substituted ring (adjacent to benzene ring) of benzoxazine
monomer were detected. Samples run by ATR had a thickness of approximately 0.2 mm.
3.2.4 Analysis of morphology
Atomic force microscopy (AFM) was used to obtain phase and height information from images of
smooth sample surfaces. The domain size and size distribution were obtained by using NanoScope imaging
software. Digital Instruments AFM was used in tapping mode with a scan size of 5 μm and scan rate of 1
Hz. In this mode, the ratio of the set amplitude to the free air vibration amplitude was used to adjust the
force applied onto the sample surface. The phase images taken in this mode gave enough contrast
differentiating the domains of various mechanical properties. Considering that interference of phase angle
difference did not occur, which might have arisen because of adhesion contrast, the soft segment and hard
segment domains were expected to appear as dark and bright spots, respectively in the images.
3.2.5 Characterization of shape memory properties
Films of PU/PB-a polymers as well as control specimen were tested. The tests were carried out
using Instron tensile testing machine equipped with a heating chamber (model 4204) and DMA in stress-
strain (SS) mode. Instron was used for sequential steps of heating, stretching and cooling under tension.
DMA was used to obtain recovery stress in “L control” SS mode. This kept the length of sample fixed
during recovery by applying tensile force, e.g., the recovery force, which was used to calculate recovery
stress. In addition to “L control”, “F control” of SS mode in DMA was used to test shape recovery
47
behavior. This mode provided a fixed zero load on specimens to have unconstrained recovery during
heating. The test conditions were as follows. The duration of preheating before each run and cooling after
each run was set to 5 min. The average thickness of specimens was less than 0.4 mm and the standard
crosshead speed was 50 mm/min. The operating temperatures for each sample are presented in Table 3- 2.
An infrared (IR) thermometer was used to measure actual surface temperature of thin films during tensile
testing to confirm the adequacy of duration for cooling and heating, which revealed very close temperature
values to the set values. Since samples were very thin, it was assumed that the actual surface temperature
represented the overall sample temperature.
Table 3- 2 Testing temperatures for evaluation of shape memory properties
Sample Temperature (°C)
I (control) 71
II 85
II 110
The tests on shape recovery and recovery stress measurements were conducted under the following
conditions. Standard heating rate was 4 °C/min. The sample dimensions were 20x3x0.5 mm. Temperature
scan covered the range between 25 °C-150 °C. For recovery stress, a separate procedure was followed that
took into account the respective transition temperature of the samples and higher heating rates to complete
activation. Table 3- 3 presents the conditions used for measurement of recovery force. It should be noted
that the activation temperature was adjusted at 20 °C above Tg of the compounds.
Table 3- 3 DMA set parameters for application recovery stress measurements.
Sample Initial Temp. (°C) Final Temp. (°C) Heating Rate (°C/min) Duration test (min)
I ~20 71 71 15
II ~20 85 85 15
III ~20 110 110 15
It will be discussed later that the actual heating rate achieved was approximately 25 °C/min,
although the set heating rates were much higher (Table 3- 3). This was due to the limitations of the DMA
machine used in this work.
48
CHAPTER IV
RESULTS AND DISCUSSIONS
4.1 Thermal properties
Differential scanning calorimetry (DSC) was used to determine the thermal properties of PU/PB-a
samples. Glass transition temperatures were also determined from the results of dynamic mechanical
analysis (DMA). In addition, DSC was used to determine the extent of polymerization of benzoxazine.
It was found that benzoxazine monomer synthesized in this research exhibited curing behavior
similar to those observed in the earlier studies [163,169,178]. From DSC curves shown in Figure 4- 1, the
exotherm was determined to have 339 J/g of heat release and a peak for heat release at 225 °C. The heat
release was computed from the area under the curve of curing of the monomer. The heat released from the
second curve was determined as 295 J/g, indicating that this specimen was approximately 13% cured
before undergoing further curing in DSC. In this partially cured specimen, a shift of peak temperature
occurred from 225 °C to 229 °C.
In Figure 4- 2, the DSC curves of Sample I, e.g., the control material, collected after Brabender
mixing and the exposure to heat in vacuum oven are illustrated. None of the curves showed any clear glass
transition behavior. The curve of Brabender product exhibited a melting peak at 219 °C with a smaller
shoulder at around 200 °C. The heat of melting was 30.4 J/g as computed from the area under the curve and
onset temperature of 206 °C. On the other hand, the curve of vacuum oven product exhibited three smaller
peaks and its heat of melting increased to 31.7 J/g and the peak maximum appeared at 223 °C with an onset
temperature at 217 °C.
49
Figure 4- 1 DSC scans of B-a monomer and precursors with 13% curing. Scan rate, 10 °C/min
Figure 4- 2 DSC scans of Sample I. Scan rate, 10 °C/min.
The peak maxima and onset temperatures of thermally aged samples are as follows. The first peak
had 6.8 J/g heat of melting, 175 °C peak temperature, and 165 °C onset temperature. The second peak had
50
10.4 J/g heat of melting, 206 °C peak temperature and 193 °C onset temperature. The third peak had 6.1 J/g
heat of melting, 223 °C peak maximum and 217 °C onset temperature.
The endothermic peak in Sample I can be attributed to the crystallinity of hard domains comprised
of hydrogen bonded urethane groups [53,130]. However, crystalline domains probably reorganized during
thermal aging in vacuum oven that led to a broader size distribution or imperfection among the crystallites.
Figure 4- 3 DSC scans of Sample II. Scan rate, 10 °C/min.
In Figure 4- 3, DSC traces of Sample II are presented. The DSC curve of sample obtained after
mixing in Brabender exhibited three distinct trends. These trends were in the following order: glass
transitions, an endotherm, and an exotherm, respectively. From the endotherm, the melting heat was
determined as 16.45 J/g; The peak temperature of 199 °C and the onset temperature of 184 °C were also
identified. The exothermic heat was calculated as 14.3 J/g with a peak temperature of the exotherm at 261
°C and the onset temperature at 229 °C. As for the glass transition behavior, two clear transitions were
observed. The first transition occurred between 25 °C and 50 °C, while the second transition occurred
between 160 °C and 175 °C.
Sample II- Vacuum oven
Sample II- Brabender
51
The second curve presented in Figure 4- 3 was obtained after keeping the specimen in vacuum oven
at 180 °C for 3 hrs. This curve shows similar trends as the first curve in Figure 4- 3. The melting peak was
observed at the same temperature range, although the exotherm got smaller. The calculated heat for melting
increased up to 24.9 J/g and the peak temperature remained almost the same at 200 °C, while the onset
temperature shifted to 189 °C. The exothermic heat was as small as 1.6 J/g and the peak occurred at 259 °C
with an onset temperature of 241 °C. The glass transition behavior showed some differences compared to
the curve of Sample II after Brabender mix up. The first transition occurred in the same range, e.g., 25-50
°C, whereas the second transition occurred between 55 °C and 75 °C. Comparing the DSC traces in Figure
4- 2 and Figure 4- 3, a few trends can be inferred in regards to the presence of benzoxazine in Sample II.
First, polybenzoxazine with its ability to react with the free isocyanates present in the polymer played a role
in the degree of crystallinity of the hard domains. Recall from Table 3- 1 that molar ratio of NCO
functionality to OH functionality of polybenzoxazine was adjusted to allow for reactions of NCO groups
with OH groups of polybenzoxazine. In view of this, the reaction between NCO groups and OH groups of
polybenzoxazine might have prevented the extent of hard segment crystallinity in Sample II from reaching
to the same level as in Sample I. Second, the exothermic peak appearing right after the melting peak in
Sample II in Figure 4- 3 can be attributed to the residual curing of benzoxazine. Even though
polybenzoxazine-polyurethane reaction via –NCO/–OH reactions, after the polymerization of benzoxazine
might occur in the same temperature range as that of benzoxazine curing, no information was available in
literature about PB-a–PU reaction temperature and specific reaction heat. Nevertheless, confirming with
benzoxazine curing curve in Figure 4- 1, a considerable difference in exothermic heat release did not occur
other than a shift to higher temperatures and a smaller exotherm. This shift can be assumed to have
originated from the presence of prepolymer and smaller amount of benzoxazine in PU/PB-a. In addition, it
should be emphasized that the percentage of curing was computed based on exothermic heat release from
only benzoxazine curing. Thus, the heat of the exotherm from Figure 4- 3 was computed and compared
with 339 J/g for 100% curing of benzoxazine. It was observed that benzoxazine in Sample II after curing in
vacuum oven was 95%.
Glass transition behavior of Sample II on both curves in Figure 4- 3 appeared to be similar, e.g., two
transition ranges at low and high temperatures. This was thought to be due to phase separation. Although
52
earlier work on PU/PB-a mostly reported the formation of single phase systems, the current study is unique
in that the formulations used in the preparation of compounds made the formation of multiple phase
systems possible. First, Table 3- 1 revealed clearly that the polymer system was primarily polyurethane
with small amounts of benzoxazine. Since pristine polyurethane underwent phase separation, this must
have affected the overall phase separation in Sample II regardless of the OH-NCO reactions after curing of
benzoxazine. Second, the product of polybenzoxazine-polyurethane reaction might have formed a separate
phase different from dominant polyurethane phases. The shift of glass transition ranges can be attributed to
glass transition of polyurethane soft segment at lower temperatures, e g., 25-50 °C, and the new hard
segment phase originated from polybenzoxazine. Such increases of glass transition temperatures have also
been reported in previous PB-a/PU systems [178-181]. Materials in this study differed from those of earlier
studies in that nearly all of the earlier work considered mixing and curing of prepolymer and benzoxazine.
This kind of preparation ultimately led to proportionate reactions between the two reactants, e.g.,
prepolymer and polybenzoxazine. On the other hand, prepolymer and benzoxazine were mixed at
predetermined molar ratios in this research. In this manner, a major portion of prepolymer underwent chain
extension reactions and formed the usual SMPU domains, while the rest of the chains reacted with
polymerized benzoxazine precursors, e.g., polybenzoxazine. As a result, a reduction of phase separation
within polyurethane hard domains should not be counted towards within domains formed from the reaction
between –NCO groups of PU chains with –OH groups formed in polybenzoxazine.
DCS traces of Sample III are presented in Figure 4- 4. Note that in this case a higher fraction of
benzoxazine was used as indicated in Table 3- 1. The curve belonging to the specimen obtained after
Brabender step exhibited a melting peak with an endothermic heat of 8 J/g, peak temperature of 188 °C,
and an onset temperature of 175 °C. The exothermic curve had a 33 J/g heat release, 260 °C peak
temperature, and a 235 °C onset temperature. The first glass transition occurred in the same range as that of
Sample II, while the second glass transition shifted to a lower temperature between 140 °C and 160 °C.
Sample III after vacuum oven showed a significant difference. In this case, no melting peak appeared, but
the exotherm was observed. In addition to this, a smaller exotherm occurred right before the usual
exotherm of benzoxazine curing. The heat release for this smaller peak was as little as 0.65 J/g. Peak
temperature of this exotherm was 185 °C and the onset temperature was 166 °C. The usual exotherm
53
produced 16.8 J/g heat release with a peak temperature of 257 °C and the onset temperature of 215 °C.
Lastly, the curing percentage for Sample III was 73%.
Figure 4- 4 DSC scans of Sample III. Scan rate, 10 °C/min.
The disappearance of the melting peak of crystalline urethane domains and the appearance of a new
second exotherm were the characteristics of Sample III. Taking the molar ratio of the components into
account, the DSC traces, especially the abundance of hard segment melting peak can be explained as
follows. It is evident that the polyurethane portion of Sample III did not have any semi crystalline domains
and this led to the absence of usual melting peak around 200 °C. In addition, X-ray results obtained from
both small angle X-ray diffraction (SAXD) and wide angle X-ray diffraction (WAXD) measurements were
used to support the findings of DSC analyses. Figure 4- 5 shows the images of WAXD and SAXD
measurements results. In the first image, the presence of only a diffuse ring is the result of the amorphous
morphology. Similarly, SAXD measurements exhibited a diffuse scattering indicating the absence of
crystallinity. This can be ascribed to the increased amount of benzoxazine interfering with the organization
of hard segment urethane groups by hydrogen bonding and consequently a reduction of the crystallinity,
due to more prevalent polyurethane-polybenzoxazine reaction. The molar ratio of benzoxazine in Sample
54
III was twice as much as it was in Sample II. While a higher portion of polybenzoxazine allowed more
reactions with polyurethane, due to incomplete curing, more benzoxazine oligomer was also formed and
led to different physical interactions with urethane domains. This will be discussed again in relation to
hydrogen bonding investigation by FT-IR/ATR.
Figure 4- 5 WAXD and SAXD measurement results of Sample III, respectively the first and the second
images.
More accurate values of Tg were obtained by conducting DSC measurements with higher heating
rates in the same temperature range of earlier tests. These results are presented in Table 4- 1. It was found
that as the benzoxazine amount in the samples was increased, the onset value of 1st Tg shifted to higher
temperatures, while that of 2nd
Tg decreased to lower values. It is known from earlier polyurethane
researches that an increase in the Tg value of soft segment is possible with higher contents of hard segment
[140]. In this regard, the reaction between PU and PB-a might have led to the same effect on the Tg values
of soft segment, since the product of this reaction was a part of the fixed phase in PU/PB-a system. On the
other hand, the fall in the 2nd
Tg values should be attributed to partial phase mixing due to the chemical
interactions between PU and PB-a, which was able to affect the miscibility of usual hard and soft segment
of polyurethane.
Table 4- 1 Glass transition onset values determined through DSC with higher heating rates.
Samples 1st
Tg (°C) 2nd
Tg (°C)
I 44 106
II 51 97
III 62 96
55
4.2 Thermomechanical properties
Viscoelastic properties of the samples were investigated by DMA in tension mode. Several
parameters such as loss tangent, storage modulus, and loss modulus were obtained as a function of
temperature.
Figure 4- 6 Loss tangent (tan δ) as a function of temperature for all samples. Heating rate, 4 °C/min;
frequency, 1 Hz.; scan rate between, -50 °C and 150 °C.
Figure 4- 6 presents shift angle, e.g., tan δ, as a function of temperature of the samples. The
temperature at the maximum value of tan δ was used to determine glass transition temperatures of the
samples. It is evident that samples exhibited only a single peak of tan δ, indicating a single value of Tg.
This was assumed to be due to the following reason. The phase separation degrees of PU-PB-a samples
were between 13% and 27% as determined from FTIR/ATR investigation, which might be hard to detect in
DMA, since the measurement was based on macro scale properties, e.g. extension response of films to heat
changes. Consequently, single Tg values were accepted as representative Tg values in this study due to the
focus on such macro properties as shape recovery and recovery stress. The Tg values inferred from the
maximum of tan δ vs. temperature plots of Sample I and Sample III were 51 °C and 65 °C, respectively,
56
whereas that of Sample III was much higher, 91 °C. Also note from Figure 4- 6 that Samples I and II
exhibited wide transitions between glassy and rubbery states within 100 °C. On the other hand, Sample III
had its transition within a larger angle range.
Figure 4- 7 Storage modulus (E') as a function of temperature for all samples. Heating rate, 4 °C/min;
frequency, 1 Hz.; scan rate between, -50 °C and 150 °C.
Figure 4- 7 presents the values of storage modulus as a function of temperature. The plateau values
of storage modulus (E') in glassy state were determined as 7.5 GPa, 5.3 GPa, and 4.7 GPa, respectively for
Sample I, II and III. The curves exhibited behavioral similarity to crosslinked polymers in the rubbery
region. It is seen that between 100 °C and 150 °C, Sample I and II led to higher rubbery plateau modulus
than Sample III. Figure 4- 8 shows loss modulus (E'') as a function of temperature for all of the Samples. It
is seen that Sample III flowed much more easily than the other two samples at high temperatures.
57
Figure 4- 8 Loss modulus (E'') as a function of temperature for each sample. Heating rate, 4 °C/min;
frequency, 1 Hz.; scan rate between, -50 °C and 150 °C.
Table 4- 2 reveals the values of storage and loss modulus changes in a specific temperature range
around the Tg of corresponding sample. It is seen that there is a very large gradient of storage modulus
values associated with change of temperature around glass transition. For Sample I, the ratio of storage
modulus values 20 °C below and above its Tg is 11. But, this number for Sample II and III drastically
changed to 147 and 271, respectively.
The ratio of E' values at 20 °C below Tg and 20 °C above Tg play critical role in determining the
shape memory properties. In view of this large and rapid temperature changes during a single cycle of
SME, SMPs must have proper modulus changes in the activation temperature range to allow for rapid
mechanical deformations [12, 22]. It should be noted that large values of E' (Tg-20 °C)/ E' (Tg+20 °C)
mean higher values of shape fixity and rapid shape recovery. In this context, Sample II and III exhibit much
better shape memory properties compared to Sample I. As a result, this was manipulated at the outset
before undertaking this project.
58
Table 4- 2 Viscoelastic properties of the samples
Samples E' (glassy) (Pa) E' (Troom) (Pa) E'(Tg+20) (Pa) E'(Tg-20)/E'(Tg+20) Tg (°C)
I 7.5x109 4.1x10
9 5.3x10
8 11.3 51
II 5.3x109 3.8x10
9 3.4x10
7 147 65
III 4.7x109 3.9x10
9 1.4x10
7 271 91
Figure 4- 9 The effect of curing time on Tg and E' of Sample II and III.
The duration of thermal treatment in vacuum oven was varied to investigate its effects on storage
modulus and glass transition values that reveal important information about how these samples would
thermally change during the thermomechanical cycles of shape memory process. The results are presented
in Figure 4- 9. It was determined that the Tg values of Sample II and III increased by 8 °C and 11 °C,
respectively, whereas the storage modulus values showed a slight overall fall. Sample II exhibited
decreased from 4.5 x108 Pa to 1.4 x10
8 Pa and Sample II exhibited a decrease from 1.4 x10
7 Pa to 1.2 x10
7
Pa. It should be noted that the changes in the values were assumed to occur due to the polymerization of
partially polymerized residues of benzoxazine precursors.
1.E+6
1.E+7
1.E+8
1.E+9
50
60
70
80
90
100
110
0 12 24 36 48 60
Sto
rag
e M
od
ulu
s (E
')
Te
mp
era
ture
(°C
)
Curing Time (hrs)
Sample II-Tg
Sample III-Tg
Sample II- E
Sample III- E
59
Figure 4- 10 shows thermogravimetric analysis (TGA) results of sample specimens subjected to air
at high temperatures. It is observed that all these samples underwent loss of weight and exhibited three
transitions. The transitions can be divided as follows. The first transition was between 225 °C and 350 °C,
the second transition between 350 °C and 450 °C, and the third transition occurred between 450 °C and 650
°C. The temperature at 5 wt.% loss, % residue, and temperature of transitions are presented in Table 4- 3.
Figure 4- 10 TGA curves of the samples. Scan rate was 20 °C /min.
Table 4- 3 TGA analysis results. Scan rate was 20 °C/min.
Sample 5wt. % loss
(°C) Residue (wt. %)
1st
Transition
(°C)
2nd
Transition
(°C)
3rd
Transition
(°C)
I 302 0.085 300-350 350-450 450-650
II 290 0.488 255-375 400-450 550-650
III 267 0.792 225-350 400-425 575-650
It was concluded from TGA analysis results that the earlier initiation of weight loss for Samples II
and III ranging from 225 °C to 375 °C was due to the susceptibility of unpolymerized entities inside the
polymer systems to degradation. This is demonstrated within the circles of individual transitions in Figure
1232 cm-1
1st Transition
2nd
Transition
3rd
Transition
60
4- 10. From Sample I to III, the slope of the first transition got larger. However, the weight loss percentage
got smaller compared to that of Sample I, which can be accepted as an indication of enhanced thermal
stability due to polybenzoxazine presence. The second and the third transitions exhibited a difference from
the first transition in terms of the temperature values at maximum weight loss within their limits. Sample II
and III exhibited lower weight loss percentages in the second transition and close loss values in the third
transition compared to Sample I. However, these weight losses always occurred at higher temperatures with
Sample II and Sample III.
4.2 Spectroscopic analysis
The spectroscopic analyses of the chemical substituents in the samples, the prepolymer and
benzoxazine monomer were carried out by FT-IR and ATR measurements. It was first confirmed that the
spectroscopy data of the prepolymer and benzoxazine monomer matched the characteristics reported by
other researchers [161-3]. Figure 4- 11 shows FT-IR spectra of benzoxazine monomer. Two characteristic
absorption bands at 942 cm-1
and 1232 cm-1
were detected. These peaks were previously assigned to the
benzene mode of the benzene ring next to the oxazine ring and trisubstituted benzene of the oxazine ring, in
that order [176, 183]. The exact structure of chemical formula of benzoxazine was given earlier in Figure 2-
11 and may be helpful for understanding the assignments of the characteristic peaks
.
Figure 4- 11 ATR spectra of benzoxazine monomer.
61
Figure 4- 12 1H-NMR analysis of benzoxazine monomer (bis (3-phenyl-3,4-dehydro-2H-1,3-
benzoxazinyl)-isopropane).
The chemical nature of benzoxazine was also confirmed by 1H-NMR that is presented in Figure 4-
12. The results are in good agreement with an earlier study that investigated the exact benzoxazine
monomer used in this research [169]. The first single peak at around 1.5 ppm was due to two methyl groups
that keep the two main benzene rings together. The other two single peaks at 4.5 ppm and 5.2 ppm are due
to the methyl groups in the oxazine ring, before and after nitrogen atom. The reason why their magnitudes
are nearly half of the other single peak on the right side is because of the number of methyl groups that
contributes to the same peak. The smaller peaks at around 7 ppm are from the aromatic rings. The reference
peak was overlapped due to the same detection range as the aromatics.
62
Figure 4- 13 FTIR spectra of polyurethane prepolymer.
Figure 4- 13 presents FT-IR spectra of polyurethane prepolymer based on MDI/PTMG. The spectra
had the following characteristic peaks; 3420 cm-1
, (N-H bonded); 3320 cm-1
, (N-H free); 2281 cm-1
,
(NCO); 1733 cm-1
, (CO free); 1703 cm-1
, (CO bonded); 1600 cm-1
, (C=C); 1530 cm-1
, (δN-H and νC-H)
and 1080 cm-1
, (C-O-C). Among the peaks illustrated in Figure 4- 13, the peak appearing at 2218 cm-1
that
belongs to NCO functional groups was used to monitor the chain extension reaction with butanediol (1,4-
BD). During chain extension reaction, the NCO peak diminishes, due to the consumption of free isocyanate
(NCO) groups. The peaks at 1733 cm-1
(free CO) and 1703 cm-1
(hydrogen bonded CO) were used to
calculate the hydrogen bonding percentages in the hard segment. This was used to infer the degree of phase
separation.
Figure 4- 14 exhibits the spectroscopy data from ATR measurements. All samples had similar
absorption peaks, also present in the prepolymer except for the NCO peak at 2281 cm-1
. The absorptions at
3420 cm-1
and 3320 cm-1
reduced to smaller magnitudes. Due to the very complex nature of
polybenzoxazine absorptions or coincidences of polybenzoxazine and polyurethane peaks, many of the
peaks from polybenzoxazine could not be detected. However, the absence of absorption peak at around
3500 cm-1
, which is assigned to hydroxyl groups of polybenzoxazine, was a clear signal that hydroxyl
functionalities of polybenzoxazine completely reacted with the free NCO groups of the prepolymer.
63
Figure 4- 14 ATR spectra of the samples.
Table 4- 4 presents the calculated values of bonding index of the samples. The areas of hydrogen
bonded CO peak at 1703 cm-1
and free CO absorption peak at 1733 cm-1
were used to determine the value
of hydrogen bonding index. The results indicate that the hydrogen bonding reduced in the presence of
benzoxazine. It may be attributed to the interactions of urethane linkages with both polybenzoxazine and
benzoxazine precursors including monomers, dimers, and oligomers. The data in Table 4- 4 also supports
an earlier observation that crystallinity reduced in samples containing benzoxazine due to lower degree of
hydrogen bonding between hard segment domains. Another attribute of hydrogen bonding was the degree
of phase separation. The values of α in Table 4- 4 also indicate of the degree of phase separation between
PU segments. It is noted that, the phase separation degree decreased to half of its original value going from
Sample I to Sample III. At this point however, it should be pointed out that the discussion so far was based
on PU hard domains, which could not reveal any information about the overall phase separation arising
from chemical interactions between PU and PB-a. In this regard, DMA results, which quantitatively
revealed thermomechanical properties of PU/PB-a, should be taken into account, for SME considerations.
1733 m-1
1703 cm-1
Sample II
Sample III
Sample I
64
Table 4- 4 Hydrogen bonding % of urethane domains. AFCO refers to 1730 cm-1
and ANCO; 1700 cm-1
Sample AFCO AHCO α=AHCO/ (AFCO+AHCO)
I 14.1 5.1 0.27
II 10.7 2.58 0.19
III 7.97 1.2 0.13
Figure 4- 15 illustrates the curve fitting worked out by FYTIK free plotting program [187]. These
fitted curves were used to calculate the area under the carbonyl peaks reported in Table 4- 4.
Figure 4- 15 Curve fitting at 1733 cm-1
and 1703 cm-1
for Sample I, II, and III, respectively.
4.3 Mechanical properties
Mechanical properties of polymers are of great significance to the applications. Material
performance and failure modes are dependent upon mechanical properties such as modulus of elasticity,
elongation at break, yield stress, ultimate stress and toughness values. Characterizations of these properties
are carried out by tensile tests with respect to the applicable standards, e.g., ASTM or ISO. For this study,
ASTM D 882 method was used.
Table 4- 5 Tensile properties of the samples.
Sample Young’s Modulus
(MPa)
Elongation %
at Break
Yield Stress
(MPa)
Ultimate Stress
(MPa)
Toughness
(MPa)
I 228 478 41 50 216
II 258 95 44 35 36
III 475 32 N/A 160 6
65
Figure 4- 16 Stress-strain diagram of the samples at room temperature.
Table 4- 5 lists important mechanical properties. It is seen that the mechanical behaviors of the
samples were greatly influenced by the incorporation of polybenzoxazine. Modulus enhancement is
obvious for Sample II and III compared to Sample I. Young’s modulus for Sample II and III were
determined as 258 MPa and 475 MPa, respectively. Elongation at break, however, reduced significantly in
the presence of benzoxazine. Sample III possessed the highest ultimate stress, although Sample II had a
smaller ultimate stress than Sample I. While Sample I and II demonstrated yielding behavior, Sample III
showed brittle failure. Toughness values consistently decreased from Sample I to Sample III. Even though
there were strong chemical interactions between polyurethane and polybenzoxazine, the toughness values
of samples having benzoxazine, e.g., Sample II and III, did not go above 36 MPa (Sample II). Toughness
value, e.g., 6 MPa for Sample III, is still higher than those reported earlier with similar benzoxazine weight
percentages and polyurethane structures [178].
Figure 4- 16 presents typical stress-strain diagrams. The tests were conducted at room temperature
at ~20 °C with dog-bone shaped specimens according to ASTM D-882 standard for thin films, e.g., for
specimens with thicknesses less than 0.4 mm. The crosshead speed was set to 50 mm/min. It is seen that the
failure behaviors of Sample I and II differed from that of Sample III. The formers exhibited yielding
behavior, while the latter revealed brittle breakage. This important distinction should be attributed to the
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absence of crystallinity in Sample III. In addition, oligomeric benzoxazine precursors, even though present
at low concentrations, might have influences on the failure behavior of Sample III. Note that benzoxazine
monomer and its precursors are brittle.
Figure 4- 17 Stress-strain diagram of the samples at deformation temperatures. Sample I, 71 °C; Sample II,
85 °C; Sample III, 110 °C. Max strain was 150%.
The tensile properties were also investigated at the deformation temperatures of each sample. The
computation results are presented in Figure 4- 17. Measurements produced reasonable results in correlation
with the data of DMA storage modulus. The samples exhibited a decreasing trend of modulus of elasticity
from Sample I to Sample III. The particular thermomechanical features of the samples and the deformation
temperatures were two main reasons of the current results. Noting that in Figure 4- 7 rubbery plateau of
Sample III was at the bottom with the lowest value above 110 °C, the current results reproduced the
correlation between tensile and thermomechanical properties.
4.4 Shape memory properties
The important parameters such as recovery stress, shape recovery, and shape fixity are evaluated to
determine shape memory performance of the polymers. In this study, these parameters were investigated as
function of deformation conditions such as stretching rate and strain rate.
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4.4.1 Recovery stress and shape recovery ratio
Figure 4- 18 illustrates how the recovery stress evolved for 100% stretched samples as function of
temperature from 20 °C to 150 °C with a heating rate of 4 °C/min. The samples were stretched with a
crosshead speed of 50 mm/min, at 71 °C, 85 °C, and 110 °C, respectively for Sample I, II and III.
It is observed that Sample I produced the lowest recovery stress of around 6.8 MPa. Note that the
maximum values of stress for each sample in Figure 4- 18 were taken as the recovery stress. The maximum
value occurred at 86 °C, which is 35 °C more than the Tg value of Sample I. This finding is in good
agreement with an earlier work [149]. As evident, the recovery stress values for Samples II and III were
much higher than that of Sample I, with values respectively at 11.2 and 13 MPa. Sample II reached its
maximum recovery stress value at a temperature very close to that of Sample III, around 94 °C. In addition,
the curves in Figure 4- 18 illustrate some differences in terms of the trends. While the recovery stress value
of sample I could not hold on to its maximum value and declined at temperatures higher than 86 °C,
Samples II and III seemed to keep the trend more firmly even well above their Tg values. Incidentally, the
recovery stress curve remained almost flat at temperatures higher than 100 °C, up to 150 °C.
Figure 4- 18 Recovery stress behaviors of 100% strained samples. Heating rate was 4 °C/min and stretching
rate was 50 mm/min.
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It is also evident from Figure 4- 18 that regardless of the extent of the curing of benzoxazine
precursors and the extent of reaction between of PB-a OH groups and PU NCO groups, in Sample II and
Sample III, benzoxazine addition to SMPU led to substantial increase of recovery stress. The reason behind
this improvement is the enhanced ability of the samples to preserve elastically stored entropic energy
possibly by the significant hard domain structures. The hard domains possibly also underwent changes in
structure, size and distribution due to the reactions between OH of PB-a and NCO of PU. Furthermore, it
was considered that a different hard domain or sub-domain apart from original polyurethane (PU) hard
domains might have emerged because of the same interactions. For now, this part of the discussion is
deferred until the morphological properties are discussed.
Figure 4- 19 Shape recovery ratio of 100% strained samples. Heating rate was 4 °C/min and stretching rate
was 50 mm/min.
Figure 4- 19 depicts the shape recovery ratio of the samples. In this context, a shape recovery ratio
of 100% strained specimens was preferred. Sample I started to recovery at 40 °C and the slope generally
decreased. The specimens of Sample I recovered only 72% of their length at up to 150 °C. Sample II
showed a sharper recovery behavior starting at around 60 °C. Note that the sample had higher Tg at 65 °C.
After 110 °C, the curve adopted a plateau shape with little increase in recovery ratio. The maximum value
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69
of recovery ratio at 150 °C was 84%. Shape recovery of Sample III started at around 65 °C and with a very
steep climb reached 93% at 100 °C. After 100 °C, a clear plateau behavior was observed.
Shape recovery behavior observed in Figure 4- 19 was reasonable, considering the transition and
heating rate into account. Sample II and III had greater modulus ratio of E'(Tg-20)/E'(Tg+20), which is
reflected into stronger shape recovery performance.
4.4.2 Effects of deformation conditions and heating rate
Several deformation conditions were found to exert strong influence on recovery stress and shape
recovery [13,17]. Particular attention was given upon heating rate, stretching rate, strain, stretching
temperature, and in some cases cooling rates [26,188, 189]. In this study, all of these conditions, other than
cooling rate, were investigated. The tests were carried out with Sample II considering that other samples
would exhibit similar trends.
4.4.2.1 Influence of strain on recovery stress and shape recovery
The influence of strain shape memory properties was addressed in several studies [53, 73]. It was
observed that shape recovery stress also increased with the strain applied to shape memory polymers. This
improvement was attributed to strain induced orientation of polymer chains. In particular for SMPUs,
orientation of polymer chains in different segments was found to be different. Hence, their behaviors under
the same strain conditions were not the same. This difference originated from the hard segment orientation,
which in turn is very much related to the orientation of urethane domains. The urethane domains orient
differently in various directions at different levels of strain, e.g., transverse orientation at low strains and
orientation in the direction of elongation at high strains. However, soft segment orientation occurs always
in the direction of elongation. In addition, the miscibility between the segments was also found to alter
orientation at different temperatures. As a general observation, the orientation improves at high
temperatures, however, especially at temperatures higher than the glass transition or melting temperatures
of hard domains, orientation declines due to diminished hydrogen bonding and improved miscibility of the
segments.
70
The influence of strain on shape memory properties are presented in Figure 4- 20, Figure 4- 21and
Figure 4- 22, respectively for Sample I, II, and III. For Sample I and II, strain increase also resulted in an
increase of recovery stress. The largest increase occurred in sample I, e.g., from nearly 4 MPa (50%), to
around 8 MPa (100%), and then to 14 MPa (150%). The curves have very similar shapes. A very sharp
increase occurred in nearly 90 seconds followed by a plateau. After reaching the plateau, the specimens of
Sample I did not show any relaxation. These results are in agreement with an earlier work that had the same
formulation and similar strain ratios, e g., 1.6, 2.0, and 2.7. The approximate recovery stress maximum
values obtained were respectively 9.5 MPa, 11 MPa, and 17 MPa [53].
Figure 4- 20 Strain effect on recovery stress trend of Sample I. Stretching rate was 50 mm/min and
stretching temperature was 71 °C.
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Figure 4- 21 Strain effect on recovery stress trend of Sample II. Stretching rate was 50 mm/min and
stretching temperature was 85 °C.
Sample II exhibited higher recovery stress values for 50% and 100% strained specimens, 6 MPa and
8 MPa, respectively. However, 100% strained specimen demonstrated a different value of recovery stress,
10 MPa, which is much lower than Sample I. The specimens strained at 100% and 150% also showed some
degree of relaxation after reaching the peaks values. Another interesting point to note is how rapidly the
maximum value of recovery stress was reached.
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Figure 4- 22 Strain effect on recovery stress trend of Sample III. Stretching rate was 50 mm/min and
stretching temperature was 110 °C.
In Figure 4- 22, the effects strain on recovery stress of Sample III specimens are depicted as a
function time and temperature. The 50% strained specimen produced a peak maximum at around 6.5 MPa
and the 100% strained specimen exhibited a higher value of about 8.3 MPa, whereas the 150% strained
specimen had an unexpected low maximum value with 7.7 MPa.
The recovery process is a ramification of a variety of thermomechanical changes. During the course
of returning back to the initial shape, it is imperative that polymer chains know their original location and
conformations inside the specimen. This occurs only when the polymer chains are not deformed beyond a
critical level, at which the chains slip pass each other and adopt new permanent conformations. In addition,
the internal factors such as heat dissipation and stress relaxation do not reduce the shape recovery
dramatically.
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Figure 4- 23 Strain effect on shape recovery behavior of Sample I. Stretching rate was 50 mm/min and
stretching temperature was 71 °C. Heating rate was 4 °C /min.
Shape recovery values of Sample I specimens are presented in Figure 4- 23. A trend is clearly seen
regarding to the influence of strain on recovery ratio. While 50% strained specimen attained 90% shape
recovery, 100% strained specimen had a maximum shape recovery of 72%. On the other hand, the 150%
strained specimen could recover only 43% of its shape. Hence, it can be stated that a inverse relationship
existed between strain and shape recovery ratio for Sample I. Another point to note that 50% strained
specimen could not reach a plateau value of recovery ratio up to 150 °C. This trend was more or less
similar for 100% strained specimen. However, a plateau like behavior was detected for 150% strained
specimen after 135 °C. As will be presented later, these trends can be explained using stress relaxation of
various specimens.
The effects of strain on recovery ratio of Sample II specimens are presented in Figure 4- 24. All
specimens followed the same trend. For example, high recovery ratio, existence of plateau, and the time
span at which shape recovery took place. The 50% strained specimen attained the highest shape recovery
ratio of 85.7%. The recovery ratios of the 100% and 150% strained specimens were 84.6% and 81.4%,
respectively.
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74
Figure 4- 24 Strain effect on shape recovery behavior of Sample II. Stretching rate was 50 mm/min and
stretching temperature was 85 °C. Heating rate was 4 °C/min.
Figure 4- 25 Strain effect on shape recovery behavior of Sample III. Stretching rate was 50 mm/min and
stretching temperature was 110 °C. Heating rate was 4 °C /min.
Figure 4- 25 illustrates the shape recovery curves of Sample III under different strain conditions.
Specimens strained at 50%, 100%, and 150% produced the shape recovery values 98%, 85%, and 88%,
respectively. In addition, it is seen that recovery started at around 50- 60 °C and ended at 110 °C. Note that
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was the stretching temperature of Sample III was also 110 °C. Furthermore, Sample III produced the
sharpest recovery of shape among these samples.
4.4.2.2 Influence of stretching rate on recovery stress
The effect of stretching rate on shape memory properties is one of the least studied deformation
conditions in SMPs literature [53]. In this study, Sample II was chosen as a representative to investigate the
role of stretching rate on recovery stress. The conditions of stretching were as follows. The crosshead speed
was varied to produce stretching rates of 10 mm/min, 50 mm/min, and 100 mm/min. The allowed strain
was 100% and heating rate was set to 85 °C/min. Figure 4- 26 presents the results of such tests. It is
observed that recovery stress showed a sharp increase within the first two minutes. Also, the recovery stress
increased with stretching rate. Specimen with 10 mm/min stretching rate produced a maximum value of 4.5
MPa, the specimen strained with 50 mm/min produced 8 MPa. Finally a recovery stress of 9.8 MPa was
obtained for 100 mm/min.
Figure 4- 26 The effect of stretching rate on recovery stress of Sample II.
The trend of the curves in Figure 4- 26 can be attributed to the extent of orientation of chains
produced by deformation at different stretching rates. If the rate of stretching is low, the oriented chains
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may undergo relaxation and cannot store much elastic energy. On the other hand, at high rates of stretching,
polymer chains primarily orient conforming to the rate of stretching and therefore do not undergo much
relaxation.
4.4.2.3 The influence of stretching temperature on recovery stress
The deformation temperatures of SMPs exert strong influence on recovery stress [60]. It was
found that SMP specimens stretched at higher temperature also produce higher recovery stress. Molecular
mobility is higher at higher stretching temperatures. This helps SMPs store most of the elastic energy due
to chain orientation. SMPs stretched at lower temperatures, on the other hand loses a part of its energy to
heat dissipation and chain breakage, due to higher viscosity and slower chain mobility [53].
Figure 4- 27 shows the effects of stretching temperature on Sample II. Three different
temperatures were chosen, 65 °C, glass transition temperature of Sample II, 85 °C (20 °C above the
transition temperature) and 100 °C (35 °C above the onset of glass transition temperature of Sample II).
The specimen deformed at 65 °C exhibited a very large relaxation after 4 minutes. This was however not
the case with the other specimens. Specimens deformed at 100 °C and 85 °C showed plateau values of
recovery stress. The different behavior for the specimen deformed at 65 °C can be explained as follows. It
is a norm to choose a deformation temperature approximately 20 °C higher than the glass transition
temperature so as to obtain the best shape memory properties. In this context, the specimen deformed at 65
°C (Tg) was outside this norm. A number of factors contributed to the behaviors seen in Figure 4- 27 for
specimen stretched at 65 °C. First, although the Tg of this material was 65 °C (Figure 4- 6) the onset of
glass transition started much earlier, e.g., at around 30 °C. Accordingly, a fraction of the chains was glassy
at 65 °C, while another fraction was rubber. It was the rubbery chains which deformed easily at 65 °C. The
glassy chains possibly did not undergo deformation. Therefore, only a small amount of elastic energy was
stored by the rubbery chains, which led to small value, e.g., around 7 MPa, recovery stress in Figure 4- 27.
Also note that the specimen may have undergone some degree of plastic deformation. The fraction of
chains which were glassy at 65 °C and which did not deform at 65 °C, now became rubbery at temperatures
above 65 °C and began flowing. This led to strong decline of recovery stress at temperatures around 80 °C.
77
Figure 4- 27 The effect of stretching temperature on recovery stress of Sample II.
4.4.2.4 The influence of heating rate on recovery stress
Heating rate here refers to the rate at which the DMA set up was programmed to raise its
temperature. Some studies already reported on the roles of heating rate on shape memory properties
[26,53,60].
In this study, 50% strained Sample II specimens produced with a stretching rate of 50 mm/min were
used. The specimens were stretched at 85 °C which was 20 °C above Tg. The test consisted of a
temperature sweep from 20 °C to 150 °C with various heating rates: 4 °C/min, 8 °C/min, and 16 °C/min.
The test results are presented in Figure 4- 28. The maximum values of recovery stress maximum were
determined as 10.2 MPa, 8.7 MPa, and 6.9 MPa, respectively for heating rates of 4 °C/min, 8 °C/min, and
16 °C/min. Thus, the recovery stress decreased with increase of heating rate. This finding was different
from what was found by Cao et al. [53]. Results of that study revealed that recovery stress increased with
higher heating rates. This was attributed to enhanced the release of stored stress upon higher heating rate.
However, in this research heating rate investigation produced exactly opposite results which were also in
agreement with those published by Gall et al. [26]. In that study, a shift of maximum recovery stress to
higher temperatures and stress values occurred, as heating rate increased. In view of this, the findings of
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78
heating rate investigation were in good confirmation with the results of this study. Because both peak
maximum values lagged to higher temperatures and recovery stress occurred at lower maximum values.
The different results of the studies were attributed to the nature of materials. Both studies utilized of quite
different SMPs. Cao et al. used SMPUs and Gall et al. used a thermoset epoxy system cured thermally and
tested it by bending method rather than tension, while this study utilized partially reacted SMPU and PB-a.
Thus, morphologies even though based on amorphous structures differed very much.
Figure 4- 28 The influence of heating rate on recovery stress of 50% strained Sample II specimens.
4.4.3 Shape fixity
Shape fixity refers to the unconstrained recovery of SMPs upon unloading of stretching load and
upon cool up of stretched specimen to room temperature. This can be explained as an instantaneous
shrinkage of a very small portion of polymer chains that cannot preserve the imposed strain. This
immediate strain relaxation is usually around 0.5 to 10 percent for many shape memory polymers [13, 16,
21].
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79
Figure 4- 29 Shape fixity of the samples.
Figure 4- 29 presents values of the shape fixity test. The shape fixity tests were conducted at
50mm/min stretch rate and 100% strain. For each sample, five specimens were used to obtain mean values
of shape fixity. The standard deviation was determined to be less than 1.5%. Results of the tests
demonstrated that Sample II and III showed superior shape fixity over Sample I. The highest average value
was for Sample III with 99.3% which was followed by Sample II with 96.6% and Sample I with 93.5%. It
should be noted that the hard segment percentage should be high for higher shape fixity. Preserving
orientation of hard segments is much easier in systems with higher hard segment content. It should be noted
that hard segment domains tend to orient much faster and easier, and remain in oriented states longer.
4.4.4 Stress relaxation behavior
Relaxation of applied stress should be avoided to obtain higher recovery stress. The stress in sample
deformed suddenly to a desired strain can be easily measured to obtain information on stress relaxation
[79]. In view of this, stress is presented as a decaying function of time and mostly presented as a ratio or
percentage. The relaxation behavior of shape memory polyurethanes (SMPUs) have been studied by
several researches [14, 53, 81]. It is considered that the relaxation is governed by three different response
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80
times under fixed strains. The first relaxation time is the quickest and is related to hard domain breakup, the
second response time is due to the soft segment chain slip offs that occur after the breakage of hard
domains. The third relaxation time is due to separation of the soft segment chains from hard domains. This
occurs only after very long durations under strain.
The stress relaxation behavior of SMPs is of great significance to recovery stress, since recovery
stress is preserved through a state of polymer at which the polymer chains are able to sustain the
deformation energy as elastically stored energy. In this state, polymer chains are restrained and ready to
return to their original conformation, if strain is removed. Thus, polymer chains should be thoroughly
cooled to preserve the strained state, otherwise they exhibit stress relaxation at higher stretching
temperatures. At this point, the competitive effects of stress relaxation and higher deformation temperature
on recovery stress should be noted. As mentioned earlier, an increase in deformation temperature increases
the value of recovery stress. On the other hand, at higher deformation temperatures stress relaxation may
occur readily and consequently less recovery stress may result.
Figure 4- 30 The influence of strain on relaxation behavior of Sample II.
Earlier work on polybenzoxazine films did not report stress relaxation. In addition, no research work
considered a study on stress relaxation on polybenzoxazine and polyurethane blend system. In this work,
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the stress relaxation tests were carried out to investigate the effect of stretching rate and the strain on stress
relaxation behaviors of Sample II and III. The effects of strain and stretching rate on stress relaxation
behavior of Sample II are presented respectively in Figure 4- 30 and Figure 4- 31.
The test conditions were as follows. The stretching rate was 50 mm/min and the stretching
temperature was 85 °C. In addition, the relaxation test duration was 30 minutes. The data in Figure 4- 30
revealed that the percentage of stress relaxed varied between 42% and 47% at strains between 50-150%.
Furthermore, the stress relaxed rapidly less than two minutes of test.
Figure 4- 31 The influence of stretching rate on relaxation behavior of Sample II.
Similar behavior was also observed in the case of stretching rate. The data on the effects of
stretching rate is presented in Figure 4- 31. The data indicates that only 40-45% of stress could be held by
the specimens of Sample II irrespective of the stretching rate of strain.
Figure 4- 32 exhibits the influence of strain on the relaxation behavior of Sample III. The specimens
were deformed with a stretching rate of 50 mm/min at 110 °C. The test duration was set to 30 minutes.
Results of the tests demonstrated that as the strain increased from 50% to 100%, stress relaxation ratio fell
from 73% to 43%. However, 150% strained specimen led to an intermediate stress relaxation ratio with
65%. Nearly 80% of relaxation was completed for each specimen after the first two minutes of test.
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Figure 4- 32 Strain effect on the relaxation behavior of Sample III.
Figure 4- 33 The influence of stretching rate on the relaxation attitude of Sample III.
The influence of stretching rate influence on stress relaxation behavior of Sample III is shown in
Figure 4- 33. It is seen that stress relaxation occurred faster in specimens stretched at higher rate. Though
the difference between recovery stress values for specimens stretched at 50mm/min and 100 mm/min was
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not very much, respectively, 73% and 79%, the specimen stretched at 10 mm/min led to much less stress
relaxation, around 44%. It can be surmised that the polymer chains find more time at slower stretching rates
to attain better molecular and segmental orientations, e.g., conformational arrangements and the alignments
of polymer chains. The same finding was also reported in another study on shape memory polymers [53].
4.5 Morphological properties
Three important tools find extensive use in determining morphology of polymer systems:
transmission electron microscopy (TEM), scanning electron microscopy (SEM), and atomic force
microscopy (AFM). Even though SEM and TEM remained to be an option for determining domain size and
distribution, AFM was preferred in this study. AFM was useful as it aimed at understanding the
dissimilarity in mechanical properties of domains that existed on the surfaces of PU/PB-a samples. In
addition to the topographical images that showed the height of local points on the surfaces, phase images
produced the distribution of soft and hard domains as dark and bright colored spots.
The basic principles of tapping mode in AFM can be briefly explained. A tip attached to a cantilever
makes oscillatory contact of with sample surface. During this movement, the interactions between the
sample and the tip change the amplitude to the tip for its oscillation to remain at the same frequency.
Besides the amplitude, resonance frequency and the phase angle of oscillation alter as well. The oscillation
frequency of the cantilever is kept very close to the resonance frequency of the cantilever for the best
contact conditions. A typical phase image is the result of the interactions mentioned above and is related to
the energetic interactions [190-192].
The following high quality images were made possible by film specimens having very smooth
surfaces. The procedure was explained in the experimental section thoroughly. All images were taken in
tapping mode with a scan rate of 1 Hz. Scan size was 5 μm. Height profile was 200 nm or smaller and
phase angle was kept between 90°and 150
°. Force amplitude value was set between 3-6 mV. As the tip on
cantilever, a silicon tip with a bending spring constant of 14 Nm-1
and resonance frequency of 138 kHz was
utilized.
84
Figure 4- 34 AFM phase (A) and height (B) images of sample I, scan size; 5um.
Figure 4- 34 respectively shows phase (A) and height (B) images of Sample I. A phase contrast
appeared in the first image with dark and bright spots distributed throughout the sample. On the left corner
of this image, a very large bright spot was identified. In addition, relatively smaller spots were found
distributed on the surface. The dark spots in phase image (first figure) corresponded to the bright spots in
topography image (second figure). The height varied from 0 to 500 nm, while phase angle shift altered
from 0 to150°. Average domain size as measured by the Nanoscope imaging software was around 200 nm.
The color difference in the phase image was attributed to the mechanical property differences of the
domains. The abundance of bright spots in the phase image was considered to be due to the hard segment
domains, since the hard segment percentage was 71.2% for Sample I. In addition, it is observed that the
distribution of bright and dark spots was mostly uniform indicating that micro-phase separation took place
throughout the surface.
AFM phase (A) and height (B) images of Sample II are illustrated in Figure 4- 35. In the height
image, the color contrast appeared very much in detail and dark and bright spots are seen very clearly. It is
seen that the phase image did not possess the same contrast. Only a little color difference was distinct.
A B
85
Figure 4- 35 AFM phase (A) and height (B) images of sample II, scan size; 5um.
Results of AFM on Sample II indicate that phase separation was substantially reduced compared to
Sample I. This is an agreement with phase separation data presented earlier from FT-IR, e.g., 19%. This
can be attributed to chemical reaction between polyurethanes and cured polybenzoxazine during thermal
treatment in the vacuum oven. This was thought to have diminished phase separation between PU hard and
soft domains. In views of the phase image of Sample II, it can be inferred that the reaction product of
PU/PB-a did not constitute a separate phase.
Figure 4- 36 AFM phase (A) and height (B) images of sample III, scan size; 5um.
A B
A B
86
The phase (A) and height (B) images of Sample III are presented in Figure 4- 36. Both images
revealed differences in comparison to Sample I and Sample II. In phase image (A), small particles within
larger aggregates appeared. The average size of the particles was less than 100 nm whereas the size of
aggregates nearly 1 μm. Another difference was the distribution of dark and bright spots. All dark spots
appeared around the aggregates while the rest remained in relatively bright spots.
The AFM images of Sample III can now be compared with other test results. The hard segment
percentage calculated for Sample III was 74.5% and DSC scans exhibited no crystalline melting behavior
for Sample III (Figure 4- 4). This led us to infer that the distribution of hard segment domains was affected
by the presence of polybenzoxazine and it precursors as well as by the urethane- polybenzoxazine reaction
products. The soft phase surrounded the aggregate in phase image (Figure 4- 36 (A)) was due to the
presence of this product. Note that hard segment phase separation reduced as much as 50% between
Sample I and Sample III. Also note that a portion of benzoxazine did not polymerize and only small
fraction of polybenzoxazine could not react with polyurethanes. Due to these facts, it was postulated that
the aggregates on the phase image were due to separation of polybenzoxazine that formed during thermal
treatment and possibly also reacted with polyurethanes. Other smaller bright particles that were well
distributed throughout the surface represent the non-reacted benzoxazine precursors or polybenzoxazine.
The average size of these smaller bright particles was measured to be around 70 nm and the average
aggregate size was approximately 1μm.
The morphological features of the three samples can be represented in a simple manner by the
following sketches in Figure 4- 37. Sample I is shown as chain extended hard domains connected with
polyol chains and hydrogen bonding between urethane groups. Sample II is illustrated as a similar system
with some differences due to the presence of polybenzoxazine and its reaction with the urethane groups.
First, the usual hard domains of polyurethane deteriorated, the crystallinity and distribution of domains
were affected. Second, phase mixing occurred because of both the PU/PB-a reaction and the lower level of
hydrogen bonding between usual the hard and soft segments of polyurethane. Sample III is presented by the
third sketch where usual polyurethane hard domains are now involved in even a higher extent of chemical
and physical interactions with polybenzoxazine and its precursors. Phase mixing is more pronounced and
no crystallinity is present. In addition, agglomerations of polybenzoxazine surrounded by the urethane
87
groups and soft segment chain are observed in response to higher levels of benzoxazine in the formulation.
Further, smaller size benzoxazine precursors are present throughout the sample.
Figure 4- 37 Simple sketches for the morphological features of the samples.
88
CHAPTER V
CONCLUSIONS
In this thesis, a thorough study of polyurethane/polybenzoxazine shape memory polymer system
was presented. Important shape memory properties such as recovery stress, shape recovery and shape fixity
were characterized. Besides, the influences of deformation conditions particularly on recovery stress were
evaluated, which included strain, stretching rate, stretching temperature, and heating rate.
Benzoxazine in polyurethane formed a partially reactive system where predetermined portions of
polyurethane were allowed to react with phenolic OH groups in polybenzoxazine molecules. In view of
this, most thermal and mechanical properties of polyurethane were affected by the changes which involved
polybenzoxazine physical interactions with the soft and hard segment domains and polybenzoxazine-
polyurethane reactions. The following conclusions can be drawn from the data presented in this thesis.
PU/PB-a samples produced better shape memory properties than polyurethanes. Sample III
exhibited the highest shape recovery performance with 13 MPa recovery stress and 93%
shape recovery ratio. The presence of polybenzoxazine led to higher level of elastically
stored energy in the polymers due to chemical and physical interactions with polyurethane
hard segment.
The deformation conditions affected shape memory performance considerably. Higher
stretching rates and stretching temperatures improved the values of recovery stress by as
much as 45%. On the other hand, higher heating rates caused nearly 30% reduction in the
values of recovery stress, e.g., 10.1 MPa for 4 °C/min to 8.7 MPa for 8 °C/min, and to 6.9
MPa for 16 °C/min.
The maximum values of recovery stress of Sample I and II showed improvements with
strain of the deformed specimens. Of three samples, Sample I exhibited the highest
89
increase in recovery stress with strain, e.g., 4.5 MPa for 50%, 5 MPa for 100 MPa, and 14
MPa for 150%.
The maximum values of shape recovery of the samples declined with an increase of strain.
The most dramatic change occurred with Sample I. The shape recovery values reduced
from 90% at 50% strain to only 43% at 150% strain. In contrast, Sample II and III retained
the shape recovery values with increase of strain. While Sample II retained its recovery
around 81% at 150% strain, Sample III demonstrated the highest recovery ratio with 90%
at 150% strain.
All samples retained a large fraction of their temporary shapes and offered high values of
shape fixity. However, a clear trend was seen with the increase of benzoxazine content in
the samples, e.g., 93 % for Sample I, 97 % for Sample II, and 99 % for Sample III.
Stress relaxation ratio of the specimens with benzoxazine was investigated in conjunction
with the influences of stretching rate and strain. The results were used to correlate the
trends of shape recovery properties. It was found that the relaxation behavior of Sample II
was not affected much by strain and stretching rate, while Sample III exhibited an increase
of stress relaxation at higher stretching rates
The chemical interactions between polybenzoxazine and polyurethanes were found to alter
the morphology of the samples. AFM phase image of Sample I demonstrated a phase
separated morphology with dark and bright spots indicating respectively the soft and hard
domains. AFM image of Sample II showed that there was not much phase difference. The
phase image of Sample III indicated the presence of cured polybenzoxazine clusters with
sizes ranging from 70 nm single particles to 1 μm aggregates. The dark spots in the
neighborhood of these aggregates were due to soft and hard segment domains.
90
CHAPTER VI
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