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 POLYURETHANE-POLYBENZOXAZINE BASED
SHAPE MEMORY POLYMERS
Numan Erden
Thesis
Approved: Accepted:
______Advisor Department Chair Dr. Sadhan C. Jana Dr. Sadhan C. Jana
______Faculty Reader Dean of the College Dr. Kevin Cavicchi Dr. Stephen Z.D. Cheng
______Faculty Reader Dean of the Graduate School Dr. Robert A. Weiss Dr. George R. Newkome
______Date
ii 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. iii 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.
iv DEDICATION
Dedicated to my parents; Casim and Şükriye Erden for their everlasting affection and love...
v 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 upon 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.
vi 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
vii 2.4.1 Synthesis of the benzoxazine ...... 30
2.4.2 Polymerization ...... 322
2.4.3 Distinguishing properties of polybenzoxazines ...... 333
2.5 Earlier polybenzoxazine/polyurethane work ...... 344
III. EXPERIMENTAL ...... 388
3.1 Materials ...... 388
3.1.1 Raw materials for synthesis of polyurethane ...... 388
3.1.2 Raw materials for benzoxazine ...... 399
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 ...... 67
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 viii
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...... 55
4- 2 Viscoelastic properties of the samples ...... 58
4- 3 TGA analysis results. Scan rate was 20 °C/min...... 59
-1 -1 4- 4 Hydrogen bonding % of urethane domains. AFCO refers to 1730 cm and ANCO; 1700 cm ...... 64
4- 5 Tensile properties of the samples...... 65
ix 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...... 333
3- 1 Chemical formula of MDI...... 388
3- 2 Chemical formula of PTMG...... 399
3- 3 Chemical formula of 1,4-butanediol...... 399
3- 4 General chemical formula for DABCO T120...... 399
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
3- 9 The time-temperature plot showing duration and temperature of mixing, compression molding and vacuum oven...... 44
x 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...... 61
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...... 68
4- 19 Shape recovery ratio of 100% strained samples. Heating rate was 4 °C/min and stretching rate was 50 mm/min...... 69
4- 20 Strain effect on recovery stress trend of Sample I. Stretching rate was 50 mm/min and stretching temperature was 71 °C...... 71
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
xi 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 images of sample I, (A) phase and (B) height, scan size: 5μm...... 84
4- 35 AFM images of sample II, (A) phase and (B) height, scan size: 5μm ...... 85
4- 36 AFM images of sample III, (A) phase and (B) height, scan size: 5μm...... 85
4- 37 Simple sketches for the morphological features of the samples...... 87
xii 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 1 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
2 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 3 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 4 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 5 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.
6 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 7 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 theses, 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 8 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
9 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 10 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
11 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.
12 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.
13
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 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]. 14 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. 15 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 16 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]. 17 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. 18 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]. 19 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. 20 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. 21 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]. 22 Table 2- 1 Frequently used reactants for polyurethane synthesis. Methylene bis (p-cyclohexyl isocyanate) (H12MDI) Castor oil Methylene bis (p-phenyl isocyanate) (MDI) Polycaprolactone diol Isophorone diisocyanates (IPDI) Poly (tetramethylene) glycol (PTMG) Hexamethylene diisocyanate (HDI) Poly (ethylene adipate) Poly (dimethyl siloxane) PDMS (hydroxy butylated) Toluene diisocyanate (TDI) 23 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 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]. 24 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. 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 25 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. 26 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 C, which was the crystalline melting range of the soft phase derived from PCL with a molecular weight ⁰of 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 molecule attachments [148], enhanced electrical conductivity and optical arrangements. Of particular 27 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 C and 40 C. A fluctuating resistance vs. temperature curve was obtained for CNT around 70 C. Furthermore,⁰ 28 ⁰ ⁰ 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 beneficial, if optimum processing conditions and viable applications are found. In those studies, the 29 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. 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 30 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]. 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. Figure 2- 10 The synthesis of mono functional 3,4-dehydro,2H,1,3-benzoxazine adapted from Ref. 38. 31 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]. 32 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]. 33 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. 34 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 35 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 36 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. 37 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. Figure 3- 1 Chemical formula of MDI. 38 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. 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. 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. 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. 39 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. 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. 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 40 utilizing a three neck reaction bottle for the synthesis of prepolymer and Brabender mixer for chain 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 reaction bottle was heated to 75 °C and the purge gas was turned on. PTMG was poured fist followed by 41 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. 42 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 43 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. 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. 44 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 45 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 46 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. 47 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. 48 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. 49 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 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 50 observed. The first transition occurred between 25 °C and 50 °C, while the second transition occurred between 160 °C and 175 °C. 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, where as 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%. 51 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 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 52 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 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 53 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 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 st that as the benzoxazine amount in the samples was increased, the onset value of 1 Tg shifted to higher nd temperatures, while that of 2 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 nd other hand, the fall in the 2 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. 54 Table 4- 1 Glass transition onset values determined through DSC with higher heating rates. st nd Samples 1 Tg (°C) 2 Tg (°C) I 44 106 II 51 97 III 62 96 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 55 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, 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. 56 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. 57 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 4.1x109 5.3x108 11.3 51 II 5.3x109 3.8x109 3.4x107 147 65 III 4.7x109 3.9x109 1.4x107 271 91 110 1.E+9 100 90 1.E+8 C) ° 80 Temperature ( Temperature 70 1.E+7 Storage Modulus (E') Modulus Storage Sample II-Tg 60 Sample III-Tg Sample II- E Sample III- E 50 1.E+6 0 12 24 36 48 60 Curing Time (hrs) 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 x108 Pa and Sample II exhibited a decrease from 1.4 x107 Pa to 1.2 x107 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. 58 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. 1st Transition 1232 cm-1 2nd Transition 3rd Transition 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. 1st 2nd 3rd 5wt. % loss Sample Residue (wt. %) Transition Transition Transition (°C) (°C) (°C) (°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 59 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. 60 1232 cm-1 942 cm-1 Figure 4- 11 FT-IR spectra of benzoxazine monomer. Figure 4- 12 1H-NMR analysis of benzoxazine monomer (bis (3-phenyl-3,4-dehydro-2H-1,3- benzoxazinyl)-isopropane). 61 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. 2281 cm-1 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 62 (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. Sample I 1703 cm-1 1733 m-1 Sample II Sample III 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 63 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. -1 -1 Table 4- 4 Hydrogen bonding % of urethane domains. AFCO refers to 1730 cm and ANCO; 1700 cm . 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 64 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 Elongation % Yield Stress Ultimate Stress Toughness (MPa) at Break (MPa) (MPa) (MPa) I 228 478 41 50 216 II 258 95 44 35 36 III 475 32 N/A 160 6 180 Sample I 160 Sample II 140 Sample III 120 100 80 Stress MPa Stress 60 40 20 0 0 1 2 3 4 5 6 Strain % (X100) 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 65 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 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. 2 1.5 1 Stress (MPa) Stress 0.5 Sample I Sample II Sample III 0 0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 Strain (x100) 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 66 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. 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. 67 14 12 10 8 6 4 Recovery (MPa) Stress Recovery Sample I 2 sample II Sample III 0 0 50 100 150 200 Temperature (°C) Figure 4- 18 Recovery stress behaviors of 100% strained samples. Heating rate was 4 °C/min and stretching rate was 50 mm/min. 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. 68 100 80 60 40 sample I Shape rcovery Shapercovery % 20 sample II sample III 0 0 50 100 150 200 Temperature (°C) 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 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 69 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. 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]. 70 16 80 14 70 12 60 C) ° 10 50 8 40 6 30 4 50% 20 ( Temperature Recovery (MPa) Stress Recovery 2 100% 10 150% 0 0 0 5 10 15 20 Time (min) Figure 4- 20 Strain effect on recovery stress trend of Sample I. Stretching rate was 50 mm/min and stretching temperature was 71 °C. 12 100 90 10 80 70 C) 8 ° 60 6 50 40 4 30 50% ( Temperature Recovery Recovery stress (MPa) 20 2 100% 10 150% 0 0 0 5 10 15 20 Time (min) 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 71 degree of relaxation after reaching the peaks values. Another interesting point to note is how rapidly the maximum value of recovery stress was reached. 10 140 120 8 100 C) ° 6 80 4 60 50% 40 Temperature ( Temperature 2 100% Recovery stress ( MPa ) stress Recovery 20 150% 0 0 0 5 10 15 20 Time (min) 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. 72 100 90 80 70 60 50 40 30 Shape recovery recovery Shape % 50% 20 100% 10 150% 0 0 50 100 150 200 Temperature (0C) 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. 73 100 90 80 70 60 50 40 30 100% Shape recovery % recovery Shape 20 150% 10 50% 0 0 50 100 150 200 Temperature (0C) 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. 100 90 80 70 60 50 40 30 50% Shape recovery % recovery Shape 20 100% 10 150% 0 0 50 100 150 200 Temperature (°C) 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 74 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. 12 100 90 10 80 ) 70 C 8 ° 60 6 50 40 4 10 mm/min 30 Temperature ( Recovery stress (MPa) 20 2 50 mm/min 100 mm/min 10 0 0 0 2 4 6 8 10 12 Time (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 75 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. 76 12 100 90 10 80 70 8 C) 60 6 50 40 4 30 ( Temperature Recovery Recovery stress (MPa) 65 °C 20 2 85 °C 10 100 °C 0 0 0 2 4 6 8 10 12 Time (min) 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 77 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. 12 10 8 6 4 4 °C/ min 8 °C/min Recovery (Mpa) Stress Recovery 2 16 °C/min 0 0 50 100 150 200 Temperature (°C) 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]. 78 100 99 98 97 96 Sample I 95 Sample II Shape fixity Shape fixity % Sample III 94 93 92 0 1 2 3 4 5 6 Specimen # 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 79 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. 50 45 40 35 30 25 20 15 50% Relaxation ratio % % ratio Relaxation 10 100% 5 150% 0 0 5 10 15 20 25 30 35 Time (min) 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, 80 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. 50 45 40 35 30 25 20 15 10 mm/min 10 50 mm/min Relaxation Ratio Ratio Relaxation % 5 100 mm/min 0 0 5 10 15 20 25 30 35 Time (min) 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. 81 80 70 60 50 40 30 Relaxation Ratio Ratio Relaxation % 20 50% strain 10 150% strain 100% strain 0 0 5 10 15 20 25 30 35 Time (min) Figure 4- 32 Strain effect on the relaxation behavior of Sample III. 90 80 70 60 50 40 30 10 mm/min Relaxation Ratio Ratio Relaxation % 20 50 mm/min 10 100 mm/min 0 0 5 10 15 20 25 30 35 Time (min) 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 82 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. 83 A B Figure 4- 34 AFM images of sample I, (A) phase and (B) height, scan size: 5μm. 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. 84 A B Figure 4- 35 AFM images of sample II, (A) phase and (B) height, scan size: 5μm. 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. A B Figure 4- 36 AFM images of sample III, (A) phase and (B) height, scan size: 5μm. 85 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 86 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. 87 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 88 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. 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